Nickel-Rich Layered Cathode Materials for ... - ACS Publications

Nickel-Rich Layered Cathode Materials for Automotive Lithium-Ion Batteries: Achievements and Perspectives Seung-Taek Myung,†,# Filippo Maglia,‡,# Kang-Joon Park,§ Chong Seung Yoon,∥ Peter Lamp,‡ Sung-Jin Kim,*,‡ and Yang-Kook Sun*,§ †

Department of Nano Engineering, Sejong University, Seoul 143-747, South Korea BMW Group, Petuelring 130, 80788 München, Germany § Department of Energy Engineering, Hanyang University, Seoul 133-791, South Korea ∥ Department of Material Science and Engineering, Hanyang University, Seoul 133-791, South Korea ‡

ABSTRACT: Future generations of electric vehicles require driving ranges of at least 300 miles to successfully penetrate the mass consumer market. A significant improvement in the energy density of lithium batteries is mandatory while also maintaining similar or improved rate capability, lifetime, cost, and safety. The vast majority of electric vehicles that will appear on the market in the next 10 years will employ nickel-rich cathode materials, LiNi1−x−yCoxAlyO2 and LiNi1−x−yCoxMnyO2 (x + y < 0.2), in particular. Here, the potential and limitations of these cathode materials are critically compared with reference to realistic target values from the automotive industry. Moreover, we show how future automotive targets can be achieved through fine control of the structural and microstructural properties. choice of the automotive industry for “near-future” EV generation will be, in most cases, a direct evolution of the cathodes already employed by the large majority of car manufacturers,8 that is, a transition metal layered oxide belonging to the LiNi1−x−yCoxAlyO2 (NCA) or LiNi1−x−yCoxMnyO2 (NCM) family. In principle, layered−layered materials (general formula xLi2MnO3(1−x)LiMO2 (M = Ni, Co, Mn) would allow one to reach even higher capacities (above 250 mAh g−1) when used at upper cutoff voltages of 4.6 V.5 Nevertheless, the voltage fade observed upon cycling and low power capabilities seem, despite the huge man and monetary effort, in particular, from the U.S. DoE, not solvable at the moment. A reduction of the abovementioned effort seems to be planned for the near future.

D

uring the past decade, the Li-ion batteries (LIBs) market has expanded from consumer electronics to other areas including the automotive industry. Currently, LIBs are applied at all levels of electrification, from hybrid (HV) to plug-in hybrid (PHEV) to fully electric (EV) vehicles. Nonetheless, the share of EVs in the world’s automotive market is increasing at a slower rate than expected and hoped.1 Key factors, cost-to-range ratio and charging time in particular, represent obstacles to the mass-market penetration of EVs. An electromobility model based on daily use in an urban environment might not be sufficient to achieve wide customer acceptance. Driving ranges of at least 300 miles are likely required to guarantee the success of future EVs.2 The quest for an extended driving range has fueled the appearance of countless publications focusing on novel electrode active materials, in particular, for the cathode compartment.3−7 These new compounds are expected to offer the potential for higher energy densities due to either higher operating voltages or larger capacities, the latter in some cases achievable via multielectron Li intercalation reactions.6 Despite the large number of novel, sometimes exotic, compounds that appear in the scientific literature nearly every week, the time restrictions imposed by the development of battery concepts for automotive applications drastically limit the choices. Cathode active materials that can realistically find application in the automotive industry in the next 10 years are most likely restricted to a small number of compounds already known and investigated.8 The cathode of © 2016 American Chemical Society

The potential and limitations of these cathode materials are critically compared with reference to realistic target values from the automotive industry. This does not mean that all targets have been achieved. Driving range requirements impose an increase in the energy density of layered oxides that requires extending their intrinsic capacity Received: November 11, 2016 Accepted: December 16, 2016 Published: December 16, 2016 196

DOI: 10.1021/acsenergylett.6b00594 ACS Energy Lett. 2017, 2, 196−223

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http://pubs.acs.org/journal/aelccp

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but also a large number of assumptions concerning, for example, cell and battery design, and inert electrode components. Moving from right to left in Figure 1, the pure cathode active material energy density is first translated into the corresponding value at the full cathode level. Parameters such as porosity, electrode thickness, and electrode composition are required and depend on the specific properties of the active material such as particle size and conductivity. Moving further to the left (cell level), the cell format must be selected including the cell type and size, number of jelly rolls, and type of anode. Three main cell formats are used in commercial EVs or HVs: prismatic hard case, cylindrical, and pouch cells. These cells show different inner volume occupancies, typically higher in the cylindrical cells than those for the other formats. For what concerns packing efficiencies at the battery pack level, the prismatic cell is the most convenient format and will be consequently used in the calculations herein. Finally, assumptions must be made concerning the strategies to connect individual cells to battery modules and battery packs in order to achieve the practical voltages (typically between 200 and 400 V) and capacities required for vehicle applications. Different choices at all of the above-described levels will lead to somewhat different target values at the cell, electrode, and material levels, which are accounted for in Figure 1 by the error bars. Figure 1 indicates an average specific cell energy target of around 300 Wh kg−1, an increase of a factor of ∼2.5 with respect to the actual value. At the cathode active material level, a target value of at least ∼750 Wh kg−1 (or around 202 mAh g−1 at 3.7 V) is required. Figure 2 shows a possible evolutionary roadmap of NCM- and NCA-based cells in terms of gravimetric (Wh kg−1) and volumetric (Wh L−1) energy density. The possibility to achieve the automotive targets for 2025 (here indicated by the dashed green lines) is presented in terms of four key parameters: nickel content in the active material (from 0.33 to 0.90 mol %), anode type (graphite vs a Si−C anode with a 1000 mAh g−1 capacity), upper cell cutoff voltage (4.2 V vs 4.4 V), and electrode load (35% porosity, 15 mg cm−2 vs 20% porosity, 25 mg cm−2). The voltage and load limits of Figure 2 are representative of a conservative scenario (values similar to those currently employed) and a more competitive one and are achievable without critical changes in cell chemistry or design, in particular, for the type of electrolyte employed.8 The cell energy density values of Figure 2 were obtained using a custom software program considering a hard case prismatic cell format comprising two jelly rolls, as in the case of the cells used in the BMW i3, but the trends would be similar for any cell format. In these calculations, the projected volume and weight of the inert components were estimated based on the trend in commercialized products observed in the past few years. For example, the thicknesses of current collectors and separators have been set to 20 μm (aluminum foil and separator) and 10 μm (copper foil), respectively. Moreover, the amount of electrolyte was estimated based on the selected porosity (20 and 35%) and the wetting characteristics of the active material. The balance between anode and cathode (N/P ratio) was the same as the current value used, with 20% over dimensioning of the anode in order to ensure the safety of future cells.8 The comparison with the target value for the year 2025 suggests that Ni-rich NCM and NCA compounds (Ni > 80 mol %) have the potential to reach, or at least approach, energy densities of 300 Wh kg−1 and 700 Wh L−1 at the cell level. Nevertheless, it is evident that cells relying on graphite anodes have to withstand loadings, densities, and voltages at which

(e.g., Ni content), practical upper cutoff voltage, and/or electrode loading with respect to their actual values. Moreover, a successful increase of energy density heavily depends on the possibility of maintaining and possibly improving all other key parameters including power, lifetime, safety, and cost. Limitations in power, lifetime, and safety will unavoidably accompany more demanding use of layered oxide cathodes and still represent challenges from both the fundamental science and engineering points of view. Material suppliers, battery manufacturers, and original equipment manufacturers (OEMs) are intensively engaged in the exploration of possible compositional and morphological modifications of NCM and NCA compounds to improve lifetime and safety properties. Several strategies have been proposed to limit the cathode reactivity toward the electrolyte and to reduce the extent of transition toward unwanted crystallographic structures during Li deintercalation. Control of the particle and crystallite microstructures has also been investigated as a means to improve the transport properties, Li diffusion in particular. In this paper, we review the most recent and promising results concerning NCA and NCM cathode materials. Published results in this field, in particular, for those compositions highest in nickel, will be compared with the projected targets of the automotive industry in terms of energy, power, lifetime, and safety.

1. OEM TARGETS AND IMPACTS OF CATHODE MATERIALS When discussing properties and requirements at the battery pack, cell, or electrode level, results do not depend on the active material alone, as is the case for most of the data published in the scientific literature, but must also include the effects of other components, including the electrolyte, separator, electrode, cell, and battery design as well as the amounts of all inactive components. At the battery pack level, a driving range of 300 miles would correspond to energies around 60−70 kWh, assuming a space utilization of the cells similar to that of most of the EV models currently on the market. To discuss the potential offered by NCA and NCM cathodes to meet future automotive requirements, energy density values should be calculated at all levels, as shown in Figure 1. This requires not only knowledge of the fundamental material properties (capacity, Li+ insertion/extraction voltage, crystallographic density, etc.)

Figure 1. Specific energy targets from pack to material level together with design steps for battery packs. 197


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battery packs will be maximized. What additionally will set higher demands for future cells will be the increased requirements for charging. Here, typical automotive targets for the next decade establish charging times well below 30 min to regain 80% of the driving range, which implies challenging continuous charging currents. Lifetime requirements are also expected to negligibly vary in terms of energy throughput. Here, the 80% energy retention end of life criteria has to be fulfilled. This means that, given the larger energy of future battery packs, lifetime requirements in terms of the number of cycles are expected to correspondingly decrease. Finally, there will be no compromises in terms of cell or pack safety, and EVs are expected to have at least the same safety level as conventional cars. To qualify for automotive use, cells have to pass a set of abuse tests, and an at least hazard level 4 classification is required by the automotive industry (Table 1). Current cell Table 1. Eucar Hazard Levels12 hazard level

description

classification criteria and effect

0 1

no effect passive protection, activated

2

defect/damage

3

leakage, Δmass < 50%

4

venting, Δmass ≥ 50% fire or flame rupture explosion

no effect. no loss of functionality no defect; no leakage; no venting, fire, or flame; no rupture; no explosion; no exothermic reaction or thermal runaway; cell reversibly damaged; repair of protection device needed no leakage; no venting, fire, or flame; no rupture; no explosion; no exothermic reaction or thermal runaway; cell irreversibly damaged; repair needed no venting, fire, or flame; no rupture; no explosion; weight loss < 50% of electrolyte weight (electrolyte = solvent + salt) no fire or flame; no rupture; no explosion; weight loss ≥ 50% of electrolyte weight (electrolyte = solvent + salt) no rupture; no explosion (i.e., no flying parts) no explosion, but flying parts of the active mass explosion (i.e., disintegration of the cell)

5 6 7

Figure 2. (a) Cell Wh kg−1 and (b) cell Wh L−1 estimations calculated in prismatic hard case format for different cell chemistries. Values resulting from conservative and competitive electrode designs are indicated by min/max values. Passive materials are not optimized.

designs mitigate these hazardous events through a mix of chemical and construction measures, including blending of active materials, special additives in electrolytes, reinforced separators, and mechanical safety elements in the cell. Unfortunately, a clear link between active material properties, such as thermal stability and temperature/voltage-dependent oxygen release, is still not available. The possibility to establish a clear and possibly quantitative relationship between the results obtained at the material and electrode levels and the test listed in Table 1 represents one of the biggest challenges for the next decade.

lifetime and safety requirements have not yet been demonstrated. On the other hand, less severe conditions would be required when graphite is substituted, at least partially, by silicon. The advantage of Si-based anodes stems from their higher specific capacity.9−11 At constant N/P ratio, high-capacity anodes can be coated as thinner layers than when using anodes of graphite. The reduced space requirement on the anode side leaves space for a larger number of electrode windings, which results in an overall increase of the cell capacity. A quantitative discussion of power requirements is more complicated than that for energy density, and the link between the different levels is less straightforward. The status of power performance in commercialized vehicles is much less critical than that for driving distance, and power is generally considered a lesser issue for customer acceptance. As an example, the BMW full-electric i3 cells show constant full discharge capacities at different current densities ranging from 10 to 550 mA g−1 with almost no capacity decrease even at higher C rates. This indicates the excellent rate capability of actual cells, and increases of pack and cell energies could mitigate the requirements in terms of mA cell−1 or mA g−1 of active material. However, future power requirements are under constant optimization and depend on vehicle type and pack size. Higher requirements are likely, meaning that the P/E ratio even for large

