Cathode for Li Ion Batteries - ACS Publications - American Chemical

Research Article www.acsami.org

Theoretical Prediction of Surface Stability and Morphology of LiNiO2 Cathode for Li Ion Batteries Eunseog Cho,* Seung-Woo Seo, and Kyoungmin Min Platform Technology Lab, Samsung Advanced Institute of Technology, 130 Samsung-ro, Suwon, Gyeonggi-do 16678, Republic of Korea S Supporting Information *

ABSTRACT: Ni-rich layered oxides are considered to be a promising cathode material with high capacity, and their surface structure should be extensively explored to understand the complex associated phenomena. We investigated the surface stability and morphology of LiNiO2 as a representative of these materials by using density functional theory calculations. The results reveal that the Li-exposed surfaces have lower energies than the oxygen surfaces, irrespective of the facets, and the Ni-exposed ones are the least stable. The equilibrium morphology can vary from truncated trigonal bipyramid to truncated egg shape, according to the chemical potential, whose range is confined by the phase diagram. Moreover, the electrochemical window of stable facets is found to strongly depend on the surface elements rather than the facet directions. Contrary to the stable Li surfaces, oxygen exposure on the surface considerably lowers the Fermi level to the level of electrolyte, thereby accelerating oxidative decomposition of the electrolyte on the cathode surface. KEYWORDS: LiNiO2, Ni-rich layered oxide, surface energy, equilibrium morphology, absolute band alignment, density functional theory calculations

I. INTRODUCTION Ni-rich layered oxides are considered to be potential cathode materials for lithium ion batteries in electric vehicles and energy storage devices, because Ni is believed to be important for improving the capacity.1 For example, it is well known that a higher Ni content in nickel−cobalt−manganese oxide (NCM) increases the capacity; Co contributes to higher rate capability, and Mn enhances the thermal stability.2−5 Thus, oxides with >80% Ni have been widely studied to achieve high storage capacities as large as 200 mAh/g.1,6−10 However, the increased Ni content in the oxide also induces several types of degradation behaviors as the cycles continue, preventing the material’s commercial application despite its high capacity.11,12 The degradation of Ni-rich structures has been attributed to various factors: lattice instability, cation disordering, phase transformation, oxygen release, and microcrack propagation.2,11,13−16 Extensive studies have been carried out to discover the origins11,13−19 and suggest strategies20−29 to alleviate such degradation phenomena. For example, our previous studies explained that the NCM with high Ni content could suffer from structural instability17 and all types of degradation behaviors can be accelerated in concert with the delithiation process.18 As a promising strategy, several types of dopants that are introduced in the bulk region have been proposed.20−27 Among them, Al20,21 and Mg22,23 have proved to be effective © XXXX American Chemical Society

for increasing the capacity retention via enhancing structural and thermal stability. In addition, such degradation behaviors depend strongly on the surface features, because the surface itself is more vulnerable to environmental changes compared to the inner region during Li intake/release. In this regard, surface modifications to the cathode such as the surface coating method have also been suggested as an effective way for suppressing degradation that originates from the surface. The surface coating approach aims to provide a physical barrier at the surface, because the cathode surface is likely to be degraded by chemical or electrochemical reactions with the electrolytes.28−30 Furthermore, such surface reactions can induce electrolyte decomposition and subsequent gas generation, which causes the battery pack to swell and permanently deteriorates the performance.31 The metal oxides and metal phosphates have been widely explored as coating materials, and recently, nanoscale layers such as the cubic phase have been reported as another way to improve the stability of layered oxides.30 Meanwhile, as the Ni content in the cathode increases, the amount of Li residues such as LiOH and Li2CO3 also increases on the cathode surface after synthesis, and these are the source of gas evolution by their decomposition reactions Received: June 14, 2017 Accepted: September 12, 2017 Published: September 12, 2017 A

DOI: 10.1021/acsami.7b08563 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

ACS Applied Materials & Interfaces

(GGA) implemented by Perdew, Burke, and Ernzerhof for the exchange correlation energy functional with the spin polarization method.39 The on-site Coulomb interaction (GGA + U approach),40 which has been proved to accurately predict the structural and redox properties,41−43 was used to describe the strongly correlated d electrons of the Ni (6.5 eV)18,44 element. Ferromagnetic ordering was performed for all calculations. For the magnetic ordering of Ni, recent muon-spin relaxation (μSR) data45 suggest that the energy values from the ferromagnetic and antiferromagnetic states are close to each other, and previous DFT calculation also predicted that the ferromagnetic ground state is just 3−5 meV per unit cell lower in energy than that from the antiferromagnetic state.42 The energy cutoff was selected to be 500 eV, and the atomic positions for bulk and slab structures were fully relaxed until the ionic force on each atom was below 0.02 eV/Å. The energy convergence criterion was set to be 10−5 eV. The cell parameters were optimized for the bulk LiNiO2 structure (R3̅m; a = 2.896 Å, c = 14.114 Å), and then used to carry out the surface relaxation. Different supercell structures were adopted for each facet. The slab thicknesses as well as the kpoint samplings over the Brillouin zone were selected to guarantee the convergence of the total energy and force. Tasker showed that based on electrostatic considerations, when the slab has a net dipole moment perpendicular to the surface, the energy diverges as the thickness of the simulated slab increases.46 This is indeed the case for the asymmetric slab model, because top and bottom surfaces have different atomic elements, and thus, the net dipole moment is different from zero. For example, the construction of the polar facets by simply cleaving the bulk crystal generates the asymmetric slab structure, satisfying the stoichiometry of the bulk structure. However, this slab structure produces a slope in the electrostatic potential along the direction of the slab thickness because of the noncompensated dipole moments, thus the vacuum level cannot be correctly defined. To prevent this undesirable dipole occurrence, we adopted the symmetric structure, that is, the atomic configuration on the top and the bottom surfaces of the slab are identical. The comparison between the two different slab structures and their electrostatic potentials is shown in the Supporting Information (SI), Figure S1. The detailed information on supercell size, surface area, slab thickness, and k-points sampling for the nonpolar and polar facets we adopted for this study is also shown in Table S1. For some structures with partially exposed Li on the (003) facet, the electrostatic potential at the top and bottom sides could be slightly different although the symmetric structures were adopted. Thus, a dipole correction in the z-direction (perpendicular to the surface) was additionally used to guarantee accuracy in the Fermi energy alignments. A vacuum region of more than 15 Å was also selected along the zdirection to avoid spurious effects stemming from the periodic boundary condition. II.II. Surface Energies and Facets. The first consideration in the modeling is which surface facets need to be chosen for the stability and morphology predictions. This selection is based on the powder X-ray diffraction (XRD) data and the Bravais−Friedel−Donnay−Harker (BFDH) method. The powder XRD data for the bulk (R3̅m) LiNiO2 structure contains strong intensities for Miller indices (003) and (104), and several minor peaks such as (012) and (101).35,47 The BFDH method gives a rough estimate of the facets that are likely to be important for the crystal structure based on two

