Conflicting Roles of Anion Doping on the ... - ACS Publications

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Conflicting Roles of Anion Doping on the Electrochemical Performance of Li-Ion Battery Cathode Materials Fantai Kong,† Chaoping Liang,† Roberto C. Longo,† Dong-Hee Yeon,‡ Yongping Zheng,† Jin-Hwan Park,‡ Seok-Gwang Doo,*,‡ and Kyeongjae Cho*,† †

Materials Science & Engineering Department, The University of Texas at Dallas, Richardson, Texas 75080, United States Energy Laboratory, Samsung Advanced Institute of Technology, Samsung Electronics, Yongin 446-712, Republic of Korea

‡

S Supporting Information *

ABSTRACT: Anion doping is one of the most widely adopted strategies to improve the electrochemical performance of cathode materials for Li-ion batteries. However, undesirable side effects are often observed together with enhanced electrochemical properties, leading to an unsatisfactory overall performance. In order to develop an anion doping strategy which enhances the positive effects and suppresses undesirable side effects, the understanding of their origin at the atomic scale is a crucial step. In this work, using density functional theory (DFT), we report a systematic study on the effects of three common anion dopants (F, S, Cl) on a wide range of properties of a model cathode material, LiNiO2, including redox potential, ionic conductivity, Li/Ni exchange, lattice distortion, and Ni migration upon delithiation. The results show that the dopants improve certain properties but worsen others, revealing some distance-dependent features. Overall, our work shows conflicting roles of anion doping on the battery voltage, rate performance, and structural stability of the cathode material. By identifying the origins of the different roles, we propose a rational anion doping strategy for the optimization of the overall electrochemical performance of the cathode material. These results for LiNiO2 can also promote anion doping studies and improved materials design in other Ni-rich layered oxide cathode materials.

1. INTRODUCTION Driven by the growing demand for clean-efficient energy storage media in cell phones, smart grid, electric vehicles, and so forth, Li-ion batteries (LIB) are currently under extensive worldwide study. To improve the electrochemical performance of cathode materials in LIB, anion doping has been one of the widely adopted strategies. Compared with cation doping, it reveals several unique advantages: (1) Cation dopants are often reported to play a limited role in preventing surface Li or transition metal (TM) ion dissolution at elevated temperatures, ascribed to the presence of HF and related acids in the liquid electrolyte.1 However, by forming oxyfluoride through F doping, cathode materials show strengthened immunity to dissolution in HF-related acids.1,2 (2) Anion doping also avoids the undesirable occupation of Li or TM ion sites, which might consume Li or redox active TM ions, thus impairing the achievement of the theoretical energy density. (3) Previous studies focused mainly on the TM ion redox activity, when determining the cathode materials capacity, while recent research based on a class of 4d TM-based Li-rich oxides highlighted the significance of the (O2)n− anionic redox process, which opens a new gate for further increase of the energy density.3,4 Therefore, an anion doping strategy that directly modifies such anionic redox processes would be a © 2016 American Chemical Society

promising direction in the near future for Li-rich oxides engineering. Among the ideal candidates for anion doping, F, S, and Cl are the most widely used, due to their similar atomic radii with O and the easy realization using conventional synthesis methods. These dopants have proven to effectively enhance the electrochemical performance of many cathode materials.5−13 Undesirable side effects have also been frequently observed after anion doping, along with enhanced electrochemical properties, which ultimately leads to an unsatisfactory overall electrochemical performance of the doped oxides. For instance, F was reported to suppress the formation of FeLi antisite defects in LiFePO4, thus enhancing the rate performance, but the achievable discharge capacity was reduced upon high doping concentration.5,14 Another typical example is the S-doped spinel LiMn2O4. In spite of the increased rate performance and thermal stability after S doping, lowered electronic conductivity was also observed.15,16 Similar conflicting effects were also observed in high energy density (>800 Wh/kg) cathode materials: LiNiO2 and its derivative LiNiO2-based oxides Received: June 28, 2016 Revised: September 6, 2016 Published: September 9, 2016 6942