2. DEVELOPMENTAL HISTORY OF NCA AND NCM CATHODES First introduced in 1991, LiCoO2 was the first layered transition metal oxide to be successfully incorporated into commercial rechargeable LIBs and is still used widely in portable electronic devices as a positive electrode. The layered LiCoO2 has a layer of Li ions and CoO6 octahedra stacked in an alternating ABCABC··· sequence, and this layered two-dimensional structure imparts high Li+ ion mobility to the material, generating a theoretical specific capacity of 272 mAh (g-oxide)−1 at a potential of 4.2 V vs Li0/Li+. However, the practical capacity is limited to >140 mAh g−1 because only ∼0.5 mol of Li+ ions can be reversibly cycled without causing capacity losses due to modifications in the LiCoO2 structure.3 Lately, LiCoO2 has been developed for an up to 4.45 V vs Li/Li+ upper cutoff voltage, showing approximately 180 mAh g−1 via surface modification.13−17 Despite the early success of the LiCoO2 cathode, the high material cost of Co has 198


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Figure 3. (a) Initial charge and discharge curves of LiNiO2 synthesized by a solid-state reaction of LiNO3 and Ni(OH)2 (reproduced with permission from ref 28, Copyright 1993, The Electrochemical Society); (b) cycling behavior of a Li/LiNi0.7Co0.3O2 cell at 0.4 mA cm−2 (reprinted from ref 32, Copyright 1993, with permission from Elsevier); (c) charge and discharge curves of LiNi0.75Al0.25O2 at a constant capacity of 150 mAh g−1 (reproduced with permission from ref 41, Copyright 1995, The Electrochemical Society); (d) comparison of the second charge− discharge cycle obtained at the C/20 rate for LixNi0.89Al0.16O2 and LixNi0.78Al0.26O2 cells (reprinted from ref 53, Copyright 2003, with permission from Elsevier); (e) specific discharge capacity of the cylindrical lithium-ion batteries of LiNi0.8Co0.15Al0.05O2 and LiNi0.75Co0.15Al0.05Mg0.05O2 with a graphite negative electrode, that is, Mg0 and Mg5, respectively, at a 2 C rate during cycling at 60 °C (reprinted from ref 54, Copyright 2005, with permission from Elsevier).

fueled the search for a new class of positive electrodes. In

As an alternative material to LiCoO2, the relatively low material cost of Ni led to investigation of LixNiO2, isostructural with LiCoO2 (when x > 0.6), as a possible candidate cathode for LIBs.20−22 Although LiNiO2 has a theoretical capacity of 275 mAh g−1, Arai et al. demonstrated that more than 200 mAh g−1 at

addition to the material cost, LiCoO2 can react violently with the liquid electrolyte in the fully charged state,18 leading to fire and explosion.19 199


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an effect similar to that of Co3+.41 It is also assumed that the Gibbs energy for the formation of Al2O3 at 25 °C (−1582.3 kJ mol−1) is sufficiently low to stabilize the crystal structure, meaning that the doped trivalent renders the crystal structure rigid. Therefore, simple S-shaped charge and discharge curves associated with a single-phase reaction are observed with excellent reversibility (Figure 3c). However, replacement of Al3+ in the transition metal layers should be carefully considered because it is an electrochemically inactive species. An example is given in Figure 3d that addition of more Al3+ decreases the capacity, which dilutes the benefit of the weight reduction achieved by Al doping. Meanwhile, the presence of Al3+ effectively increases the operation voltage due to weakening of the Ni−O bond by the stronger Al−O bond through the inductive effect; a similar tendency was observed in a LiCo1−xAlxO2 system.42,43 The balance between reduced capacity and increased operation voltage based on the amount of Al3+ should be considered in the design of Al-doped LiNiO2. Other metallic elements substituted in the Ni sites to form LiNiyM1−yO2 (M = Mg, Ti, Fe) were investigated, but the results were inferior to those previously discussed.44−51 The best overall performance of LiNiO2 was achieved when it was co-doped with either Co and Mn (NCM) or Co and Al (NCA). Various combinations of compositions of both NCM and NCA have been tested over the years, and both NCM and NCA cathodes have reached the stage where two electrodes can be safely commercialized, especially for large-format LIBs. In the following sections, NCM and NCA properties will be discussed in detail. 2.1.1. Development of NCA Materials. Because the abovementioned LiNi1−xCoxO2 and LiNi1−xAlxO2 compounds exhibited notable improvement in electrochemical properties compared to those of LiNiO2, hybridization of both compounds can lead to synergistic effects on the structure and electrode performance. In particular, cation mixing in the Li layer was reduced by the introduction of both trivalent Co and Al in the transition metal layer. Also, replacement by these elements successfully prevented multistep phase transition during Li+ insertion and extraction. Although Al3+ is electrochemically inactive, a redox reaction of Co3+/4+ is possible in the voltage range of 2.5−4.3 V, so that the added Co instead of Ni in the NCA compensates for the capacity, while Al3+ keeps the crystal structure rigid during repetitive Li+ insertion and extraction. However, a large amount of Al3+ is not favored due to its inactivity, which tends to lower the reversible capacity.52,53 The LiNi0.8Co0.15Al0.05O2 delivers a large discharge capacity of approximately 200 mAh g−1. Due to its stabilized crystal structure aided by Co3+ and Al3+, LiNi0.8Co0.15Al0.05O2 exhibits excellent cycling performance not only at room temperature but also at 60 °C. A recent report by Kondo et al.54 suggested that Mg doping into LiNi0.8Co0.15Al0.05O2 was further effective to reduce Ni2+ ions in Li layers, increase the average valence of Ni ions, and suppress phase transition during reaction with Li+ ions, showing 91% of its initial capacity even after 500 cycles at 60 °C (Figure 3e). LiNi0.8Co0.15Al0.05O2 is a currently widely used commercial material. The double doping of Co3+ and Al3+ is known to maintain better layer structure. Madhavi et al.55 have attempted to find an optimized composition using LiNi0.7Co0.3−zAlzO2 (0 < z < 0.15). Formation of solid solution was possible in the range, whereas formation of impurity was inevitable when z was greater than 0.15. In terms of capacity and retention, LiNi0.7Co0.25Al0.05O2 was the best composition. Although the chemical composition slightly deviated from Madhavi’s

4.5 V vs Li0/Li+ can be extracted from stoichiometric LiNiO2. However, a 20% capacity loss is observed in the first charge− discharge cycle23 due to the structural modification when charged below x = 0.5. Formation of stoichiometric LiNiO2 with a Li/Ni ratio of 1:1 is still difficult because it often results in lithium-deficient Li1−xNi1+xO2 with part of the Ni present as Ni2+ in the Li layers, primarily due to the similarity of ionic radii of Li+ (0.76 Å) and Ni2+ (0.69 Å).24−27 In addition, LixNiO2 is often Lideficient, and it is difficult to prepare stoichiometric LiNiO2. Ohzuku’s group28 noted the importance of the synthesis conditions of starting materials and O2 gas flow during synthesis, a slight excess of LiNO3 and Ni(OH)2 at 750 °C in O2. Namely, addition of excess lithium salt for the synthesis and consecutive washing of the product was a critical condition to deliver high capacity with good reversibility, delivering a discharge capacity of approximately 150 mAh g−1 in the voltage range of 2.5−4.2 V (Figure 3a). Further work by Arai et al.29,30 confirmed the importance of excess lithium, obtaining a large rechargeable capacity higher than 200 mAh g−1 in the voltage range of 3.0−4.5 V. To improve the cyclability of LiNiO2 and alleviate the synthesis problem of binary solid, solutions with the general formula LiNi1−xMXO2 have been investigated. Partial introduction of Co into the LiNiO2 crystal structure, LiNi1−xCoxO2, is readily derived due to the structural similarity between LiCoO2 and LiNiO2. Hence, a phase-pure LiNi1−xCoxO2 can be obtained in the whole range (x = 0−1).31−33 Structural stabilization is accomplished by the Co3+, which hinders formation of Ni2+ in the Li layer.32,33 In particular, LiNi0.5Co0.5O2 does not need an oxygen environment during heat treatment due to the air-stable character of LiCoO2. The stabilized structure provides a singlephase reaction instead of multistep phase transition, which is dominant in LiNiO2, during Li+ insertion and extraction, enabling improvement in the electrode performance (Figure 3b). LiNiO2−LiCoO2 solid solution was suggested, and LiNi1−yCoyO2 indeed greatly reduced the capacity loss during cycling.31,33−36 LiNi0.74Co0.26O2 was able to retain over 200 mAh g−1 for 70 cycles without significant capacity fading.35 It was shown that Co substitution at the Ni sites tends to reduce cation mixing and stabilizes the layered structure during cycling,31,34 demonstrating gradual decreases in both a- and c-axes but an increase in the c/a ratio. However, LiNiO2 in LiNi1−yCoyO2 in the electrochemically or chemically delithiated state was observed to transform to the rock-salt structure above 200 °C, accompanied by oxygen evolution,30 which initiated a violent exothermic reaction with the volatile organic solvent in the carbonate electrolyte.18,29 On the other hand, a small amount of Mn doping of LiNiO2 deteriorated the electrochemical properties (reversible capacity of only 110 mAh g−1 in the 3−4.15 V range)37 but improved the thermal stability of LiNiO2 in the delithiated state.38 The Mn doping suppressed the large exothermic reaction and imparted thermal stability to the LiNiO2 electrode. Doping of LiNiO2 with Al was also observed to minimize the detrimental phase transition and enhance the thermal behavior of the cathode.39,40 Al is an interesting dopant because its atomic weight is lighter than those of other transition metal elements. Hence, Al-doped LiNiO2 demonstrates an increased gravimetric capacity. In addition, Al is usually stabilized as Al3+, so that the presence of Al3+ in the transition metal layer tends to reduce the a-axis but increase the c-axis parameters because α-LiAlO2 (a = 2.80 Å and c = 14.23 Å) has smaller a-axis but larger c-axis parameters compared to those of LiNiO2 (a = 2.88 Å and c = 14.19 Å), which decreases the cationic disorder in the Li layers, 200


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Figure 4. (a) Interval storage and cycle test performance of NCR18650 (NCA cathode) and CGR18650E (LCO cathode) at 45 °C (reprinted from ref 61, Copyright 2014, with permission from Elsevier); (b) cycle performance of model cells under four ΔDOD conditions at 25 °C: (●) 0−60, (◆) 10−70, (■) 40−100, and (▲) 0−100% (reprinted from ref 62, Copyright 2014, with permission from Elsevier); (c) average power and average % power fade as a function of cycle-life test time (i.e., number of cycle-life test cycles) for ATD Gen 2 Baseline cells tested at 45 °C. Fits of the data to a square root function of the test time are shown. Test data are shown for a test period of 44 weeks. The rate of change in power and % power fade increases between the 28 and 32 week test periods (reprinted from ref 67, Copyright 2003, with permission from Elsevier); (d) average power of NCA cells at 300 Wh vs time (reprinted from ref 69, Copyright 2003, with permission from Elsevier); (e) EIS data (10 kHz− 0.01 Hz) at 25 °C for the full cell and positive and negative electrodes harvested from Gen2 and VarC cells (reprinted from ref 76, Copyright 2007, with permission from Elsevier). 201


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Figure 5. (a) In situ synchrotron X-ray diffraction patterns at selected 2θ regions for the first charge (top) and second charge (bottom) for NCA with Bragg peaks according to space group R3̅m. R1 and R2 denote phase 1 and phase 2, respectively; x denotes the amount of Li that is deinserted from Li1−xNi0.80Co0.15Al0.05O2 (reproduced from ref 90); (b) HRTEM image, convergent beam electron diffraction patterns, and structure schematics illustrating the crystal structure difference between the bulk and surface of LiNi0.8Co0.2O2 particles (reproduced with permission from ref 70, Copyright 2003, The Electrochemical Society); (c) image intensity ratios of Li layers to neighboring transition metal layers showing the increase in transition metal ion concentration at the Li sites. The ratios were measured for the original and the single-cycled NCA samples (reproduced with permission from ref 83, Copyright 2011, The Electrochemical Society); (d) top: cross section of an embedded particle in a coin cell after one charge−discharge cycled NCA in which cracking is observed throughout the particle and cracks in several locations are indicated by arrows; bottom: cross section of an embedded particle in a coin cell after 4500 charge−discharge cycles showing extensive cracking, similar to that observed in single particles (reproduced from ref 91 with permission, Copyright 2013, Wiley−VCH Verlag GmbH & Co. KGaA); (e) calorimetry profiles (at 100% SOC) for an 18650 cell with a Mag-10 graphite anode and a LiNi0.8Co0.15Al0.05O2 cathode in 1.2 M EC/ EMC (3:7 by weight) electrolyte: ARC profile showing the three stages of thermal runaway (reprinted from ref 98, Copyright 2006, with permission from Elsevier). 202