mainly with electrolytes. Thus, optimal coating materials have been newly devised that can effectively remove the Li residues in addition to the role of a physical barrier preventing unwanted contact with the electrolytes.28,29 Although various approaches to modify the surfaces have been demonstrated on Ni-rich layered oxides, the atomistic level of insight into the surface structure is still lacking. On the theoretical side, a few studies have been carried out for the surface structures of LiNiO2 to obtain atomic-scale surface features of Ni-rich layered oxides.32−34 However, most of these studies only focus on one or two specific facets of the LiNiO2 structure,32,33 or do not properly reflect the characteristics of the polar facets, whose specific atoms can be exclusively exposed on the surface and whose stability depends on the variations of the chemical potentials.34 Hence, the detailed surface information of Ni-rich layered oxides remains unclear, even though a precise understanding of the surface is needed for controlling the various surface features. This study is designed to theoretically predict the structural and electrochemical properties of the LiNiO2 surface. Current research has focused on increasing the Ni content in electrode materials, even up to >90% of the total amount of transition metals.28 Hence, LiNiO2 can be a proper model for the general features of recently developed Ni-rich oxides. Next, the preparation of the pure LiNiO2 phase with good capacity retention is difficult, owing to the large amounts of intrinsic defects, such as Li−Ni mixing and vacant oxygen. Very recently, Xu et al. used a solidstate method to synthesize stoichiometric LiNiO2 with excellent capacity retention.35,36 The development of such new synthetic methods can further facilitate the adoption of LiNiO2 in batteries, considering that it has the maximum capacity among Ni-rich layered oxides as well as better cost efficiency compared to that of the conventional cathode material of LiCoO2. In this article, we examine the surface energy, equilibrium morphology, and absolute band alignment of LiNiO2 surfaces within the density functional theory (DFT) framework. For the surface structure of Ni-rich layered oxides, a rough experimental consensus is that as the cycles go on, the original rhombohedral phase (R3̅m) partially transforms to a spinel-like phase (Fd3̅m) 3,6,12,31 or NiO-like rocksalt (Fm3m Although the surface ̅ ) one. phases vary somewhat depending on the degradation status, the original rhombohedral phase (R3̅m) was chosen for this study because it is the fundamental and dominant phase. Contrary to previous studies, both polar and nonpolar facets are investigated here. Environmental effects closely linked to the surface properties, such as temperature and oxygen flow, are included as a form of chemical potential. From the constructed Li−Ni−O phase diagram, the stable region of the LiNiO2 phase is thermodynamically defined by the value of chemical potential. Thus, we could explain the surface energy and morphological change within the definite range of chemical potential obtained from the phase diagram. Moreover, the energy band alignment of each facet as well as the exposed atoms is calculated, and the corresponding electrochemical stability is also revealed by comparing with the energy level of the electrolyte.

II. COMPUTATIONAL DETAILS II.I. First-Principles Calculations. The first-principles calculations were performed by the pseudopotential plane wave method, using the Vienna ab initio simulation package.37,38 We adopted the generalized gradient approximation B


Research Article


change of Δμ(O) at finite temperature T (K) and oxygen partial pressure p (atm) is generally represented as

assumptions: the center-to-facet distance is inversely proportional to the lattice plane spacing (BF rule), and higher order planes are preferentially grown over the lower order ones (DH rule).48−50 Using this method, we estimate that the surface of the LiNiO2 structure is composed of (003), (101), and (104) facets (Figure S2 in SI). However, the BFDH method only reflects the crystal symmetry and the lattice parameters and it neglects all of the chemical nature and the related energetics. Therefore, the stronger the chemical bonding effects in the crystal, the less accurate results that this method produces. On the basis of both the XRD and BFDH information, we chose the (104), (110), and (100) facets for the nonpolar surfaces and the (003), (012), and (101) facets for the polar surfaces to predict the surface properties. For the nonpolar facets, the exposed atoms on the surface contain all three elements (Li, Ni, and O) satisfying the bulk stoichiometry of LiNiO2. In contrast, the polar surface has fewer elements at the outmost layer; for the (003) facet, Li, Ni, or O can be exclusively exposed on the surface. Recently, it has been proposed that for the (111) facet of LiMn2O4, the surface Mn can be exchanged with the subsurface Li atoms to form a partial inverse spinel arrangement.51−53 Thus, we also checked the possibility that the exposed Ni(Li) on the surface can be exchanged with the subsurface Li(Ni). Although the Li−Ni exchange in the inner region has been widely reported for Ni-rich layered oxides,11,13,18 the Li−Ni exchange on the surface has not been confirmed. The surface energy (γ) of each facet is defined as the energy difference between the slab and bulk structures. It is calculated with the following equation. γ=

Gslab − ∑i n(i)μ(i) 2A

∼

Δμ(O) =

(2)

Here, μ(O2; T, p), the chemical potential of molecular oxygen, is expressed as μ(O2 ; T , p) = μ(O2 , 0 K) + ΔG°(O2 , T ) + RT ln p(O2 )

(3)

where ΔG°(O2, T) is the change in O2 free energy, and RT ln p(O2) is the contribution of O2 partial pressure (e.g., p(O2) = 0.2 atm in atmospheric conditions). For the GGA-based calculation, the binding energy of O2 is well known to be overestimated, which is associated with adding electrons to the oxygen p orbital to form anions from the O2 molecule, and such an error affects the energies of oxide materials.54,55 Wang et al. calculated this constant error as 1.36 eV per O2 molecule in the oxidation reaction within the GGA + U framework, by fitting the formation enthalpy of transition metal oxides.56 Their correction factor is used here to calculate μ(O2, 0 K) by shifting the oxygen energy obtained from the GGA + U calculation. The ΔG°(O2, T) of gaseous oxygen is affected by the change of both enthalpy and entropy according to the temperature, and its value is obtained from the JANAF thermochemical table.58 Thus, the overall equation of μ(O2, T, p) can be rewritten as μ(O2 ; T , p) = μ(O2 , GGA‐fit) + ΔH °(O2 , T ; JANAF) − T ΔS°(O2 , T ; JANAF) + RT ln p(O2 ) (4)

Eslab − ∑i n(i)μ(i) 2A

1 [μ(O2 ; T , p) − μ(O2 , 0 K)] 2

Δμ(O) decreases almost linearly with increasing temperature, and this trend at p(O2) = 0.2 atm is plotted in Figure S3. Similarly, Δμ(Li) is represented as the difference between the chemical potential of Li and that of pure Li metal: Δμ(Li) = μ(Li) − μ(Li; metal). Meanwhile, the Ni energy in the Nicontaining oxides (e.g., LiNiO2, LiNi2O4, and NiO) should also be corrected by different amounts according to the chosen U value. The reason for this is that the formation energy of transition metal oxides should be calculated using the energies obtained by the different methods. For example, the formation energy of LiNiO2 is defined by the energy difference between LiNiO2 and the pure elements (Ni and Li metals and O2 molecules). However, here the energy of LiNiO2 is calculated using the GGA + U framework, whereas that of pure Ni metal is calculated using just the GGA method. Hence, their energy difference gives an incorrect result. Jain et al. suggested a way to correct the energy, by fitting the enthalpy difference between the calculation and the experiment as a function of the fraction of transition metal element.57 Using Jain’s approach, the correction value of Ni is obtained as 2.912 eV/Ni with the U value of 6.5 eV. The fitting graph is shown in Figure S4.