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suggestion, Lee et al. 56 found a similar composition, LiNi0.8Co0.15Al0.05O2, which preserved a stable phase up to a highly oxidized state. These cathode materials delivered almost the same discharge capacity of approximately 200 mAh g−1. Later work by Gulimard et al.53 found a quasi-ideal lamella structure with no extra nickel ions in the interslab space (less than 1%). However, a higher level of Al3+ in the transition metal layer interrupted facile expansion and contraction of the crystal structure during Li+ extraction and insertion. This resulted in a smaller capacity than that of the above two materials, which was ascribed to lack of electroactive Ni3+ species in the transition metal layer. These findings confirm that 5% Al3+ in the transition metal layer is empirically sufficient to stabilize the layer character of LiNi1−xCoxAl0.05O2. Provided that the electrochemical properties are similar in terms of capacity, retention, and rate capability, a lower Co content is preferable because Ni is costeffective compared to Co. Therefore, it is deduced that the optimal composition is LiNi0.8Co0.15Al0.05O2, the original crystal structure of which is maintained in the voltage range of 2.5−4.3 V. Recent results indicate that increasing the solubility of Al to 9% in the transition metal layer, LiNi0.81Co0.1Al0.09O2, enables extension of the upper voltage cutoff to 4.5 V, with good cycling performance at 60 °C.57 It is generally known that Al replacement at the transition metal site is not easy because of the formation of undesired oxide phases such as β-LiAlO2, Al2O3, and γ-LiAlO2 when calcined above 600 °C. This is because the phase formation reaction rate of Al progresses slower than those of Li, Ni, and Co when a conventional solid-state reaction method is employed. For this reason, coprecipitation,39,52 sol−gel,58 or solution-based synthetic methods59 are preferred to obtain phase-pure NCA. Because α-LiAlO2 (a = 2.800 Å, c = 14.22 Å) has a smaller a-axis and larger c-axis than LiNiO2 (a = 2.88 Å, c = 14.19 Å) and LiCoO2 (a = 2.88 Å, c = 14.19 Å),40,60 anisotropic change of the lattice parameters is natural, provided that a solid solution is formed. 2.1.2. Performance of NCA Cathode Batteries. Cell Performances. Watanabe et al.61 compared cell performances of NCA cathode batteries (NCR18650) and LCO cathode batteries (CGR18650E) (Figure 4a) at 45 °C. The NCA cathode batteries exhibited obviously better capacity retention during the cycling test at 45 °C, although the retentions were similar for the cells when cycled at 25 °C. This indicates better suitability of the NCA cathode batteries for vehicles and energy storage applications. They also cycled the NCA cathode batteries at 25 °C in four ΔDOD (depth of discharge) conditions (Figure 4b). The capacity fading in the ΔDOD range of 0−100% was significantly large compared to that in the other limited ΔDOD conditions. There was almost no notable difference in the degradation behavior of discharge capacity among the three limited ΔDOD conditions, in which the tested voltage ranges were different though the DOD condition was fixed to 60%. Fixing the ΔDOD to 60%, the resulting cycling performances could be significantly enhanced even at 60 °C, maintaining more than 90% of the initial capacity after 2500 cycles. In addition, in the condition of 100% ΔDOD, significant capacity fading has been observed.61−89 These results suggest that the capacity fade behavior is highly dependent on the degree of ΔDOD, not on the upper or lower limits of DOD even at elevated temperatures. Power Fading Phenomena of NCA Batteries. The LiNi0.8Co0.15Al0.05O2 cathode (Gen 2) and LiNi0.8Co0.10Al0.10O2 cathode (Variant C) were tested by Argonne National Laboratory as a part of the Advanced Technology Development

(ATD) program (Figure 4c,d). The present study adopted the same cell chemistries as in that study except for the cathode materials. Figure 4c shows the change in average power and the average % power fade of Gen 2 through 44 weeks of cycling at 45 °C. Power fade had a relation with time, and the rate at 45 °C was almost twice that at 25 °C for a period of 28 weeks, after which the fade further accelerated. Similar power fade was also observed for Variant C, although the cell exhibited the highest power during the test period (Figure 4d). Similar to the power fade, impedance rise was confirmed for both Gen 2 and Variant C (Figure 4e). A slight decrease in the impedance was noted in the positive electrode of Variant C using the material with a higher amount of Al, LiNi0.8Co0.10Al0.10O2. This indicates that the cathode is the main factor that affects the impedance, namely, progressive increases in charge transfer and mass transfer at the interface between the cathode and electrolyte. According to their further experiment using rinsed electrodes harvested from an aged cell, the anode returned to its original capacity, while the cathode failed to deliver a capacity as high as that in its fresh state. This indicates that the cathode is mainly responsible for the power fade in NCA cathode batteries. 2.1.3. Degradation of NCA Materials during Li + Extraction/Insertion. During Li+ extraction during charging, the structure of Li1−δNiO2 is basically maintained. However, production of monoclinic and another hexagonal structure in the intermediate and deeply delithiated states is inevitable due to the intrinsic character of LiNiO2.13 These repetitive structural changes are not desirable because the evolution can eventually cause breakdown of the crystal structure during long-term cycling. To minimize such detrimental structural variation, earlier works focused on partial substitution at the Ni site by other elements, such as Co3+ or Al3+.33,34,52 Ohzuku et al.41 and Delmas et al.52 observed a single-phase reaction in Li1−δNi1−xAlxO2 and Li1−δNi1−xCoxO2 during Li+ extraction from the host structure, respectively, enabling structural stabilization by the effect of Co3+ or Al 3+ . Such an effect is also observed in Ni-rich LiNi0.8Co0.15Al0.05O2, even though a two-phase reaction, hexagonal 1 and hexagonal 2 phases, is associated with Li+ extraction and insertion. Very recently, Robert et al.,90 however, insisted that the two-phase reaction is observed only at the first cycle (Figure 5a), after which Li+ extraction from LiNi0.8Co0.15Al0.05O2 is related to a reversible single-phase reaction in the second cycle, as confirmed by in situ synchrotron X-ray diffraction. This reversibility is a possible reason for the better capacity retention of NCA electrodes relative to that of LiCoO2 (Figure 4a), indicating that NCA electrodes should have shown good capacity and power retention in consideration of their structural stability. Therefore, the capacity and power fades are related to other factors rather than structural reversibility. From a LiNi0.8Co0.2O2 electrode demonstrating a 43% fade in power, Abraham et al.70 found the formation of NiO-like or LixNi1−xO-like (rock-salt structure) byproducts (Figure 5b) on the surfaces of LiNi0.8Co0.2O2 particles. NiO is known to have poor conductivity of lithium ions and electrons, and a resistance layer is formed on the primary particles. This surface change would produce an increase in charge transfer resistance during cycling (Figure 4e), which was confirmed by neutron and synchrotron XRD studies. Shikano et al.78,79 correlated the power fade and cationic disordering of the NCA electrode. According to their XANES and XPS investigation, the valence of Ni ions did not vary with degradation, and the structural change from the layer structure to the lithium-deficient cubic phase, probably, LixNi1−xO, was located near the surface of the NCA 203


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materials. Sasaki et al.80 and Muto et al.81 found that an inactive NiO-like phase was detected near the grain surface and boundaries in NCA based on a spatial distribution map. Consecutive microscopic work by Zheng et al.83 demonstrated that the region with a disordered rock-salt structure, presumably a NiO-like or LixNi1−xO-like phase, was about 5 nm in width, while the transition region with a partially ordered structure was about 20 nm in width (Figure 5c) at the first cycle. A substantial volume of material with the transitional structure was formed near the grain boundary. These researchers confirmed this result using a conducting agent and binder-free NCA electrode.85 Another aspect is the formation of Li2CO3 on the surface of NCA particles. Earlier work by Kostecki et al.72,74 found that the concentration of surface carbon species increased on the aged NCA electrode. This also caused interparticle resistance to dominate cathode interfacial charge transfer resistance and accounted for power and capacity fades. Later work by Kobayashi et al.77 suggested that the carbon species was a film of Li2CO3 decomposed from the electrolyte, the formation of which was accelerated at elevated temperature. The formation of a Li2CO3 layer was more evident when the structure of NCA was transformed to a lithium-deficient cubic structure on the surface. Hence, decomposition of the layer structure due to a deficiency of lithium on the surface of NCA is linked to the formation of Li2CO3 and to the power fade. Zhuang et al.71 intentionally exposed NCA particles to air to produce Li2CO3 layers. The insulating layer isolated NCA particles, greatly lowering both capacity and power. Recently, morphological change of NCA particles was suggested as a possible cause of capacity and power fade.79−86,91 Such particle cracking has been associated with anisotropic changes in the lattice, resulting from the extraction and insertion of lithium ions. Substitution of Mg for metal elements in the NCA material can reduce such lattice change.54 Miller et al.91 investigated NCA particles after 1 and 4500 cycles (Figure 5d). The particle underwent intergrain cracking after the first cycle. Almost all grain boundaries showed evidence of separation (Figure 5d-1). Correlating the above results, it is likely that the formation of NiO-like and Li2CO3 phases on the surfaces of individual primary particles accelerates the particle cracking even at the first cycle. Electrolyte can penetrate along the propagated cracks of the secondary particles. Thus, each primary particle is always exposed to electrolyte. Indeed, LiPF6-based electrolyte always contains a small amount of water as an impurity. The existence of a small amount of water consequently causes breakdown of the electrolyte accompanied by acceleration of HF generation. This means that, during extensive cycling, the active materials are always exposed to HF. The produced HF continuously attacks the active materials and causes their decomposition. Transition metal ingredients are gradually dissolved into the electrolyte and tend to adhere on the surfaces of the cathode and anode. Although the extended grain boundary layers such as NiO and Li2CO3 formed at the first cycles act as protection layers to prevent degradation of morphology during cycling, these inorganic layers are insufficient to protect active materials during extensive cycling. As a result, the cracking was more obvious by electrochemical cycling after 4500 cycles (Figure 5d-2). Furthermore, the destroyed particles lessen the electrical contact between the active material and conducting agent, so that capacity fading and impedance increase readily occur during cycling. To minimize such morphological degradation, it is very important to control the inner surface area of the secondary

particle. Recent reports by Sun et al.92−97 suggested a radial alignment of nanorod primary particle assembly in a full concentration gradient (FCG) secondary particle of NCM. This configuration of nanorods ultimately reduces voids in the spherical particle to minimize the contact area with the electrolyte, which substantially decreases the side reactions with the electrolyte during cycling. Although these approaches have not been applied in a NCA system, efforts are necessary to incorporate these NCA materials in vehicles and energy storage applications. 2.1.4. Safety. Safety is another important issue that has to be considered for widespread commercialization of materials in vehicle applications. Cells are composed of organic electrolytes and are prone to thermal runaway at temperatures above 180 °C (Figure 5e).98 In particular, for deeply delithiated cathode materials, transition metal substructures are exothermically decomposed at that temperature due to the release of oxygen from the crystal structure. The removal of oxygen from the delithiated crystal structure, which occurs at the onset temperature of the exothermic reaction, causes structural reorganization via the continuous migration of cations to more stable sites such as vacant tetrahedral lithium sites,78,99,100 and the structure is finally stabilized into the cubic spinel. At this point, the violent exothermic reaction is terminated below 300−350 °C for delithiated NCA materials. Further heating induces the formation of rock-salt (Fm3m) phase, and the structure does not return to its original layer structure after cooling to room temperature. Although the occurrence of exothermic behavior is inevitable, many efforts have been made to reduce such thermal decomposition for NCA materials. Incrementing the Al content while decreasing the Ni and Co concentration or modifying the surface using electroinactive heterogeneous thin layers is a good approach to improve safety-related characteristics.101 2.2.1. Development of NCM Materials. In 2001, Ohzuku et al. revisited lithium nickel manganese oxides with or without cobalt as possible alternatives to LiCoO 2 . 102,103 LiNi1/3Co1/3Mn1/3O2 consists of Ni2+, Co3+, and Mn4+,104 while LiNi1/2Mn1/2O2 is composed of Ni2+ and Mn4+.105 Although similar discharge capacities can be obtained from both materials, less occupation of Ni2+ in the Li layer, cation exchange, and improved reversibility and polarization are observed in LiNi1/3Co1/3Mn1/3O2. The presence of Co in the transition metal improves such properties as electric conductivity, which fosters the high rate capability. Later, MacNeil et al.106 studied a solid solution of LiNixCo1−2xMnxO2 (x = 0−0.5), in which both the a- and c-axis parameters decreased linearly with increasing Co content. Despite the extensive study, a single optimal composition has not yet been found for the NCM system. It is generally understood that increasing the Ni content will increase the initial discharge capacity but incurs severe capacity fading during cycling. In addition, Ni-rich NCM compounds are susceptible to thermally induced phase transition to spinel and rock-salt structures, which is accompanied by oxygen release. The released oxygen can lead to thermal runaway of the battery, especially at the overcharged state.107−109 Increasing the Co content substantially reduces the capacity loss during cycling at the expense of significantly increased material cost and reduction in capacity. The Mn content particularly improves the thermal stability, but Mn doping substantially reduces the discharge capacity of NCM. To clarify the role of each transition metal in NCM, a comprehensive electrochemical evaluation of LiNi1/3Co1/3Mn1/3O2, LiNi0.5Co0.2Mn0.3O2, LiNi0.6Co0.2Mn0.2O2, LiNi 0.75 Co 0.15 Mn 0.15 O 2 , LiNi 0.8 Co 0.1 Mn 0.1 O 2 , and Li204