(1)

Gslab, the free energy of the slab structure, can be approximated by Eslab (slab structure energy) assuming that the volume change is trivial. n(i) and μ(i) are the number of atoms and the chemical potential of species i, respectively. The surface area (A) is multiplied by 2, because the slab structure has two surfaces (top and bottom) in the model. Here, μ(i) depends on the environmental conditions, and accordingly, the surface energy is not constant under environmental changes such as the variance of temperature. For the nonpolar slab structures, ∑in(i)μ(i) is easily obtained from the energy of the bulk LiNiO2 structure, so that individual values of μ(i) are not required for determining the surface energy. However, for the general polar facets, the element composition of the slab structure no longer follows the stoichiometry of bulk LiNiO2. For example, in the O-exposed surface, the number of O atoms is not twice that of the Li or Ni atoms in the slab structure of LiNiO2. Because the surface and bulk structures are in equilibrium, the relation μ(Li) + μ(Ni) + 2μ(O) = ELiNiO2 always holds by the thermodynamic stability condition of the LiNiO2 structure. Thus, the surface energies of the polar facets depend on just two independent chemical potentials (μ(Li) and μ(O)) due to the above constraint. II.III. Phase Diagram Construction. The phase diagram of Li−Ni−O was constructed with the formation energies obtained from DFT calculations. To acquire reliable formation energies, the energies of both O and Ni for LiNiO2 should be corrected, because the GGA-based calculations have been reported to generate poor values for these elements.54−57 The

III. RESULTS Figure 1 shows the phase diagram of Li−Ni−O as a function of both Δμ(O) and Δμ(Li). To construct the diagram, we calculated the formation energies of 10 structures including five Ni-containing oxides. The space groups and the total energies are tabulated in Table S2 (SI). As explained in Section II.III, Δμ(O) can be physically represented as the change of both temperature and oxygen partial pressure, that is, a decrease in Δμ(O) corresponds to an increase in T and/or a decrease in C


Research Article


the neighboring oxide phases. Moreover, the phase of LiNiO2 disappears when Δμ(O) < −1.14 eV irrespective of Δμ(Li), which corresponds to 980 K at p(O2) = 0.2 atm, and thus the stable rocksalt phase of NiO is expected to form at over 980 K. This prediction is supported by the experimental finding that the hexagonal (layered) LiNiO2 could be easily synthesized at 873 K, but the cubic (rocksalt) phase was obtained at 1053 K.47,48 Meanwhile, the LiNiO2 phase does not form an equilibrium phase (no tie line appears) with O2 gas at 0 K, because both the spinel LiNi2O4 and monoclinic Li2NiO3 phases block its formation. This is clearly represented in the ternary phase diagram at 0 K shown in Figure S5. Because there is no tie line between LiNiO2 and O2, the surface of LiNiO2 cannot also be contacted with gaseous O2 in equilibrium conditions at 0 K. However, the surface of LiNiO2 can be transformed to Li2NiO3 or LiNi2O4 because the LiNiO2 phase has a tie line with both Li2NiO3 and LiNi2O4. Hence, a plausible equilibrium structure at 0 K is one where the surface Li2NiO3 or LiNi2O4 phases are placed on the LiNiO2 interior (bulk) phase if enough oxygen gas is provided. However, after the spinel phase disappears over around 300 K, the LiNiO2 phase can eventually form an equilibrium state with O2 gas. (Note that the Li2NiO3 phase disappears after the tie line between LiNiO2 and O2 is formed.) This phase diagram implies that the surface structure of the LiNiO2 phase can exist thermodynamically in conjunction with O2 over around 300 K, although this diagram was constructed using the information of the bulk structures instead of the surface phases. Figure 2 shows the relaxed geometries of the nonpolar facets ((104), (110), and (100) surfaces) together with several representative structures of the polar facets ((003), (102), and (101) surfaces). All polar facets that we considered here are also shown in Figures S7−S9. For the nonpolar facets, each layer satisfies the stoichiometry of bulk LiNiO2 structure, so the structures of the surface layers (e.g., interlayer spacing and relative atomic positions) are similar to those in the interior region. For the (104) and (110) facets, the interlayer spacing on the surface only expands by 4% and diminishes by 8% compared to that of the interior region, respectively. Moreover,

Figure 1. Phase diagram of Li−Ni−O as a function of changes in O and Li chemical potentials. The red dotted line corresponds to Δμ(O) at 300 K at p(O2) = 0.2 atm.

p(O2). Under a synthetic environment of LiNiO2, the temperature and oxygen partial pressure can play a role in controlling parameters, and their effects can be expressed as the oxygen chemical potential. In other words, when we increase the temperature during synthesis, the value of Δμ(O) decreases, and when we raise the partial pressure of oxygen gas through the oxygen blowing process, the value of Δμ(O) increases. For convenience, we assume that the synthesis is carried out in atmospheric conditions with p(O2) fixed at 0.2 atm, then Δμ(O) only reflects the temperature. In the phase diagram, the LiNiO2 phase is found to exist in an extremely narrow area (−3.44 eV < Δμ(Li) < −2.53 eV, −1.14 eV < Δμ(O) < −0.16 eV) enclosed by other Ni-containing oxides such as LiNi2O4, NiO, and Li2NiO3. The Ni3O4 spinel structure is not stable in the diagram. Considering that the stable area of LiCoO2 in the phase diagram (−3.63 eV < Δμ(Li) < −1.65 eV, −2.65 eV < Δμ(O) < 0.11 eV) is 6 times as large as that of LiNiO2,59 LiNiO2 should be synthesized under rigorously controlled experimental conditions to prevent its transition to

Figure 2. Relaxed surface geometries of nonpolar surfaces, (104), (110), and (100) facets, together with representative relaxed structures of polar surfaces, (003), (012), and (101) facets. Violet: Li, red: O, royal blue: Ni. D


Research Article


this result confirms that LiNiO2 forms smaller particles than those of LiCoO2 in experiments.34 Meanwhile, the polar facets represent considerably different relaxed surface geometries from their interior structures. The surface atoms for both the (003) and (012) facets are positioned only on the top layer in the figure, meaning that only atoms in the top layer have dangling bonds. It should be noted that (003)Li and (012)LiNi represent Li for the (003) facet and both Li and Ni for the (012) facets, respectively, playing the role of surface atoms. In contrast, for the (101) facet, all atoms on the three layers including the top one have dangling bonds. Therefore, (101)LiONi is used in this case to denote that Li, O, and Ni on the three layers are all treated as surface atoms, but Li is the outmost one followed by O and Ni. After geometry relaxation, the surface atoms for the polar facets usually move toward the subsurface layer, resulting in a substantial reduction of interlayer spacing compared to that in the interior region (Figure 2). For instance, for the (003) facet, the interlayer spacing between the outmost (Li) and subsurface (O) layers is reduced to 50% of the interior spacing. Next, we investigated the surface energies and stable geometries of the polar facets as a function of the oxygen

the release/intercalation of the Li atoms on nonpolar facets may be facilitated during the charging/discharging process (clearly seen in the top view of Figure S6), owing to the lack of blocking atoms. In comparison, the passage of Li atoms on the (003) facet is likely to be hindered by the Ni−O octahedral structure. The surface energies of the nonpolar facets were calculated and compared with those of LiCoO259 in Table I. Table I. Surface Energies (J/m2) of the Nonpolar Facets facet

LiNiO2

LiCoO259

(104) (110) (100)

0.640 1.244 1.648

1.048 2.241 2.943

The (104) facet shows the lowest surface energy, and the value is considerably lower than those of other nonpolar facets. Interestingly, the sequence of surface energies in LiNiO2 is identical with that in LiCoO2 ((104) < (110) < (100)), but the values of LiNiO2 are half the value of those of LiCoO2. Considering that the slow crystal growth is thermodynamically attributed to the lower surface energy (i.e., more stable surface),

Figure 3. Surface energies of polar facet (a) (003), (b) (012), and (c) (101) surfaces as functions of Δμ(O). (d) Overall surface energies represented as the lowest values for each facet. The vertical black dotted line gives the boundary divided by the equilibrium regions of LiNiO2 phase with spinel (O-rich, left) and with rocksalt (O-poor, right) phases. E


Research Article


Figure 4. (a) Exposed surface areas of each facet. (104): Dark cyan diamond, (003): black square, (101): blue triangle, and (102): red circle. Equilibrium morphologies viewed from top (left) and side (right) direction when Δμ(O) is (b) 0 eV (O-rich), (c) −0.51 eV (the largest area of (104) facet), (d) −0.93 eV (emergence of (102) facet), and (e) −1.14 eV (O-poor). Top view means when looking at the (003) facet from its perpendicular direction.

chemical potential, as shown in Figure 3. The range of Δμ(O) chosen here in the following discussion corresponds to the line along which the LiNiO2 phase has equilibrium with the LiNi2O4 spinel (−0.24 eV < Δμ(O) < −0.16 eV) and NiO rocksalt phases (−1.14 eV < Δμ(O) < −0.24 eV) in the phase diagram (Figure 1). (Note that Δμ(O) = −1.14 and −0.16 eV are defined as the O-poor and O-rich states, respectively.) It should be noted that Δμ(O) is linked to both the temperature and oxygen partial pressure. Thus, Δμ(O) decreases (i.e., moves toward the right direction in the x-axis) upon increasing (decreasing) the temperature (oxygen pressure). The surface energies of polar facets strongly depend on the chemical potential. Accordingly, the stable surface geometries also depend on the environmental change represented by the chemical potential. Figure 3a gives the surface energies of the (003) facet according to the exposed atoms and their ratio on the surface. For the Li surface, the exposed ratio of surface atoms was chosen from fully covered Li (1.0-Li) to 8% Li coverage. For the O surface, in contrast, the ratio of the exposed O atoms was varied from 50% (0.5-O) to 100% (1.0O). The overall tendency for this facet shows that the Liexposed structures have lower surface energies than those of the oxygen-exposed ones irrespective of the ratio of exposed atoms.

Thus, the stable (003) facet has Li atoms on its surface. The Niexposed surface for the (003) facet is highly unstable and cannot maintain a layered structure during the relaxation calculation (not shown in the plots). Meanwhile, as Δμ(O) decreases (e.g., by increasing temperature), the amount of exposed Li increases on the stable surface because the reduction of Δμ(O) induces increased Δμ(Li), namely, a Li-rich environment, in the phase diagram where the surface with more exposed Li becomes more stable. Specifically, in the Orich region where LiNiO2 is in equilibrium with the spinel LiNi2O4 phase, 17% of the fully covered Li monolayer (0.17-Li) remains on the surface. As Δμ(O) is reduced, the stable surfaces are obtained in the order of 33%, 50%, and then 100% of Li on the surface (0.33-Li, 0.5-Li, and 1.0-Li, respectively) at thermodynamic equilibrium. Conversely, for the O-exposed surface (003)O, as Δμ(O) increases, the surface with more exposed O becomes more stable, because the O-rich region favors more exposed O atoms on the surface. Figure 3b shows the changes of surface energies for the (012) facet, which is classified by the O-terminated (012)O and both Li- and Niterminated surfaces (012)LiNi. For (012)O, as Δμ(O) decreases, the surface energy decreases for 25% of oxygen on the surface (0.25-O), and increases for the 75% (0.75-O) and F


Research Article


Figure 4 shows the changes of (a) surface area for each facet and (b−e) the corresponding equilibrium morphology according to Δμ(O). The equilibrium morphology, called the Wulff shape,60 is generated to minimize the total surface energy of a fixed volume. This minimum energy shape is determined by the symmetry of the crystal and the surface energies of various facets. For the equilibrium morphology, both (104) and (003) facets are always exposed for the entire range of oxygen chemical potential that satisfies the existence of the LiNiO2 phase. When LiNiO2 equilibrates with the spinel phase (O-rich, low T or high p(O2)), the surface area consists of 88.4% of (104) and 11.6% of (003) facets. When Δμ(O) decreases, the area percentage of the (104) facet grows to 95.9%, and then it is rapidly reduced to 24.8% in the O-poor region (Δμ(O) = −1.14 eV). The (003) facet, whose surface contains only Li termination, occupies 1.3−12.0% of the surface area for the entire range, which is quite small compared to that of the (003) facet of LiCoO2 (around 20−50%).59 Considering that it is difficult for Li atoms to move through the large Ni−O octahedral blocks toward the (003) facet, the ideal LiNiO2 with less (003) area can theoretically have faster charging/ discharging rates than those of LiCoO2. For the (012) facet, the Li-exposed surface area (LiNi_ex) appears at Δμ(O) < −0.88 eV and eventually attains 20.2%. The exposed (101) facet also has only Li atoms in the outmost surface, but it appears in the morphology at Δμ(O) < −0.47 eV and reaches 42.9%, the largest among all facets in the O-poor region. It should be noted that all polar facets exposed on the equilibrium shape are Li-terminated surfaces, but only the nonpolar (104) facet has all of the elements (Ni, O, and Li) on its surface. Considering that the surface atoms are more prone to being released compared to when they are in the bulk position, exposure of the (104) facet should be reduced to prevent O release as well as Ni dissolution. Reduction of the (104) facet can be achieved by decreasing Δμ(O), which can be adjusted by a decrease in oxygen partial pressure, an increase in temperature, or directly adding a reducing agent. The Wulff shape generating the equilibrium morphology can be categorized into four distinct points according to Δμ(O). In the O-rich region (Figure 4b), the shape is truncated trigonal bipyramid composed of both six (104) and two (003) faces, resembling a hexagon in the top view and an octagon in the side view, with a flat (003) facet on the top/bottom. When the area of the (104) facet is the largest at Δμ(O) = −0.51 eV (Figure 4c), the area of the (003) facet is noticeably reduced, and the corners of the polygon are slightly truncated by the emergence of the (101) facet, generating a dodecahedron. As the (012) facet emerges on the surface at Δμ(O) = −0.93 eV (Figure 4d), the shape starts to possess four different facets (18 faces), although most of the area is occupied by the (104) and (101) facets. The shape appears spherical in the top view and egg-shaped in the side view, surrounded by six (101) facets around the middle region. In the O-poor region where Δμ(O) is the lowest (Figure 4e), four different facets occupying similar portions of surface area are clearly shown, and the overall shape looks like an egg (side view) truncated by the (003) facet. If we consider a synthetic temperature above 900 K, the Wulff shape of LiNiO2 corresponds to a truncated egg surrounded by four distinct facets. After cooling to ambient conditions, the truncated egg shape can be maintained if a high kinetic barrier exists, otherwise it changes to the thermodynamically stable form of the truncated trigonal bipyramid.