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Figure 6. (a) Calculated open-circuit voltage versus capacity for LiNi1−xMO2 (M: Co and Mn) (reprinted from ref 110, Copyright 2016, with permission from Elsevier); (b) cyclability of Ni-rich Li[Ni0.85CoxMn1−x]O2 in the voltage range of 2.7−4.5 V measured at a current of 0.1 C at 30 °C (reprinted from ref 111, Copyright 2016, with permission from Elsevier); (c) capacity retention of a graphite/Li[Ni0.6Co0.2Mn0.2]O2 full cell with different electrolytic additives, VC: vinylene carbonate, SN: succinonitrile, PST: propene sulfone, PS: propane sulfone (reprinted from ref 118, Copyright 2014, with permission from Elsevier); (d) discharge curves of the Li/Li[NixCoyMnz]O2 cells with x = 1/3 (left top), 0.5 (right top), 0.6 (left middle), 0.7 (right middle), 0.8 (left bottom), and 0.85 (right bottom) as a function of C rate from 0.2 to 5 C (reprinted from ref 107, Copyright 2013, with permission from Elsevier); (e) continuous discharge capacity of Li[Li0.333Ni0.855Co0.045Mn0.067]O2 and Li[Li0.05Ni0.807Co0.043Mn0.1]O2 at 55 °C (reproduced with permission from ref 124, Copyright 2015, The Electrochemical Society).

Ni0.85Co0.075Mn0.075O2 was carried out.107 Of the tested cathodes, LiNi1/3Co1/3Mn1/3O2 performed best in terms of capacity retention and thermal safety; however, its discharge capacity was limited to 160 mAh g−1. Increasing the relative fraction of Ni increased the discharge capacity as LiNi0.85Co0.075Mn0.075O2 exhibited an initial discharge capacity greater than 200 mAh g−1 in the voltage range of 2.7−4.3 V, which is comparable to that

of NCA materials. This was empirically demonstrated by Min et al.110 using density functional calculation (DFT). Specifically, when the same amount of Li+ ions was extracted, the cathode with more Ni exhibited the lower voltage value (Figure 6a). Unfortunately, the resulting capacity retention of Ni-rich LiNi0.85CoxMn0.15−xO2 (x = 0−0.15) was disappointing, exhibiting only 65% of the initial discharge capacity after 100 205


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Figure 7. (a) Cycling performance of 18650 graphite/Li[Ni1.3Co1/3Mn1/3]O2 high-power batteries with a charge rate of 1 C and discharge rates of 5 and 10 C at 25 °C (reprinted from ref 125, Copyright 2011, with permission from Elsevier); (b) capacity retention as a function of cycle number for graphite/NCM full cells, A: NCM523, B: NCM433 (N/P ratio: 1.06), and C: NCM433′ (N/P ratio: 1.19) (reprinted from ref 126, Copyright 2014, with permission from Elsevier); (c) discharge curves of the LTO/NCM cell at various rates; and (d) cell life performance of the LTO/NCM cell for charge−discharge cycling at 3 C between 1.5 and 2.7 V (reprinted from ref 127, Copyright 2013, with permission from Elsevier); (e) X-ray diffraction patterns of an extensively cycled graphite electrode recovered after 1500 cycles operated at 45 °C/60 °C and a commercial Kokam cell graphite/NCM cell (reprinted from ref 128, Copyright 2015, with permission from Elsevier); (f) ds/dQ and dV/dQ at various C rates on the center of the surface of the graphite/NCM batteries for discharge (reprinted from ref 130, Copyright 2014, with permission from Elsevier).

cycles (Figure 6b).111 Therefore, it is imperative to obtain the optimal voltage window with desired capacity and control it to prevent excessive structural changes in Ni-rich cathode materials. Another way to potentially increase the discharge capacity from the NCM cathodes is to increase the upper cutoff voltage above 4.3 V; however, fast deterioration was observed when cycled at 4.5 V mostly due to an increase in charge transfer

resistance caused by HF attack on the cathode surface.112 LiNi0.4Co0.2Mn0.4O2 has also been charged to 4.7 V but experienced rapid capacity loss upon cycling.113 Similar to LiCoO2, LiNi0.4Co0.2Mn0.4O2 required a surface coating layer to improve the capacity retention at high voltage cycling. LiNi1/3Co1/3Mn1/3O2, which exhibits optimal performance in terms of capacity retention and thermal safety, was been charged 206


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Ni0.5Co0.2Mn0.3O2 and LiNi0.4Co0.3Mn0.3O2 at 1 C (Figure 7b). The NCM523 cell exhibited faster capacity fading than the NCM433 cell. Even though the long-term cycling performance suggests satisfactory structural and chemical performance of both active materials, higher Ni content in the active material evidently accelerates the capacity fading during cycling. Takami et al.127 demonstrated the superiority of LTO/NCM cells in terms of capacity and energy retention, rate capability, power capability, self-discharge, and so on. As can be seen in Figure 7c, the operation voltage of the LTO/NCM cell (20 Ah) is lower than that of the graphite/NCM cell due to the higher operating voltage of the LTO anode. However, the resulting capacity at 8 C was almost the same as that obtained at 1 C. The cell could be operated even at −30 °C, showing 80% of the capacity obtained at 25 °C. The expected capacity retention after 6000 cycles was approximately 86% (Figure 7d). Jalkanen et al.128 tested commercial graphite/NCM cells with a nominal capacity of 40 Ah, fabricated by Kokam, whose capacity retention after 2600 cycles was approximately 80% (32 Ah). This comparison indicates that graphite/NCM cells are acceptable for EV application at cold temperatures. Power Fading Phenomena of NCM Batteries. As shown in section 2.1.2, the main difference between NCA and NCM batteries is the use of Al and Mn, respectively. It is interesting to note that only 5% Al3+ is needed in the transition metal layer to stabilize the layer structure, while 20−33% Mn4+ is adopted in the crystal structure for NCM622 and NMC111. Our literature survey concluded that the cathode is the main problem for the power fade of NCA batteries. Later works on NCM batteries investigated other parameters that affected performance deterioration. Wang et al.129 found that capacity loss mainly occurred in the NCM cathode, 4% at 10 °C after 2000 cycles, while the graphite anode experienced only 1% loss. At 46 °C, however, the capacity loss was accelerated, 11% for the NCM cathode and 5% for the graphite anode. They also observed fast capacity fading when the cell was cycled at high rates. Graphite material loss was accelerated by high discharge rates, while cathode loss was moderate for all rates. Overall, the loss of active lithium outpaces both the cathode and anode materials, indicating that the capacity is limited by the amount of useable lithium. As a result, the material loss and loss of available lithium inventory are responsible for resistance increase upon cycling. Jalkanen et al.128 observed the formation of Li2CO3 and LiF from the extensively cycled graphite anode surface as a result of formation and growth of an SEI layer (Figure 7e), which supports the above explanation on loss of available lithium inventory. The N/P ratio is another important factor because the anode is charged to a lower state of charge (SOC) at a higher N/P ratio, which results in less volume change or less strain during cycling. Hence, a higher N/P ratio can reduce the formation of Li2CO3 and LiF, so that swelling of the cell can be delayed in pouch-type cells.126 Oh et al.130 observed swelling upon charge and contraction during discharge (Figure 7f): The contraction observed upon discharge was highly dependent on the C rate. Therefore, further effort is necessary to minimize the cathode material loss and stabilize the electrolyte to prevent the sedimentation of Li2CO3 and LiF toward the surface of the anode, which causes chain reactions such as impedance increase, capacity and power (energy) fading, and swelling. It is thought that one of the main reasons for the above-mentioned degradation is failure of cathode materials after reacting with electrolyte. Unfortunately, Ni-rich NCM cathode materials are

to 4.9 V to further extract the capacity; even the chemically stable LiNi1/3Co1/3Mn1/3O2 showed poor first cycle efficiency and a large increase in impedance at 4.5 V, largely affected by electrolyte oxidation on the cathode surface.114−116 Gallus et al.117 compared several electrolytic salts to enhance cycling stability at high voltage. They demonstrated better cycling performances when using LiBF4 and LiClO4 salts instead of LiPF6 because of reduction in metal dissolution of LiNi1/3Co1/3Mn1/3O2 at the expense of capacity. Selection of electrolyte additive is important to retain a high capacity. Kang et al.118 suggested that 1,3-propane sulfone was highly desirable for enhancing the electrochemical performance and decreasing swelling, metal dissolution, and gas generation at the fully charged state for a graphite/LiNi0.6Co0.2Mn0.2O2 full cell (Figure 6c). Thin and weaker formation of LiF in the surface film of LiNi1/3Co1/3Mn1/3O2 was achieved by addition of trimethylboroxine (TMB) and produced excellent cycling performance when cycled in the voltage range of 3−4.5 V.119 Recently, Dahn’s group120 reported the successful passivation of NCM 811 using diphenylcarbonate (DPC), retaining more than 91% of the first discharge capacity. Hence, it is likely that a significant breakthrough in the high-voltage electrolyte121 and ionic liquid122 or economically viable surface coating scheme are required to widen the operating voltage window of NCM cathodes. Noh et al.107 have reported the rate capability of NCM cathodes as a function of Ni content and showed that the rate capability substantially improved with increasing Ni content (Figure 6d). In fact, the electronic conductivity of LiNi0.85Co0.075Mn0.075O2 (2.8 × 10−5 S cm−1) was 3 orders of magnitude higher than that of LiNi1/3Co1/3Mn1/3O2 (5.2 × 10−8 S cm−1). The galvanostatic intermittent titration technique was used to confirm that Li+ diffusivity for LiNi0.85Co0.075Mn0.075O2 was indeed 10−8 cm2 s−1, while that of LiNi1/3Co1/3Mn1/3O2 was limited to 10−11 cm2 s−1; hence, Ni enrichment of the NCM cathodes tended to enhance their rate capability. Zheng et al.123 highlighted the importance of the surface structural stability of Ni-rich cathode materials at high voltage cutoff; Mn ions in the Ni-rich cathode material significantly reduced the irreversible side reactions between the electrode surface and electrolyte so that the surface structure was stabilized, ensuring good cycling stability. Choi et al.124 also found formation of solid solution for Ni-rich compounds such as (1−x)LiNi 0.95 Co0.05O2 −xLiLi0.33Mn0.67O2 (Figure 6e). An interesting finding is that the average oxidation states of Ni and Mn were +3 and +4, respectively. As a result, the material (217 mAh g−1 at the first cycle) could maintain 90% of its capacity retention after 100 cycles in the voltage range of 2.7−4.5 V. This result indicates that tetravalent Mn ions are substantially effective in stabilizing the electrochemical performance, even in small amounts. Therefore, it is imperative to obtain the optimal voltage window with desired capacity to prevent excessive structural changes in Ni-rich cathode materials. 2.2.2. Performance of NCM Cathode Batteries. Cell Performances. Structural stability is related to long-term cycling performances. Figure 7a shows the cycling stability of a graphite/ LiNi1/3Co1/3Mn1/3O2 full cell operated in the voltage range of 2.7−4.2 V.125 Even at a high charge rate (5 C), the NCM cell was able to deliver a high discharge capacity at a high rate discharge rate (10 C). This indicates that the NCM cell is suitable for highpower applications. Even at elevated temperature (50 °C), similar cell performances were observed. Liu et al.126 compared cycling performances of the graphite/NCM cells Li207


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Figure 8. Dissolved concentrations of Ni, Co, and Mn elements in 1 M LiPF6/EC+DEC (1:1) electrolyte when the NCM electrode was polarized to different potentials (reprinted from ref 135, Copyright 2012, with permission from Elsevier); (b) Li[Ni0.4Mn0.4Co0.18Ti0.02]O2 that shows reconstruction of the surface after one cycle (2−4.7 V) (reproduced with permission from ref 143, Copyright 2014, Macmillan Publishers Limited); (c) high-resolution transmission electron microscopic image analysis of the structural degradation in Li[Ni0.6Co0.2Mn0.2]O2 cathode materials after 5 cycles with the image color-overlaid based on the structural transition from the interior to the surface region (top) and the depth of disorder structure during 10 cycles (bottom); and (d) propagation of cracks observed after 2500 cycles of a graphite/Li[Ni0.6Co0.2Mn0.2]O2 full cell (reprinted from ref 144, Copyright 2016, with permission from Elsevier).