100% oxygen (1.0-O) remaining surfaces. The structure with 50% of oxygen remaining on the surface presents a constant energy with respect to Δμ(O). It is also clear that as Δμ(O) decreases, the equilibrium state favors the oxygen-deficient condition, and accordingly, the O-poor surfaces become more stable and the O-rich surfaces become unstable. In the region of stable LiNiO2 phase, the surface with 50% oxygen vacancy is the most stable among the different O-terminated surfaces, which suggests that a substantial amount of surface oxygen atoms may already be released before battery operation if the (012)O facet becomes uncovered. However, the (012)O facet is not exposed in the equilibrium morphology, because its surface energy is relatively high (∼1.1 J/m2) compared to those of the other facets (see Figures 3d and 4a). For (012)LiNi, the ideal structure (LiNi) in which the surface is composed of both Li and Ni atoms is extremely unstable in the entire region. However, when the surface Ni is all exchanged with the subsurface Li atoms (LiNi_ex, i.e., only Li atoms exposed on surface), the surface energy is markedly reduced by 1.07 J/m2 than that of the ideal (102)LiNi case. Furthermore, in the Opoor region (Δμ(O) < −0.73 eV), the exchanged surface is the most stable one for the (012) facet. This also demonstrates that the exposed surface Ni atoms destabilize the facet, like the case of (003). However, unlike the (003) facet, whose stable surfaces are all Li-exposed ones in the entire region, (012)O is the most stable for the (012) facet in the O-rich region, although the stability of the Li-exposed surface noticeably increases as Δμ(O) goes to the O-poor region. The change of exposed elements on the stable surface occurs more abruptly for the (101) facet with decreasing Δμ(O), as shown in Figure 3c. The ONiO surface (O is at the outmost, followed by Ni and O) is the most stable in the O-rich region, specifically when LiNiO2 is in equilibrium with the spinel LiNi2O4 phase. In the O-poor region (Δμ(O) < −0.34 eV), the Li-exposed surface (LiONi) becomes the most stable, which is similar to the other polar facets of (003) and (012). Meanwhile, when the Ni is exposed on the outmost surface (NiOLi), the structure is very unstable. However, the exchange of Ni on the top layer with the sublayer Li (NiOLi_ex) reduces the surface energy by 0.94 J/m2, which is similar to the Li−Ni exchange on the (012) facet. Overall, the surface energies for the three polar facets reveal the following findings: The Li-exposed surface is the most stable, and the Ni-exposed one is the least stable. The exposed Ni atoms on the surface destabilize the facet, so that these Ni atoms can be exchanged with the subsurface Li atoms to lower the surface energy. For the (003) facet, the amount of Li atoms on the surface can be varied by tuning Δμ(O). However, for both (012) and (101) facets, the fully covered Li surface is more stable compared to the partially covered one irrespective of Δμ(O). Figure 3d represents the overall surface energies for all facets considered here, showing the lowest value for each facet according to Δμ(O). The nonpolar (104) facet has the lowest surface energy in most regions where the LiNiO2 phase is stable. In the O-poor region of Δμ(O) < −1.0 eV, the (101) facet gives the lowest value instead, and the (012) facet also has comparable surface energy to (104). In the spinel equilibrium region, the surface energy of the (003) facet is comparable to that of the (104) facet. Yet as Δμ(O) decreases, the gap between these facets becomes large (0.17-Li and 0.33Li), levels off (0.5-Li), and then is reduced (1.0-Li). Meanwhile, other nonpolar facets, such as the (100) and (110) surfaces, have considerably higher surface energies than those of the polar facets. G


Research Article


noted that the LiNi_ex surface of the (012) facet is composed of only Li atoms, like the 1.0-Li of (003). Moreover, for the LiONi surface of the (101) facet, the Li, O, and Ni elements located on the first three layers are all exposed on the surface although Li resides on the outmost, so that the Fermi level of this surface can be comparable to the level of the 0.33-Li surface of the (003) facet. For the oxygen-exposed surfaces, which are thermodynamically less stable than the Li-exposed ones, the Fermi energies are positioned quite below those of the Li surfaces, and the energy levels decrease with increasing surface O content. On the basis of the level alignments, it is implied that the Fermi level of the surface is affected by the character of exposed atoms as well as their contents, rather than the facet direction. Specifically, the surface Li atoms prevent electron transfer from the electrolyte to the cathode, thereby maintaining an electrochemically stable surface. However, once the surface Li is extracted and then oxygen is exposed on the surface, the Fermi energy becomes lower to be comparable to the HOMO energy. The electron transfer from the electrolyte then becomes easier, which can induce electrolyte decomposition reactions. Several experiments propose that the high impedance and reactivity of LiNiO2 can be attributed to the surface oxygen, which can act as a nucleophile to cause electrolyte decomposition.63,64 In other words, the surface oxygen can donate electrons to the electrolyte in the form of nucleophilic attack. This hypothesis is based on the assumption that contrary to the anode, chemical reaction is dominant at the cathode whereas electrochemical reaction hardly occurs.62 However, our calculation strongly suggests that the electrochemical reaction by thermodynamic driving forces can also occur, if the oxygen atoms are exposed on the surface after the Li atoms are removed during battery operation.

Figure 5 shows the aligned Fermi levels of each facet with respect to the vacuum energy (0 eV), according to the exposed

Figure 5. Aligned Fermi level of each facet with respect to the vacuum level of 0 eV. The solid lines represent the stable surfaces on the morphology, whereas dotted lines correspond to thermodynamically unstable surfaces. The experimental HOMO level of electrolyte (EC/ DMC with LiPF6 salt) is indicated by the green solid line for comparison.

atoms. Solid lines indicate those dominant surface terminations following the Wulff analysis, whereas dotted lines correspond to those surfaces that are thermodynamically unstable within the computed range of Δμ(O). To assign the vacuum level, the electrostatic potentials were averaged in the direction parallel and plotted in the direction perpendicular to the surface on each facet, and the stationary potential values were set to the vacuum energy (Figure S1). For the electrolyte, the highest occupied molecular orbital (HOMO) energy (−7.5 eV) corresponding to the Fermi energy of solid material was adopted from the experimental value61 of a common mixed electrolyte (ethylene carbonate (EC) and dimethyl carbonate (DMC) with 1.2 M LiPF6 salt). From the electrochemical viewpoint, the electron transfer between the electrode and the electrolyte facilitates electrolyte decomposition on the electrode surface. On the anode side, the electrons at the Fermi level in the anode can be transferred to the lowest unoccupied molecular orbital (LUMO) level of the electrolyte, leading to reductive decomposition of the electrolyte. Actually, the Fermi level of graphite (a representative anode material) is positioned higher than the LUMO of electrolytes, so that the formation of a solid−electrolyte interphase from the electrolyte decomposition on the anode surface is frequently observed.62 On the contrary, on the cathode side, the HOMO electron of electrolyte can be transferred to the cathode, inducing oxidative decomposition of the electrolyte. Therefore, the Fermi level of the cathode should be higher in energy than the HOMO of electrolyte to prevent electron transfer from the electrolyte to the cathode. For the nonpolar facets in Figure 5, the Fermi levels do not show any facet dependence, and the energies (between −5.2 and −5.4 eV) are about 2 eV higher than the HOMO energy of electrolyte. This indicates that the electrolyte can be stable around the nonpolar surfaces (at least prior to battery operation), because an energy barrier of 2 eV should make it difficult for the HOMO electron of the electrolyte to transfer to the nonpolar facets. Interestingly, for the polar facets, as the amount of surface Li decreases, the energy of the Fermi level goes down irrespective of the facets. It should be

IV. CONCLUSIONS We investigated the surface energy, equilibrium morphology, and electrochemical window of LiNiO2 cathode materials within the DFT framework. Both polar and nonpolar facets were considered to predict the change of surface properties, and their electrochemical stability with respect to the electrolyte under different environments. Contrary to the nonpolar facets, the surface properties of the polar facets can be affected by the environment, which is expressed as the oxygen chemical potential. The calculated phase diagram of Li−Ni−O materials delineates the chemical potential range permitting the stable LiNiO2 phase and accordingly, the surface properties of each facet could be interpreted within the meaningful range of the potential. It was found that Li-terminated surfaces have lower surface energies than those of the oxygen surfaces irrespective of facet direction, and Ni-exposed ones are the least stable. Thus, when the surface Ni atoms are exchanged with sublayer Li, a substantial increase in the surface stability is predicted for the (012) and (101) facets. The equilibrium morphology was suggested based on Wulff theory, using both the symmetry of the crystal and the surface energies of various facets. In the Orich region (low T and high p(O2)), the morphology is composed of mainly the (104) facet and a minority of the (003) facet, generating a truncated trigonal bipyramid. As Δμ(O) decreases (high T and low p(O2)), all four ((104), (003), (102), and (101)) facets apparently occupy considerable portions of the morphology, which is shaped as a truncated egg. Finally, we also calculated the Fermi energies of stable surfaces on each facet, and the values were aligned with respect to the H


Research Article


(10) Ding, Y.; Mu, D.; Wu, B.; Wang, R.; Zhao, Z.; Wu, F. Recent Progresses on Nickel-Rich Layered Oxide Positive Electrode Materials used in Lithium-Ion Batteries for Electric Vehicles. Appl. Energy 2017, 195, 586−599. (11) Li, J.; Downie, L. E.; Ma, L.; Qiu, W.; Dahn, J. R. Study of the Failure Mechanisms of LiNi0.8Mn0.1Co0.1O2 Cathode Material for Lithium Ion Batteries. J. Electrochem. Soc. 2015, 162, A1401−A1408. (12) Schipper, F.; Erickson, E. M.; Erk, C.; Shin, J.-Y.; Chesneau, F. F.; Aurbach, D. ReviewRecent Advances and Remaining Challenges for Lithium Ion Battery Cathodes I. Nickel-Rich, LiNixCoyMnzO2. J. Electrochem. Soc. 2017, 164, A6220−A6228. (13) Hausbrand, R.; Cherkashinin, G.; Ehrenberg, H.; Gröting, M.; Albe, K.; Hess, C.; Jaegermann, W. Fundamental Degradation Mechanisms of Layered Oxide Li-Ion Battery Cathode Materials: Methodology, Insights and Novel Approaches. Mater. Sci. Eng., B 2015, 192, 3−25. (14) Spotnitz, R.; Franklin, J. Abuse behavior of High-Power, Lithium-Ion Cells. J. Power Sources 2003, 113, 81−100. (15) Bandhauer, T. M.; Garimella, S.; Fuller, T. F. A Critical Review of Thermal Issues in Lithium-Ion Batteries. J. Electrochem. Soc. 2011, 158, R1−R25. (16) Wang, L.; Maxisch, T.; Ceder, G. A First-Principles Approach to Studying the Thermal Stability of Oxide Cathode Materials. Chem. Mater. 2007, 19, 543−552. (17) Min, K.; Kim, K.; Jung, C.; Seo, S.-W.; Song, Y. Y.; Lee, H.; Shin, J.; Cho, E. A Comparative Study of Structural Changes in Lithium Nickel Cobalt Manganese Oxide as a Function of Ni Content during Delithiation Process. J. Power Sources 2016, 315, 111−119. (18) Min, K.; Seo, S.-W.; Song, Y. Y.; Lee, H.; Cho, E. A FirstPrinciples Study of the Preventive Effects of Al and Mg Doping on the Degradation in LiNi0.8Co0.1Mn0.1O2 Cathode Materials. Phys. Chem. Chem. Phys. 2017, 19, 1762−1769. (19) Johannes, M. D.; Swider-Lyons, K.; Love, C. T. Oxygen Character in the Density of States as an Indicator of the Stability of LiIon Battery Cathode Materials. Solid State Ionics 2016, 286, 83−89. (20) Nayak, P. K.; Grinblat, J.; Levi, M.; Levi, E.; Kim, S.; Choi, J. W.; Aurbach, D. Al Doping for Mitigating the Capacity Fading and Voltage Decay of Layered Li and Mn-Rich Cathodes for Li-Ion Batteries. Adv. Energy Mater. 2016, 6, No. 1502398. (21) Aurbach, D.; Srur-Lavi, O.; Ghanty, C.; Dixit, M.; Haik, O.; Talianker, M.; Grinblat, Y.; Leifer, N.; Lavi, R.; Major, D. T.; et al. Studies of Aluminum-Doped LiNi0.5Co0.2Mn0.3O2: Electrochemical Behavior, Aging, Structural Transformations, and Thermal Characteristics. J. Electrochem. Soc. 2015, 162, A1014−A1027. (22) Pouillerie, C.; Croguennec, L.; Biensan, P.; Willmann, P.; Delmas, C. Synthesis and Characterization of New LiNi1−yMgyO2 Positive Electrode Materials for Lithium Ion Batteries. J. Electrochem. Soc. 2000, 147, 2061−2069. (23) Liu, W.; Oh, P.; Liu, X.; Myeong, S.; Cho, W.; Cho, J. Countering Voltage Decay and Capacity Fading of Lithium-Rich Cathode Material at 60 °C by Hybrid Surface Protection Layers. Adv. Energy Mater. 2015, 5, No. 1500274. (24) Kong, F.; Liang, C.; Longo, R. C.; Yeon, D.-H.; Zheng, Y.; Park, J.-H.; Doo, S.-G.; Cho, K. Conflicting Roles of Anion Doping on the Electrochemical Performance of Li-Ion Battery Cathode Materials. Chem. Mater. 2016, 28, 6942−6952. (25) Yamamoto, S.; Noguchi, H.; Zhao, W. Improvement of Cycling Performance in Ti Substituted 0.5Li 2MnO3−0.5LiNi0.5 Mn0.5 O2 through Suppressing Metal Dissolution. J. Power Sources 2015, 278, 76−86. (26) Song, B.; Zhou, C.; Wang, H.; Liu, H.; Liu, Z.; Lai, M. O.; Lu, L. Advances in Sustain Stable Voltage of Cr-Doped Li-Rich Layered Cathodes for Lithium Ion Batteries. J. Electrochem. Soc. 2014, 161, A1723−A1730. (27) Fey, G. T. K.; Chen, G.; Subramanian, V.; Osaka, T. Preparation and Electrochemical Properties of Zn-Doped LiNi0.8Co0.2O2. J. Power Sources 2002, 112, 384−394. (28) Min, K.; Park, K.; Park, S. Y.; Seo, S.-W.; Choi, B.; Cho, E. Improved Electrochemical Properties of LiNi0.91Co0.06Mn0.03O2