Li[Ni0.4CoxMn0.6−x]O2.131 Chemical reaction with the electrolyte leads to the formation of a surface film, acting like a passive layer, on the surfaces of active materials, which impedes Li+ diffusion from the bulk to the electrolyte. Cherkashinin et al.132 explained that oxidative Co4+ and Ni3+/4+ in a highly delithiated state, which were dissolved during the electrochemical reaction, are inherently related to the formation of the passive layer. With increasing Ni content, Co3+ and Mn4+ persisted while Ni4+ tended to reduce to Ni3+.133 In general, the presence of stable Mn4+ in the NCM cathodes reduces the transition metal dissolution compared to that of layered Li[Mn0.8Cr0.2]O2 or orthorhombic LiMnO2, in which the average oxidation of Mn is 3+.134 For LiNi1/3Co1/3Mn1/3O2, the dissolution rates of Ni, Co, and Mn were determined in transient mode in a conventional 1 M LiPF6/EC+DEC (1:1) electrolyte for 4 weeks.135 There were abrupt increases in the dissolutions of Ni, Co, and Mn at

still under investigation for use in commercial LIBs due to an important technical issue of residual lithium compounds such as Li2CO3 and LiOH on the surfaces of active materials. Provided that this issue is resolved, the use of Ni-rich NCM would be available for high-energy LIBs. Therefore, robust cathode materials with high tap density, high rupture strength, low specific surface area, and low inner pore volume are expected to resolve the aforementioned difficulties of NCA and NCM batteries, which will be mentioned later in section 2.2.3. 2.2.3. Degradation of NCM Materials during Li+ Extraction/Insertion. Mechanisms of capacity fading of the NCM cathodes have been attributed to transition metal dissolution and intermixing of Li and Ni ions from their respective lattice sites. X-ray absorption near the edge structure of the NCM cathode has verified that the oxidation state of each transition metal in NCM cathodes is Ni2+, Co3+, and Mn4+ in 208


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Figure 9. (a) Correlation among discharge capacity, thermal stability, and capacity retention for Li1−δ[NixCoyMnz]O2 materials (x = 1/3, 0.5, 0.6, 0.7, 0.8, 0.85) (reprinted from ref 107, Copyright 2013, with permission from Elsevier); (b) mass spectroscopy profiles for the oxygen (O2, m/z = 32) collected simultaneously during measurement of TR-XRD and the corresponding temperature region of the phase transitions for NMC samples (lower panel) (reproduced from ref 145); (c) particle-to-particle variations in the transition metal L-edges of charged NMC cathode materials, expressed by L3/L2 intensity ratios with increasing temperature (reproduced from ref 146).

ionic sizes (0.76 Å for Li+ and 0.69 Å for Ni2+). It seems that the effect of Li/Ni mixing is linked to the deterioration of electrochemical performance. Wu et al.142 suggested that a low degree of Ni2+ disorder is indeed helpful to enhance the electrochemical performance of LiNi0.8Co0.1Mn0.1O2. Controlling and reducing Li/Ni mixing remains a challenging task, especially for Ni-rich NCM cathodes. Similar to NCA materials, Lin et al.143 recently observed that NCM material underwent surface reconstruction after exposure to the electrolyte; the surface of this material is composed of cubic rock-salt Fm3̅m and layer R3̅m structure (Figure 8b). The surface reconstruction was promoted by Li+ removal during charging and was more severe at the outermost surface region. This disorder further progressed with cycles (Figure 8c).144 Indeed, these NCM materials underwent repeated anisotropic expansion and contraction during cycles. As a result, considerable

potentials above 4.5 V (Figure 8a). The Mn dissolution was considerably more severe than those of Co and Ni, with minimal Ni dissolution observed up to 4.6 V. Because the Mn dissolution mostly occurred above 4.5 V, it is surmised that the dissolution was caused by acid corrosion (HF from the electrolyte breakdown). Meanwhile, recent work by Gallus et al.117 has shown that using LiBF4 instead of LiPF6 substantially reduced the metal dissolution of LiNi1/3Co1/3Mn1/3O2, which cycled at 4.6 V with low capacity fading using LiBF4. In the Ni-rich end of NCM cathodes, because of the high Ni fraction, Ni dissolution becomes more serious compared to that of Co or Mn.136 As for cation mixing, occupation of Ni ions in the Li layer in the NCM cathodes can potentially impair the electrochemical performance by blocking the Li+ diffusion path.133,137−141 In NCM cathodes, especially Ni-rich NCM cathodes, disordered distribution of Li and Ni in 3a and 3b sites can easily occur because of the similar 209


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order to avoid decomposition of electrolyte solvent as a catalyst. These coating materials (Al2O3, ZnO, ZrO2, AlPO4, AlF3) have very poor electronic conductivity and Li+ ion conductivity. In general, a thick coating of such materials produces better protection but poor power capability. Despite the insulating properties of the coating materials, a smooth, thin coating layer (5−10 nm in thickness) showed better performance than a thick coating layer.155 Of the electrolyte salts, LiPF6 is the most widely used in lithium-ion batteries. However, it can readily decompose to produce HF (PF5 + H2O → POF3 + 2 HF) in the presence of moisture.156 The generated HF readily attacks the layered oxide cathode surface to produce gradual dissolution of metal ions (Ni, Co, and Mn) from the particle surface; the edges of primary particles were shown to be severely serrated and thus increased interfacial impedance upon cycling.157 When amphoteric metal oxides are uniformly coated on the cathode surface, they act as HF scavengers and greatly reduce the HF concentration in the electrolyte near the cathode surface.148,158 As an example, ZnO,148 Al2O3,155,159 and Li3PO4160 coating layers scavenge the acidic HF in the electrolyte according to the following respective reactions because the Gibbs formation energies of the metal fluorides are much lower than those of metal oxides161

strain may cause volume shrinkage of the particles, followed by pore formation between primary particles. Electrolyte is subject to infiltration into the pores, and dissolution of the active materials near the grain boundaries results in isolation of the primary particles (Figure 8d), which further degrades the interior of the secondary particles and causes degradation of the electrical properties. 2.2.4. Safety. In the NCM system, as summarized in Figure 9a, the LiNi1/3Co1/3Mn1/3O2 cathode exhibited the best capacity retention and thermal stability among the LiNi1−2xCoxMnxO2 compositions.107 By contrast, its discharge capacity was limited due to the smallest amount of Ni in the composition. The LiNi1−2xCoxMnxO2 (x = 0.075) cathode that contained the largest Ni content experienced severe capacity fading. The capacity fading mechanism is believed to be due to volume expansion of the electrode during cycling and surface structural degradation (transformation to the rock-salt structure). In a highly oxidized (charged) state, the resulting unit cell volume is contracted, which is related to phase transition from the hexagonal 1 to hexagonal 2 phase. The instability of the crystal structure in highly delithiated states directly reflects the thermal stability, and this, in turn, triggers the phase transition, releasing oxygen from the crystal structure toward the cubic spinel structure. The onset temperature of the exothermic reaction is the initiation temperature of oxygen removal from the crystal structure, and the reaction occurs at lower temperature for high Ni-rich compounds due to structural instability at the highly delithiated state (Figure 9b).99,107,145 As seen in Figure 9b,c, the poor thermal property of LiNi0.8Co0.1Mn0.1O2 relative to LiNi0.4Co0.3Mn0.3O2 is caused by the large fraction of Ni4+ present in the charged state.146 By contrast, the presence of Mn4+ is very helpful to stabilize the crystal structure. Hence, the composition and microstructure of the LiNi1−2xCoxMnxO2 needs to be further optimized or modified by heteroelements, in particular, on the surface, to attain an ideal cathode material with high capacity and thermal stability.

The metal oxide layer gradually transforms to metal fluoride layers due to the selective scavenging of HF in the electrolyte upon cycling; this reduces the acidity of the electrolyte, leading to much less degradation of the cathode surface (Figure 10a).159 In comparison, the metal fluoride layer will be thickened with cycling due to the byproducts of reactions 1148 and 2155 and act as

3. SURFACE MODIFICATION OF NCA AND NCM MATERIALS Surface coating of cathode materials has shown to be effective to improve the electrochemical performances in terms of capacity retention, rate capability, and thermal stability of lithium-ion batteries without sacrificing their initial capacity. The coating materials studied to date are metal oxides,14,16,147,148 metal phosphates,149 metal fluorides,150,151 metal oxyfluorides,152 and metal hydroxides.153 The role of coating materials is to act as (1) a protective layer that reduces the parasitic reaction between the cathode surface and liquid electrolyte154 and (2) a HF scavenger of amphoteric metal oxide (Al2O3, ZnO, TiO2, ZrO2, etc.) that suppresses transition metal (Ni, Co, and Mn) dissolution from the cathode surface into electrolytes. The principal role of the surface coating layer is physical protection of the cathode surface to prevent parasitic chemical reactions between the cathode surface and liquid electrolyte. These reactions are accelerated by direct contact of reactive Ni4+ and Co4+ under a highly delithiated cathode state with electrolyte; this exposure leads to chemical oxidation of the electrolyte solvent and degradation of the cathode surface. Coating of electrochemically and chemically inactive material on the cathode surface can act as a physical protection barrier to suppress the parasitic reactions and hence improve the electrochemical performance of lithium-ion batteries.154 The general feature of the coating materials is that the metal in the coating layer has only one stable valence state in

Figure 10. (a) Schematic drawing of the interface between the cathode and electrolyte (adapted from ref 159); (b) schematic illustration of chemistry on the surface of bare (left) and lithium phosphate-coated Li[Ni0.6Co0.2Mn0.2]O2 (right) (reproduced with permission from ref 160, Copyright 2015, Springer).

ZnO + 2HF → ZnF2 + H 2O

(1)

and

210

Al 2O3 + 6HF → 2AlF3 + 3H 2O

(2)

Li3PO4 + HF → LixHyPO4 (or POx Hy) + LiF

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Figure 11. (Left) Schematic drawings of a core−shell (Gen 1) (adapted from ref 168); (center) a Ni-rich core surrounded by a concentration gradient outer layer (Gen 2) (adapted from ref 170); (right) a FCG lithium transition metal oxide particle (Gen 3) with the nickel concentration decreasing from the center toward the outer layer and the concentration of manganese increasing accordingly (adapted from ref 92).

surfaces of these active materials, presumably as the oxide form Li2O. Also, the outer part of the Li2O is contaminated with moisture and CO2 from air, forming LiOH and Li2CO3. The residual lithium compounds usually range in concentration from 6000 to 25 000 ppm when the Ni content is greater than 60% in the transition metal layer.107 More Ni in the transition metal layer results in formation of a larger amount of residual lithium compounds. These compounds can cause gelation of the slurry in electrode preparation. Even when the electrode was successfully fabricated, these compounds underwent oxidative decomposition at high voltage, generating dangerous gases. Therefore, the level of residual lithium compounds should be lowered to an acceptable level, 3000 ppm or lower. Jo et al.160 successfully reduced the concentration of residual lithium compound in LiNi0.6Co0.2Mn0.2O2 by modifying the surface with H3PO4 to produce a Li3PO4 coating layer. Specifically, the added starting H3PO4 was transformed to Li3PO4 during heating, and three Li atoms were trapped to produce Li3PO4 from the starting H3PO4 (Figure 10b), decreasing the amount of residual lithium compounds after coating. Interestingly, the Li3PO4 layer scavenges not only HF (reaction 3160) but also water molecules present in the electrolyte160

a protective coating layer to suppress degradation of the cathode surface. The formed metal fluoride coating layer is so resistant to HF attack that the cathode surface is well preserved. However, once the metal oxide coating layers are reacted with HF, water is generated and facilitates generation of HF in the electrolyte. The regenerated HF attacks the cathode surface, and this series of processes continues, further deteriorating cell performance. Thus, if the cathode materials are covered with stable metal fluoride layers such as AlF3, the cathode materials have less possibility to react with HF during electrochemical cycling, and the HF level in the electrolyte will be maintained, leading to further improvement of cycle life because there is no generation of water molecules. For the AlF3 coating, the outer surfaces of the active materials are surrounded by the AlF3; therefore, no further reactions occur. Unlike the Al2O3 coating layer, the AlF3 coating layer does not transform into other compounds but remains as AlF3 during extensive electrochemical cycling. Similar inorganic coating materials have also been applied to LiNiO2, NCA, and NCM cathodes to determine whether the coating materials can resolve the poor capacity retention of Ni-enriched NCM(A) cathodes. In the case of LiNiO2, La2O3162 and SiO2163 coatings have been shown to substantially increase the discharge capacity; however, the capacity retention was hardly improved as the initial capacity of the coated LiNiO2 decreased at nearly the same rate as that of pristine LiNiO2. Relatively little research exists on coating of Ni-enriched NCM(A) cathodes. MgO,164 TiO2,165 and V2O5136,138 have been used to coat LiNi0.85Co0.15O2, LiNi0.8Co0.2O2, and LiNi0.8Co0.1Mn0.1O2 cathodes. The MgO coating somewhat reduced the initial reversible capacity loss but delivered reduced discharge capacity compared to the pristine material. On the other hand, TiO2 and V2O5 coatings did improve the capacity retention but still experienced large irreversible capacity loss in the first charge/discharge cycle. Surface fluorination improved both capacity retention and initial columbic efficiency, to 86.8 from 82.7%.166 The surface coating of LiNiO2 of Ni-enriched NCM cathodes has shown varying degrees of success. In general, the coating lowered the charge transfer resistance by suppressing electrolyte breakdown and subsequent side reactions. Recently, demand for Ni-rich NCM(A) cathode materials is increasing because of their high capacity, which is required for extending the mileage of EVs. Because an excessive amount of lithium is necessary to synthesize highly crystalline Ni-rich layer compounds, unreacted lithium ingredients can remain on the

Li3PO4 + H 2O → LixHyPO4 (or POx Hy) + Li 2O

(4)

Water molecules are always present in the electrolyte as an impurity at a level less than 50 ppm, and this accelerates HF formation due to the decomposition of electrolytic salt, LiPF6. Once dissolution of active materials occurs, transition metal oxides are produced on the surfaces of active materials and are fluorinated by the HF to form water molecules. The produced water is again used for the propagation of HF, and the series of reactions ceaselessly continues. These undesirable reactions can be slowed with the help of surface modification, drastically improving cycling stability and rate performances. Attempting to lower the concentration of residual lithium compounds, Xiong et al.167 washed LiNi0.8Co0.1Mn0.1O2 with water and succeeded in lowering the pH of the active material. This action was effective for improvement in cycling performance and structural stability; however, the washed materials were less resistive to H2O and CO2 when stored in air. This emphasizes the necessity of heterogeneous layers on the surfaces of active materials in order to minimize contamination and prohibit the formation of undesired LiOH and Li2CO3 compounds. 211


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Figure 12. (a) TEM image and (a′) EDS compositional line scan from a fractured core−shell [(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2](OH)2 particle with the corresponding TEM image; (b) SEM images of Li[Ni0.8Co0.2]O2 and (c) core−shell Li[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 powders; (d) specific discharge capacity vs cycling number of (a) C/Li[Ni0.8Co0.2]O2 and (b) C/Li[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 cells in the voltage range of 3.0− 4.3 V; and (e) differential scanning calorimetry (DSC) traces of (a) Li[Ni0.8Co0.2]O2 and (b) Li[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 (adapted form ref 169).

4. SPATIAL PARTITIONING/GRADING OF COMPOSITIONS FOR NCM CATHODES The passivating effect of the surface coating helped to improve the cycle life of the cathodes. However, none of the cited works on coating of LiNiO2 or Ni-enriched NCM cathodes addressed the problem of potential thermal decomposition in the delithiated state. It is doubtful that merely protecting the surface of a cathode can fully resolve the inherent chemical and structural instability of Ni-rich NCM cathode materials. Therefore, there is a great need to develop the next generation of NCM cathodes that can utilize the benefits of each element (Ni, Co, and Mn) without compromising the energy density or safety. 4.1. Core/Shell Cathode. Another proposed solution to counter the trade-off between capacity and battery safety is to develop a core/shell NCM particle in which the particle core is enriched with Ni to provide a high discharge capacity (Figure 11, left).168 Additionally, the surface of the particle should be covered with a thin layer of NCM material enriched with Mn to provide chemical stability. This will significantly improve the structural stability during cycling by partially reducing the Ni3+ ions to more thermodynamically stable Ni2+ (Figure 11, center).170 On the basis of this idea, a fully functional microscale core/shell cathode was realized using a coprecipitation method (Figure 11, right).92 The core−shell NCM particles, whose diameter was 12 μm, were composed of a Li[Ni0.8Co0.1Mn0.1]O2 core encapsulated with a 1 μm thick Li[Ni0.5Mn0.5]O2 shell. The cathode was able to deliver 188 mAh g−1 when cycled between 3.0 and 4.3 V vs Li0/ Li+. The Mn-rich shell layer suppressed the potential HF attack from the electrolyte during long-term cycling and increased the onset temperature of the exothermic decomposition from 220 °C for the core material to 250 °C for the core/shell cathode.92 In another example of a core−shell cathode, Li-

[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 was also prepared via a coprecipitation route.92 Figure 12a shows a TEM image of the hydroxide precursor with the corresponding compositional line using energy-dispersive X-ray spectroscopy (EDS), clearly illustrating the partitioning of the 1 μm thick Mn-rich shell and Co-rich interior. A typical cross-sectional SEM image of the core−shell particle after heat treatment with LiOH is compared with that of the corresponding core material, Li[Ni0.8Co0.2]O2, without the shell in Figure 12b,c. A demarcation clearly distinguishing the core and shell after lithiation is shown in Figure 12c, whereas, rather expectedly, no such boundary was seen in the bulk material shown in Figure 12b. Extended cycling data of both core−shell Li[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 and Li[Ni0.8Co0.2]O2 in a full cell configuration in the cutoff range of 3.0−4.3 V, shown in Figure 12d, substantiate the superiority of the core−shell concept as Li[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 maintained 83% of the initial capacity after 800 cycles versus 64% for Li[Ni0.8Co0.2]O2 without the protective Mn-rich phase. Even more impressive was the thermal stability of the core−shell Li[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 cathode compared to Li[Ni0.8Co0.2]O2. As can be seen in Figure 12e, the onset temperature for the exothermic reaction of the delithiated cathode was increased by ∼60 °C, while the heat released by the reaction was substantially reduced by the presence of the Li[Ni0.5Mn0.5]O2 layer.92 Hence, the core−shell concept paved the way for developing a new series of cathode particles with compositional partitioning and demonstrated that the overall electrochemical performance of the core−shell cathode was superior to that of a NCM electrode with a uniform composition.171,172 4.2. Core/Compositionally Graded Shell Cathode. In further advancing the core−shell concept, a compositionally graded shell material was presented by our group. One motivation behind a compositional gradient in the shell material 212


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Figure 13. (a) SEM and electron-probe X-ray microanalysis (EPMA) results of the final lithiated oxide Li[Ni0.64Co0.18Mn0.18]O2. The gradual concentration changes of Ni, Mn, and Co in the interlayer are clearly evident. The Ni concentration decreases and the Co and Mn concentrations increase toward the surface. (b) Cycling performance at 1 C rate (75 mA corresponds to 190 mA g−1) of laminated-type lithium-ion batteries with an Al-pouch full cell (75 mAh) using mesocarbon microbead (MCMB) graphite as the anode and either Li[Ni0.8Co0.1Mn0.1]O2 or concentration gradient material as the cathode (upper cutoff voltage of 4.2 V), and (c) DSC traces showing heat flow from the reaction of the electrolyte with Li1−δ[Ni0.8Co0.1Mn0.1]O2, concentration gradient material Li1−δ[Ni0.64Co0.18Mn0.18]O2, and Li1−δ[Ni0.46Co0.23Mn0.31]O2 charged to 4.3 V (adapted from ref 170).

40 cycles because the thick shell layer was able to accommodate the interface strain, whereas a thin shell would be mechanically unstable when repeatedly stressed. To negate the structuredamaging effect of the volume expansion mismatch and to further increase the discharge capacity of the core−shell cathode, a concentration gradient was incorporated into the shell layer to achieve a smooth transition in composition from the core edge to the outer surface. The compositionally graded shell layer was formed on the Nirich NCM core by injecting a mixture of the Ni, Co, and Mn hydroxide solutions while the ratio of Ni, Co, and Mn was continuously changed. Figure 13a illustrates the Li[Ni0.8Co0.1Mn0.1]O2 cathode coated with the compositionally graded shell. In the shell, the concentrations of the molar fraction of Ni, Co, and Mn changed continuously from 8:1:1, finally reaching 4.6:2.3:3.1 at the surface. The shell thickness was increased to 3 μm to accommodate the steep concentration gradient of Ni from the core material to the particle surface. The cathode with the gradient shell at 4.4 V delivered an initial discharge capacity of 209 mAh g−1, which was only slightly lower than that of the core material Li[Ni0.8Co0.1Mn0.1]O2 cycled by itself, as can be seen from the inset in Figure 13b. More importantly, the cycle life of the Li[Ni0.8Co0.1Mn0.1]O2 electrode was greatly enhanced by the presence of the concentration gradient shell when tested using an Al-pouch full cell with graphite as the anode. As shown in Figure 13b, only 84% of the initial discharge capacity remained after 500 cycles for Li[Ni0.8Co0.1Mn0.1]O2, whereas the cathode coated with the

was to extract extra capacity from the shell material because the core material, Li[Ni0.8Co0.1Mn0.1]O2, by itself exhibited a discharge capacity of 200 mAh g−1 at 4.3 V compared to 188 mAh g−1 for the core−shell Li[(Ni0.8Co0.2)0.8(Ni0.5Mn0.5)0.2]O2 cathode. The reduced capacity of the core−shell cathode was largely due to the fraction of Li[Ni0.5Mn0.5]O2, whose discharge capacity is limited to 150 mAh g−1.173 Thus, minimizing the Mn content in the shell should increase the overall capacity of the core−shell cathode. The other reason for having a compositional gradient in the shell layer is the difference in volume expansion between the Ni-rich core and the Mn-rich shell during Li deintercalation. The Ni-rich NMC core typically experiences 9− 10% volume change during Li insertion and extraction, whereas Li[Ni0.5Mn0.5]O2 will undergo a much smaller 2−3% change in lattice cell volume.174,175 The disparity in volume change creates mechanical strain at the sharp interface between the core and the shell layer. During long-term cycling, the mechanical strain at the interface can lead to progressive microfracture of the shell layer from the core. The physical separation of the shell from the core is detrimental to Li+ diffusion and electronic conductivity and eventually causes degradation of the cathode capacity. To study the effect of shell thickness on cycling behavior, the thickness of the Li[Ni0.5Mn0.5]O2 outer layer was varied from 200 nm to 1.5 μm. The study indicated that as the shell thickness decreased the capacity retention of the core−shell cathode progressively deteriorated, most likely due to the interface strain developed from the volume change. In fact, the cathode with a 1.5 μm thick Li[Ni0.5Mn0.5]O2 shell retained 98% of the initial capacity after 213


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Figure 14. (a) TEM image and the corresponding electron diffraction pattern of an FCG Mn−F primary particle, illustrating the crystallographic alignment of the primary particle in the radial direction; (b) integrated atomic ratio of transition metals as a function of distance from the center of the particle for the lithiated FCG Mn−F material; (c) initial charge−discharge curves at 25 °C obtained from a 2032 coin-type half cell using lithium metal as the anode (a current density rate of 0.1 C corresponds to 22.2 mA g−1); (d) cycling performance of half cells between 2.7 and 4.5 V at 55 °C when applying a constant current rate of 0.5 C (114 mA g−1); and (e) cycling performance of a laminated-type Al-pouch cell (35 mAh) using MCMB graphite as the anode and FCG Mn−F as the cathode at a rate of 1 C, corresponding to 200 mAg−1 (upper cutoff voltage of 4.4 V) (adapted from ref 93).