vacuum level to predict the electrochemical stability. The surface Li pushes up the Fermi level for the polar facets, therefore it could prevent the electrons of the electrolyte from transferring to the facet to maintain an electrochemically stable surface. However, exposure of the oxygen atoms lowers the level and thus, electron transfer from the electrolyte becomes easier, which can induce electrolyte decomposition reactions.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b08563. Comparison of symmetric with asymmetric structures, predicted morphology by BFDH theory, inverse relation between Δμ(O) and T, correction value of Ni, ternary phase diagram of Li−Ni−O at 0 K, top view of nonpolar facets, surface geometries of polar facets, total energies for the construction of Li−Ni−O phase diagram, slab information for modeling (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Eunseog Cho: 0000-0001-5308-8278 Kyoungmin Min: 0000-0002-1041-6005 Notes

The authors declare no competing financial interest.

■

REFERENCES

(1) Liu, W.; Oh, P.; Liu, X.; Lee, M.-J.; Cho, W.; Chae, S.; Kim, Y.; Cho, J. Nickel-Rich Layered Lithium Transition-Metal Oxide for HighEnergy Lithium-Ion Batteries. Angew. Chem., Int. Ed. 2015, 54, 4440− 4457. (2) Bak, S.-M.; Hu, E.; Zhou, Y.; Yu, X.; Senanayake, S. D.; Cho, S.-J.; Kim, K.-B.; Chung, K. Y.; Yang, X.-Q.; Nam, K.-W. Structural Changes and Thermal Stability of Charged LiNixMnyCozO2 Cathode Materials Studied by Combined In Situ Time-Resolved XRD and Mass Spectroscopy. ACS Appl. Mater. Interfaces 2014, 6, 22594−22601. (3) Noh, H.-J.; Youn, S.; Yoon, C. S.; Sun, Y.-K. Comparison of the Structural and Electrochemical Properties of Layered Li[NixCoyMnz]O2 (x = 1/3, 0.5, 0.6, 0.7, 0.8 and 0.85) Cathode Material for LithiumIon Batteries. J. Power Sources 2013, 233, 121−130. (4) Zheng, J.; Kan, W. H.; Manthiram, A. Role of Mn Content on the Electrochemical Properties of Nickel-Rich Layered LiNi0.8− xCo0.1Mn0.1+xO2 (0.0 ≤ x ≤ 0.08) Cathodes for Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 2015, 7, 6926−6934. (5) Liang, C.; Kong, F.; Longo, R. C.; KC, S.; Kim, J.-S.; Jeon, S.; Choi, S.; Cho, K. Unraveling the Origin of Instability in Ni-Rich LiNi1−2xCoxMnxO2 (NCM) Cathode Materials. J. Phys. Chem. C. 2016, 120, 6383−6393. (6) Yang, J.; Xia, Y. Enhancement on the Cycling Stability of the Layered Ni-Rich Oxide Cathode by In-Situ Fabricating NanoThickness Cation-Mixing Layers. J. Electrochem. Soc. 2016, 163, A2665−A2672. (7) Ma, X.; Wang, C.; Cheng, J.; Sun, J. Synthesis and Electrochemical Properties of LiNi0.9‑xCo0.1SnxO2 as Cathode Materials for Lithium Secondary Batteries. J. Solid State Electrochem. 2007, 11, 1139−1143. (8) Manthiram, A.; Knight, J. C.; Myung, S.-T.; Oh, S.-M.; Sun, Y.-K. Nickel-Rich and Lithium-Rich Layered Oxide Cathodes: Progress and Perspectives. Adv. Energy Mater. 2016, 6, No. 1501010. (9) Jo, M.; Noh, M.; Oh, P.; Kim, Y.; Cho, J. A New High Power LiNi0.81Co0.1Al0.09O2 Cathode Material for Lithium-Ion Batteries. Adv. Energy Mater. 2014, 4, No. 1301583. I