Figure 11 graphically illustrates the concept of a NCM cathode with a FCG. Superimposed on Figure 14b are the relative compositions of the Ni, Co, and Mn probed by EPMA as a function of distance from the particle center. The compositions of Ni and Co linearly decreased from the center, while the Mn concentration smoothly increased such that the particle interior became increasingly enriched in Ni and Co and deficient in Mn. As summarized in Figure 14c−e, the discharge capacity, longterm cycling, thermal stability, and rate capability of the FCG were convincingly superior to those of the conventional cathode (CC) without a concentration gradient.93,179 Two of the unexpected features of the FCG cathode were the radially aligned primary particle morphology and crystallographic texture developed within each primary particle. Figure 14a shows a typical TEM image of a single FCG cathode whose thin section was prepared by fast ion bombardment. The FCG particle had a compact microstructure composed of elongated primary particles (∼100 nm in width and ∼2.5 μm in length) radially aligned, pointing toward the particle center. Furthermore, each elongated primary particle had strong crystallographic texture, with the c-axis of the unit cell aligned in the transverse direction to the primary particles so that the Li planes in the layered structure were radially aligned, fanning out from the particle center. The compact particle morphology and the strong [001] crystallographic texture facilitated Li+ transport and minimized contact with the electrolyte, additionally stabilizing the surface chemistry of the FCG electrode upon cycling. In an attempt to further optimize the energy density of the FCG cathode, we have engineered the processing conditions to introduce a concentration gradient with two slopes that span the entire particle.180,181 The two-sloped full-concentration gradient (TSFCG) was designed to maximize the Ni concentration near the core and the Mn concentration at the surface. The EPMA

gradient shell retained 96.5% of the initial capacity. The reduction of the Ni concentration while increasing the Mn concentration at the particle surface using the gradient shell limited the unwanted reaction with electrolyte during cycling to improve the cyclability of the electrode. Moreover, the concentration gradient within the shell likely relaxed the core/ shell interface strain that could have formed due to the sharp composition change and contributed to improving the long-term cycling of the electrode. Figure 13c confirms that, similar to the case of the core−shell electrode, the Mn-rich composition near the surface of the gradient shell significantly reduced the safety hazard of the Li[Ni0.8Co0.1Mn0.1]O2 cathode. The onset of the exothermic reaction resulting from reaction with the electrolyte was delayed by 90 °C compared to that of Li[Ni0.8Co0.1Mn0.1]O2 and very close to that of the Li[Ni0.46Co0.23Mn0.31]O2 cathode whose bulk composition was equal to that of the surface of the gradient shell.170 Various surface and core compositions have been tested to firmly establish that the composition gradation can be integrated into a NCM cathode using a coprecipitation synthesis route.176−178 The introduction of the compositional gradation was a unique approach; to the best of our knowledge, no other laboratories have produced a reliable working cathode with a concentration gradient. It is surmised that a graded composition is probably the only solution to attain all of the Li-battery design objectives for EVs: large discharge capacity, battery safety, and long-term cycling for the NCM cathode. Having successfully introduced a concentration gradient into a thin-shell layer, the next logical step was to extend the concentration gradient deeper into the particle core. 4.3. Full Concentration Gradient Cathode. The concentration gradient of Ni, Co, and Mn was stretched through the entire length of the particle to produce a so-called FCG cathode. 214


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Figure 15. (a) EPMA line scan of the integrated atomic ratio of transition metals as a function of the distance from the particle center to the surface for the lithiated TSFCG material; (b) TEM image (mosaic of several images) of an as-prepared TSFCG electrode, illustrating the rodshaped primary particle emanating from the equiaxed particles at the particle center and a bright-field TEM image showing a single elongated primary particle (dark particle at the center of the image); (c) cycling performance of a laminated pouch-type cell (25 mAh) using MCMB graphite as the anode and CC, NCA, and TSFCG as the cathode at a rate of 1 C, corresponding to 200 mA g−1 (upper cutoff voltage of 4.2 V); (d) comparison of rate capabilities of the TSFCG with CC and NCA cathodes (upper cutoff voltage of 4.3 V versus Li+ /Li); (e) DSC traces showing heat flow from the reaction of the electrolyte with CC Li1−δ [Ni0.65 Co0.13 Mn0.22 ]O2, TSFCG Li1−δ[Ni0.65Co 0.13Mn0.22]O2, and NCA Li1−δ[Ni0.8Co0.15Al0.05]O2 charged to 4.3 V (adapted from ref 180); (f) comparison of the CNT−Si/TSFCG against the currently developed LIBs; LCO: LiCoO2, NCM333: Li[Ni1/3Co1/3Mn1/3]O2, NCM523: Li[Ni0.5Co0.2Mn0.3]O2, NCM751114: Li[Ni0.75Co0.11Mn0.14]O2, and TSFCG: Li[Ni0.85Co0.05Mn0.15]O2. Schematic illustration of a Li-rechargeable battery system based on the CNT−Si anode and TSFCG cathode for vehicle application (inset) (adapted from ref 182).

compositional analysis of the particle center to the outer edge for the precursor hydroxide and lithiated TSFCG electrode, shown in Figure 15a, clearly shows the well-defined change in concentration gradient near the particle surface.181 The Ni concentration of the TSFCG electrode at the center was 72 mol %, which progressively decreased to 65 mol % at the junction and then to 60 mol % at the particle surface. Therefore, the average concentration of the interior was highly enriched in Ni, whereas the Mn concentration was maintained below ∼17 mol % in the core region and gradually increased to 28 mol % at the surface.

The TSFCG concept enabled the core region to be highly enriched in Ni to generate extra discharge capacity, with the Mnrich surface layer protecting the Ni-rich interior without producing any breaks in the rod-shaped morphology observed in the FCG cathode. The TEM images of the TSFCG cathode shown in Figure 15b attest to the fact that the unique morphology of FCG consisting of centrally aligned microscale primary particles that emanated from the particle center was indeed retained in the TSFCG electrode. Figure 15c,d,e summarize the superior electrochemical performance of the 215


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TSFCG in comparison with the CC with the same average composition of the TSFCG cathode. When cycled at 4.3 V, the TSFCG cathode delivered 200 mAh g−1, while the capacity of the conventional electrode was limited to 187 mAh g−1. Similarly, the cycle retention (increased from 71 to 88% after 1500 cycles at a 1 C rate), rate capability (increased from 110 to 150 mAh g−1 at a 5 C rate), and thermal stability (onset of the electrolyte reaction increased from 261 to 274 °C) were unequivocally improved by the TSFCG design, confirming the successful application of the TSFCG concept. The TSFCG cathode further pushed the boundaries that were set by poor cycle retention and thermal stability of the Ni-rich cathodes and perhaps enabled extraction of maximum possible capacity from the NCM cathode by forming a concentration gradient starting from LiNiO2 at the particle center. Development of the TSFCG is still in progress to maximize the Ni average content greater than 80%, for example, Li[Ni0.85Co0.05Mn0.1]O2 varying from Li[Ni0.89Co0.05Mn0.06]O2 to Li[Ni0.79Co0.06Mn0.15]O2.182 A full cell configured with a CNT−Si anode produced an energy density of 350 Wh kg−1 with excellent capacity retention for 500 cycles (Figure 15f), affirming the stability of TSFCG for delivery of high energy density suitable for EV application. This approach provides opportunities to optimize NCM battery performance by selectively distributing each transition metal within a single particle to meet the requirements of capacity, rate capability, and safety for next-generation EVs. It is also envisioned that the FCG concept can be applied to other areas to produce a wide range of multifunctional materials for application such as catalyst support, gas sensing, and drug delivery. For example, particles with a magnetic or ferromagnetic core covered with an antiferromagnetic layer can form a basis for synthesizing an exchange-spring-based permanent magnet.

Figure 16. Radar chart that shows the performance comparison of NCA (Ni 80%) with the blue line and NCM (Ni 80%) with the red line.

content in the compound. This means that process costs and higher energy density can be obtained for Ni-rich NCM to optimize the price per Wh. Other than the capacity, operation voltage, and price concerns, NCA materials exhibit better capacity retention than that of Ni-rich NCM.157 In particular, the cycling performance of Ni-rich NCM is lacking, presumably due to dissolution of Mn from the parent oxide, which is fatal for cycling at elevated temperature. Meanwhile, the presence of Mn provides significant thermal stability that elevates the onset exothermic reaction to 220 °C,107 relative to the 180 °C of Al in NCA.98,156 Another aspect that needs to be considered is that, despite their chemical similarity, NCM and NCA cathodes are typically produced using different synthetic routes. The introduction of Al into the NC structure is usually achieved via heat treatment, whereas Mn can be more easily added via coprecipitation. This may introduce a homogeneity advantage for the NCM system. Figure 16 summarizes representative properties of both NCA and Ni-rich NCM cathode materials. Synergetic effects of NCA and NCM were demonstrated in Aldoped TSFCG, Li[Ni0.84Co0.06Mn0.09Al0.01]O2.183 This material was able to deliver a high discharge capacity of approximately 240 mAh g−1 in the voltage range of 2.7−4.5 V (0.5 C rate) at the first discharge and retained a capacity as high as 202 mAh g−1 (89%) after 100 cycles. Nevertheless, the stabilizing activities of Al and Mn are not identical (Al3+ is lighter, facilitates Li+ diffusion, and replaces mostly Ni3+; Mn4+ introduces more Ni2+ in the structure, allows slightly higher capacities, but leads to a higher degree of cation mixing), and layered oxides could benefit from their simultaneous presence. In addition, it was demonstrated that long-term cycling behavior of a Ni-rich NCM cathode can be greatly improved by Al-doping as an Al-doped FCG cathode with an average concentration of Li[Ni0.60Co0.12Mn0.27Al0.01]O2 remained stable even after 3000 cycles (representing a single daily charge/discharge cycle during typical service life of an EV (ca. 10 years)) and maintained 83% of the initial capacity.184 On the other hand, an undoped FCG cathode with identical composition experienced a substantial capacity loss after only 500 cycles. SEM analysis of the cycled cathodes confirmed the mechanical stability of the Al-doped FCG cathode during cycling by revealing that the Al-doped FCG cathode particles remained intact whereas the particles from the undoped FCG cathode were completely pulverized after 3000 cycles (Figure 17). The superior long-term cyclability of the Al-doped FCG cathode

5. COMPARISON OF NI-RICH NCA AND NCM Both NCA and Ni-rich NCM have 80% Ni in the transition metal layer. As is well-known, it is rather difficult to synthesize stoichiometric LiNiO2 with a Li/Ni ratio of 1:1. The synthesis often results in the formation of lithium-deficient Li1−xNi1+xO2 with part of the Ni present as Ni2+ and cation mixing between Ni and Li as it is difficult to maintain all Ni as Ni3+ at high synthesis temperatures of >700 °C.7−10 Because of the large amount of Ni, similar to LiNiO2, additional lithium should be added to produce NCA and Ni-rich NCM to minimize cation mixing. However, unreacted additional lithium remains on the surfaces of both active materials as double-layer lithium residues such as a LiOH inner layer and Li2CO3 outer layer. The presence of such residual lithium compounds is not desirable because oxidative decomposition of those compounds facilitates gas evolution at higher voltage. Also, the presence of Li2CO3 is harmful for long-term cyclability, as mentioned in section 2.1.3. For these reasons, elimination of such lithium residues by washing and/or coating is required for high-performance batteries. Unfortunately, ionic exchange between Li+ and H+ occurs violently when NCA and Ni-rich NCM are washed in water. Care is thus necessary to minimize ionic exchange during removal of residual lithium compounds. Ni-rich O3-type layer cathode materials appear to be the ultimate step for achieving optimal energy density because of high operation voltage and high capacity delivery of approximately 200 mAh g−1. In comparison of the two prototype materials, NCA and LiNi0.8Co0.1Mn0.1O2 (Ni-rich NCM) (Figure 16), there is no difference in Ni content. One merit of the Ni-rich NCM is its cost-effectiveness due to reduction of Co 216


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necessary both to achieve a better understanding of the fundamental mechanisms behind the beneficial actions of these methods and to explore possible synergistic effects between them.

We show how future automotive targets can be achieved through fine control of the structural and microstructural properties. Surface stabilization is mandatory to diminish the reactivity and reduce the amount of highly reactive Ni4+ and/or impede the reactivity toward the electrolyte. The simultaneous adoption of a compositional gradient, with a Mn-rich shell region, and a surface coating appear at the moment as the most promising technical solution. Further improvements could include further refining of the gradient profile to further optimize the trade-off between energy density and cycle stability. Coatings deposited with more sophisticated deposition techniques, like atomic layer deposition, could also prove beneficial due to superior uniformity and potential for much better thickness control. Promising results on this topic have been recently reported for both NCA and NCM 811 cathodes. Nevertheless, the impacts of these kinds of deposition techniques on the final material cost remain to be fully evaluated.185 Morphology control has also been actively investigated to achieve faster Li+ intercalation kinetics, as demonstrated for the compositionally graded NCM materials with an acicular primary particle microstructure. Morphology control is also expected to affect cycle lifetime, for example, at prolonged cycling counts where the detachment of primary particles becomes one of the most severe causes of capacity loss. We believe that this is one of the aspects with the largest room for improvement. Further refinement of the primary and secondary particle morphologies with the additional possibility to introduce morphological gradients coupled with the chemical gradient could produce large material improvements. A noteworthy difference between NCA and NCM systems exists in their modification strategies, the compositional gradient in particular. Early experiments showed that, in contrast to the relative ease with which concentration gradients are formed in NCM materials and preserved upon cycling, concentration gradients tend to disappear in NCA upon high-temperature calcination. Although this finding requires an in-depth investigation, it suggests that, in direct comparison between NCA and NCM cathode materials as possible candidates for future automotive electrification, the ability to withstand surface modification processes could be one of the most decisive factors.

Figure 17. (a) Cycle performance of Al-doped and undoped FCG Li[Ni0.60Co0.12Mn0.27Al0.01]O2 cathode at 1.0 C in a full cell vs graphite in the voltage range of 3.0−4.2 V; SEM images of (b,c) the undoped FCG cathode, (d,e) the Al-doped FCG cathode after 3000 cycles (reproduced from ref 184).

was further notable for the fact that the Al FCG cathode was cycled at a 100% DOD for 3000 cycles to fully utilize its available capacity for a maximum energy density because the DOD of a NCA cathode for long-term cycling is typically limited to 60%, which adds dead weight to the battery and increases the material cost.62,86 Further work on optimization of the composition and the choice of electrolyte with adequate additives is necessary to increase the energy density and extend the cyclability through the combination of NCA and NCM.

6. CONCLUSIONS AND PERSPECTIVE In comparison with other types of novel cathode materials, NCAand NCM-based cathode families have an obvious advantage in terms of technological maturity. However, there is far less room for further optimization of these cathodes. In other words, using NCA and NCM cathodes to fill the gap between the present generation of automotive batteries and the 2025 target requires fine-tuning of the available variables of nickel content, upper cutoff voltage, packing density, and loading level. In parallel, the development of efficient modification techniques must offer even better mitigation of the degradation processes that unavoidably will become harsher as the cathode energy density is increased. With regard to the cycling and safety-related performance of layered cathode materials, it is evident that there is no chance for pristine materials to be implemented in future cells without sacrificing a large amount of energy density. Apart from possible co-doping, the results presented in this Review clearly indicate that no Ni-rich cathode material will be able to meet the energy density, power, lifetime, and safety targets unless one or more stabilization techniques are adopted. Further investigations are

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (Y.-K.S.). *E-mail: [email protected] (S.-J.K.). ORCID

Yang-Kook Sun: 0000-0002-0117-0170 Author Contributions #

S.-T.M. and F.M. contributed equally to this work.

Notes

The authors declare no competing financial interest. Biographies Seung-Taek Myung is a Professor of Nano Engineering at Sejong University, South Korea. He received his Ph.D. degree in Chemical 217


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(7) Cabana, J.; Monconduit, L.; Larcher, D.; Palacin, R. M. Beyond Intercalation-Based Li-Ion Batteries: The State of the Art and Challenges of Electrode Materials Reacting Through Conversion Reactions. Adv. Mater. 2010, 22, E170−E192. (8) Andre, D. Future Generations of Cathode Materials: An Automotive Industry Perspective. J. Mater. Chem. A 2015, 3, 6709− 6732. (9) Lee, J. K.; Oh, C.; Kim, N.; Hwang, J.-Y.; Sun, Y.-K. Rational Design of Silicon-based Composites for High-Energy Storage Devices. J. Mater. Chem. A 2016, 4, 5366−5384. (10) Kim, H.; Lee, E.-J.; Sun, Y.-K. Recent Advances in the Si-based Nanocomposite Materials as High Capacity Anode Materials for Lithium Ion Batteries. Mater. Today 2014, 17, 285−297. (11) McDowell, M. T.; Lee, S. W.; Nix, W. D.; Cui, Y. 25th Anniversary Article: Understanding the Lithiation of Silicon and Other Alloying Anodes for Lithium-Ion Batteries. Adv. Mater. 2013, 25, 4966−4985. (12) Freedom CAR Electrical Energy Storage System Abuse Test Manual for Electric and Hybrid Electric Vehicle Applications. Sandia Report SAD2005-3123; Sandia National Laboratories: Albuquerque, NM, 2006. (13) Chen, Z.; Dahn, J. R. Effect of a ZrO2 Coating on the Structure and Electrochemistry of LixCoO2 When Cycled to 4.5 V. Electrochem. Solid-State Lett. 2002, 5, A213−A216. (14) Chen, Z.; Dahn, J. R. Studies of LiCoO2 Coated with Metal Oxides. Electrochem. Solid-State Lett. 2003, 6, A221−A224. (15) Chen, Z.; Dahn, J. R. Improving the Capacity Retention of LiCoO2 Cycled to 4.5 V by Heat-Treatment. Electrochem. Solid-State Lett. 2004, 7, A11−A14. (16) Miyashiro, M.; Kobayashi, Y.; Seki, S.; Mita, Y.; Usami, A.; Nakayama, M.; Wakihara, M. Fabrication of All-Solid-State Lithium Polymer Secondary Batteries Using Al2O3-Coated LiCoO2. Chem. Mater. 2005, 17, 5603−5605. (17) Miyashiro, H.; Yamanaka, A.; Tabuchi, M.; Seki, S.; Nakayama, M.; Ohno, Y.; Kobayashi, Y.; Mita, Y.; Usami, A.; Wakihara, M. Improvement of Degradation at Elevated Temperature and at High State-of-Charge Storage by ZrO2 Coating on LiCoO2. J. Electrochem. Soc. 2006, 153, A348−A353. (18) MacNeil, D. D.; Dahn, J. R. The Reactions of Li0.5CoO2 with Nonaqueous Solvents at Elevated Temperatures. J. Electrochem. Soc. 2002, 149, A912−A919. (19) http://www.telegraph.co.uk/news/1526424/Exploding-laptopsprompt-Dell-battery-recall.html (Aug 2006). (20) Goodenough, J. B.; Wickham, D. G.; Croft, W. J. J. Phys. Chem. Solids 1958, 5, 107−116. (21) Morales, J.; Perez-vicente, C.; Tirado, J. L. Cation Distribution and Chemical Deintercalation of Li1−xNi1+xO2. Mater. Res. Bull. 1990, 25, 623−630. (22) Hirano, A.; Kanno, R.; Kawamoto, Y.; Takeda, T.; Yamaura, K.; Takano, M.; Ohyama, K.; Ohashi, M.; Yamaguchi, Y. Relationship between Non-stoichiometry and Physical Properties in LiNiO2. Solid State Ionics 1995, 78, 123−131. (23) Arai, H.; Okada, S.; Sakurai, Y.; Yamaki, J. Reversibility of LiNiO2 cathode. Solid State Ionics 1997, 95, 275−282. (24) Thomas, M. G. R. S.; David, W. I. F.; Goodenough, J. B.; Groves, P. Synthesis and Structural Characterization of the Normal Spinel Li[Ni2]O4. Mater. Res. Bull. 1985, 20, 1137−1146. (25) Dahn, J. R.; von Sacken, U.; Michal, C. A. Structure and Electrochemistry of Li1±yNiO2 and a New Li2NiO2 Phase with the Ni (OH)2 Structure. Solid State Ionics 1990, 44, 87−97. (26) Dahn, J. R.; von Sacken, U.; Juzkow, M. W.; Al-Janaby, H. Rechargeable LiNiO2/Carbon Cells. J. Electrochem. Soc. 1991, 138, 2207−2211. (27) Ohzuku, T.; Komori, H.; Nagayama, M.; Sawai, K.; Hirai, T. Electrochemical Characteristics of LiNiO2. Chem. Express 1991, 6, 161. (28) Ohzuku, T.; Ueda, A.; Nagayama, M. Electrochemistry and Structural Chemistry of LiNiO2 (R3m ̅ ) for 4 V Secondary Lithium Cells. J. Electrochem. Soc. 1993, 140, 1862−1870.

Engineering from Iwate University, Japan, in 2003. His research interests embrace development of electroactive materials and corrosion of current collectors of rechargeable lithium and sodium batteries. Filippo Maglia received his Ph.D. from the University of Pavia and was assistant professor in the Chemistry Department (2005−2012). He was visiting scientist at UC Davis and a DAAD fellow at the Technische Universität München. He is currently working on battery materials in the Research Battery Technology department at BMW. Kang-Joon Park is a Ph.D. student with Prof. Yang-Kook Sun in energy engineering at Hanyang University. He received a B.S. degree from Hanyang University in 2013. His current research is focused on synthesis of functional cathode materials for lithium-ion batteries. Chong Seung Yoon received his Ph.D. degree from Massachusetts Institute of Technology. He joined the Department of Materials Science and Engineering at Hanyang University in 2003 and has worked on cathodes materials for LIBs since then. His research interests include microstructure characterization using transmission electron microscopy, energy-storage materials, and magnetic materials. Peter Lamp received his Ph.D. from the Max Planck Institute for Physics, Munich. He was group leader at the Department of Energy Conversion and Storage of the Bayerischen-Zentrum für AngewandteEnergieforschung and project leader for fuel-cell systems at Webasto Thermo Systems International GmbH. At BMW since 2001, he is now leader of the Research Battery Technology department. Sung-Jin Kim received his Ph.D. in Chemistry from the Technische Universität München. He was a postdoc at Ames Laboratory and then joined LG Chem as a developer for Li-ion cells (2009−2012). At BMW AG, he is currently working on material and cell projects in the Research Battery Technology department. Yang-Kook Sun received his Ph.D. degree from the Seoul National University, Korea. He was group leader at the Samsung Advanced Institute of Technology and contributed to commercialization of the lithium polymer battery. He has worked at Hanyang University in Korea as a professor since 2000. His research interests are the synthesis of new electrode materials for lithium-ion batteries, Na-ion batteries, Li−S batteries, and Li−air batteries.

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ACKNOWLEDGMENTS This work was mainly supported by the Global Frontier R&D Program (2013M3A6B1078875) at the Center for Hybrid Interface Materials (HIM) funded by the Ministry of Science, Information & Communication Technology (ICT) and the National Research Foundation of Korea (NRF) grant funded by the Korea government (MEST) (No. 2014R1A2A1A13050479).

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