Research Article

ACS Applied Materials & Interfaces Cathode Material via Li-Reactive Coating with Metal Phosphates. Sci. Rep. 2017, 7, No. 7151. (29) Min, K.; Seo, S.-W.; Choi, B.; Park, K.; Cho, E. Computational Screening for Design of Optimal Coating Materials to Suppress Gas Evolution in Li-Ion Battery Cathodes. ACS Appl. Mater. Interfaces 2017, 9, 17822−17834. (30) Kim, H.; Kim, M. G.; Jeong, H. Y.; Nam, H.; Cho, J. A New Coating Method for Alleviating Surface Degradation of LiNi0.6Co0.2Mn0.2O2 Cathode Material: Nanoscale Surface Treatment of Primary Particles. Nano Lett. 2015, 15, 2111−2119. (31) Xiong, X.; Wang, Z.; Yue, P.; Guo, H.; Wu, F.; Wang, J.; Li, X. Washing Effects on Electrochemical Performance and Storage Characteristics of LiNi0.8Co0.1Mn0.1O2 as Cathode Material for Lithium-ion Batteries. J. Power Sources 2013, 222, 318−325. (32) Vallverdu, G.; Minvielle, M.; Andreu, N.; Gonbeau, D.; Baraille, I. First Principle Study of the Surface Reactivity of Layered Lithium Oxides LiMO2 (M = Ni, Mn, Co). Surf. Sci. 2016, 649, 46−55. (33) Kang, J.; Han, B. First-Principles Study on the Thermal Stability of LiNiO2 Materials Coated by Amorphous Al2O3 with Atomic Layer Thickness. ACS Appl. Mater. Interfaces 2015, 7, 11599−11603. (34) Kim, Y.; Lee, H.; Kang, S. First-Principles and Experimental Investigation of the Morphology of Layer-Structured LiNiO2 and LiCoO2. J. Mater. Chem. 2012, 22, 12874−12881. (35) Xu, J.; Lin, F.; Nordlund, D.; Crumlin, E. J.; Wang, F.; Bai, J.; Doeff, M. M.; Tong, W. Elucidation of the Surface Characteristics and Electrochemistry of High-performance LiNiO2. Chem. Commun. 2016, 52, 4239−4242. (36) Xu, J.; Hu, E.; Nordlund, D.; Mehta, A.; Ehrlich, S. N.; Yang, X.Q.; Tong, W. Understanding the Degradation Mechanism of Lithium Nickel Oxide Cathodes for Li-Ion Batteries. ACS Appl. Mater. Interfaces 2016, 8, 31677−31683. (37) Kresse, G.; Furthmüller, J. Efficiency of Ab-initio Total Energy Calculations for Metals and Semiconductors Using a Plane-Wave Basis Set. Comput. Mater. Sci. 1996, 6, 15−50. (38) Kresse, G.; Furthmüller, J. Efficient Iterative Schemes for Ab initio Total-Energy Calculations Using a Plane-Wave Basis Set. Phys. Rev. B 1996, 54, 11169. (39) Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77, 3865. (40) Dudarev, S. L.; Botton, G. A.; Savrasov, S. Y.; Humphreys, C. J.; Sutton, A. P. Electron-Energy-Loss Spectra and the Structural Stability of Nickel Oxide: An LSDA + U Study. Phys. Rev. B 1998, 57, 1505− 1509. (41) Shishkin, M.; Sato, H. Self-Consistent Parametrization of DFT + U Framework Using Linear Response Approach: Application to Evaluation of Redox Potentials of Battery Cathodes. Phys. Rev. B 2016, 93, No. 085135. (42) Chen, H.; Freeman, C. L.; Harding, J. H. Charge Disproportionation and Jahn-Teller Distortion in LiNiO2 and NaNiO2: A Density Functional Theory Study. Phys. Rev. B 2011, 84, No. 085108. (43) Zhou, F.; Cococcioni, M.; Marianetti, C. A.; Morgan, D.; Ceder, G. First-principles Prediction of Redox Potentials in Transition-Metal Compounds with LDA + U. Phys. Rev. B 2004, 70, No. 235121. (44) Chen, H.; Dawson, J. A.; Harding, J. H. Effects of Cationic Substitution on Structural Defects in Layered Cathode Materials LiNiO2. J. Mater. Chem. A 2014, 2, 7988. (45) Sugiyama, J.; Ikedo, Y.; Mukai, K.; Nozaki, H.; Månsson, M.; Ofer, O.; Harada, M.; Kamazawa, K.; Miyake, Y.; Brewer, J. H.; et al. Low-temperature magnetic properties and high-temperature diffusive behavior of LiNiO2 investigated by muon-spin spectroscopy. Phys. Rev. B 2010, 82, No. 224412. (46) Tasker, P. W. The Stability of Ionic Crystal Surfaces. J. Phys. C: Solid State Phys. 1979, 12, 4977. (47) Moses, A. W.; Flores, H. G. G.; Kim, J.-G.; Langell, M. A. Surface properties of LiCoO2, LiNiO2 and LiNi1−xCoxO2. Appl. Surf. Sci. 2007, 253, 4782−4791. (48) Donnay, J. D. H.; Harker, D. A New Law of Crystal Morphology Extending the Law of Bravais. Am. Miner. 1937, 22, 446−467.

(49) Bravais, A. Du Cristal Considere Comme un Simple Assemblage de Points. In Etude Cristallographiques; Gauthier-Villars: Paris, France, 1866. (50) Friedel, G. Studies on the law of Bravais. Bull. Soc. Fr. Mineral. 1907, 30, 326. (51) Kim, S.; Aykol, M.; Wolverton, C. Surface Phase Diagram and Stability of (001) and (111) LiM2O4 spinel oxides. Phys. Rev. B 2015, 92, No. 115411. (52) Bhattacharya, J.; Wolverton, C. Relative Stability of Normal vs. Inverse Spinel for 3d Transition Metal Oxides as Lithium Intercalation Cathodes. Phys. Chem. Chem. Phys. 2013, 15, 6486. (53) Wolverton, C.; Hass, K. C. Phase Stability and Structure of Spinel-based Transition Aluminas. Phys. Rev. B 2000, 63, No. 024102. (54) Patton, D. C.; Porezag, D. V.; Pederson, M. R. Simplified Generalized-Gradient Approximation and Anharmonicity: Benchmark Calculations on Molecules. Phys. Rev. B 1997, 55, 7454−7459. (55) Hammer, B.; Hansen, L. B.; Nørskov, J. K. Improved Adsorption Energetics Within Density-Functional Theory Using Revised Perdew-Burke-Ernzerhof Functionals. Phys. Rev. B 1999, 59, 7413−7421. (56) Wang, L.; Maxisch, T.; Ceder, G. Oxidation Energies of Transition Metal Oxides within the GGA + U Framework. Phys. Rev. B 2006, 73, No. 195107. (57) Jain, A.; Hautier, G.; Ong, S. P.; Moore, C. J.; Fischer, C. C.; Persson, K. A.; Ceder, G. Formation Enthalpies by Mixing GGA and GGA + U calculations. Phys. Rev. B 2011, 84, No. 045115. (58) NIST-JANAF Thermochemical Tables, 4th ed.; Chase, M. W., Jr., Ed.; The American Chemical Society, 1998. (59) Kramer, D.; Ceder, G. Tailoring the Morphology of LiCoO2: A First Principles Study. Chem. Mater. 2009, 21, 3799−3809. (60) Roosen, A. R.; McCormack, R. P.; Carter, W. C. Wulffman: A Tool for the Calculation and Display of Crystal Shapes. Comput. Mater. Sci. 1998, 11, 16−26. (61) Borodin, O.; Behl, W.; Jow, T. R. Oxidative Stability and Initial Decomposition Reactions of Carbonate, Sulfone, and Alkyl PhosphateBased Electrolytes. J. Phys. Chem. C 2013, 117, 8661−8682. (62) Gauthier, M.; Carney, T. J.; Grimaud, A.; Giordano, L.; Pour, N.; Chang, H.-H.; Fenning, D. P.; Lux, S. F.; Paschos, O.; Bauer, C.; et al. Electrode−Electrolyte Interface in Li-Ion Batteries: Current Understanding and New Insights. J. Phys. Chem. Lett. 2015, 6, 4653− 4672. (63) Aurbach, D.; Markovsky, B.; Salitra, G.; Markevich, E.; Talyossef, Y.; Koltypin, M.; Nazar, L.; Ellis, B.; Kovacheva, D. Review on Electrode-Electrolyte Solution Interactions, Related to cathode Materials for Li-Ion Batteries. J. Power Sources 2007, 165, 491−499. (64) Aurbach, D.; Gamolsky, K.; Markovsky, B.; Salitra, G.; Gofer, Y.; Heider, U.; Oesten, R.; Schmidt, M. The Study of Surface Phenomena Related to Electrochemical Lithium Intercalation into LixMOy Host Materials (M = Ni, Mn). J. Electrochem. Soc. 2000, 147, 1322−1331.

J


Cathode for Li Ion Batteries - ACS Publications - American Chemical

Recommend Documents