DOI: 10.1021/acs.chemmater.6b02627 Chem. Mater. 2016, 28, 6942−6952

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Chemistry of Materials Li[Ni1−xMx]O2 (x = 0.1−0.2, M = Ni, Co, Mn, Al, and so forth).17−19 It was reported that F increased the capacity retention,20−22 but also incorporated more Li/Ni antisite defects, increased the charge-transfer resistance,23 and reduced the initial delivered capacity.22 F-doped LiNi0.8Co0.15Al0.05O2 was shown to improve the rate performance from 0.2C to 2C current rates, but it deteriorated the performance substantially at high rates beyond 5C.21 Finally, S revealed the promising effect of suppressing the capacity degradation of LiNiO2. However, it also lowered the battery voltage, which reduces the energy density.24 Therefore, the effects of anion doping on cathode materials are somehow conflicting: some properties are improved, while others are deteriorated. From the materials design viewpoint, positive effects should be enhanced whereas undesirable side effects should be suppressed. To achieve this goal, the understanding of the origin of both positive and side effects at the atomic scale is the first crucial step. In this work, we report a systematic study of the anion (F, S, Cl) doping effects on a model layered cathode material, LiNiO2, using densityfunctional theory (DFT). The properties through examination include redox potential, ionic conductivity, Li/Ni exchange, interlayer Ni migration in the delithiated Li1−xNiO2 and lattice distortion during cycling. Our results reveal the conflicting effects of the dopants on the different properties, explaining many experimental observations. Furthermore, by analyzing the origin of the conflicting roles, we propose possible doping strategies that might take advantage of the positive effects while addressing undesirable side effects.

Figure 1. Configuration of Li18Ni18O36 with one oxygen atom substituted by an anion dopant X (X = F, Cl, or S).

3. RESULTS AND DISCUSSION 3.1. Battery Voltage. Li-ion battery voltage is determined by the redox potential difference between the cathode and the anode. The redox potential describes the energetic effort required to extract Li ions from the electrodes, with the optimized range for cathode materials located between 3.5 and 4.8 V. A potential lower than 3.5 V would affect the delivered energy density, while a value higher than 4.8 V would cause severe electrolyte degradation and electrode/electrolyte side reactions. In this section, the anion doping effect on the redox potential during the whole delithiation range has been estimated using the following formula:

2. COMPUTATIONAL METHODS DFT calculations are performed using the Vienna ab initio simulation package (VASP)25,26 with the projector-augmented wave (PAW) method.27 The generalized gradient approximation Perdew−Burke− Ernzerhof (GGA-PBE) exchange and correlation functional is used throughout all the calculations.28 A plane wave cutoff energy of 500 eV is employed to guarantee the accuracy of the obtained wave functions. The structure optimization is performed until the force on each atom is less than 0.05 eV/Å, and the total energy is converged up to 10−4 eV. The self-consistent calculations are carried out with the tetrahedron method with Blöchl corrections.29 The energy convergence criteria is set to be 10−5 eV. To describe the localized nature of 3d electrons more accurately, we also employed the effective on-site Hubbard Ueff correction on the 3d electrons of Ni.30 The Ueff value used is 6.4 eV, taken from previous theoretical reports.6,31 The supercell used in this work contains 18 unit cells: Li18Ni18O36. The anion doping is modeled by replacing one oxygen atom with one anion atom X (X = F, Cl, or S). Therefore, the model represents an anion doping concentration of 2.78%, corresponding to the LiNiO1.944X0.056 stoichiometry. The Li18Ni18O36 supercell and the doping site are illustrated in Figure 1. It has been reported that DFT+U cannot accurately calculate the interlayer distance of layered oxides under highly delithiated stages, due to insufficient consideration of van der Waals (vdW) interactions.32 Therefore, in this work, the DFT+D3 method33 for which the dispersion coefficients depend on local structural geometry, was also included in the calculation of highly delithiated states (for Li concentrations equal or below 25%). All the structural figures in this work were plotted using the VESTA software.34 The Climbing ImageNudged Elastic Band (CI-NEB) method is adopted to calculate the kinetic energy barriers during Li ion and Ni ion migrations.35 Analysis of Bader charges is a widely adopted method to study charge distribution. Bader charges are determined by the zero-gradient surface of electron density between ions. In this work, the Bader charge analysis was performed using grid-based charge density decomposition, as developed by Henkelman et al.36

V=−

E(LiNiO1.944 X 0.056) − E(Li1 − xNiO1.944 X 0.056) − x·E(Li) x·e

where E(LiNiO1.944X0.056), E(Li), and E(Li1−xNiO1.944X0.056) denote the total energy of the lithiated state, bulk BCC Li metal, and delithiated state, respectively. On the basis of previous experimental observations and theoretical predictions, the entire charging process can be approximately divided into four stages with the formation of three intermediate stable structures at Li concentrations of 0.75, 0.5, and 0.25.37,38 The full charging profiles of pristine and doped LiNiO2 have been determined accordingly, and plotted in Figure 2, where the experimental charging profile of pristine LiNiO2 is also included for comparison, showing good consistency with the current result. During the first delithiation stage, x < 0.25, S and Cl reduce the redox potential by 0.2 V, and F shows a relatively weaker effect, lowering the potential only 0.06 V. In the second stage, 0.25 < x < 0.5, S and Cl lower the potential by 0.17 and 0.07 V, respectively, with F showing the opposite effect of increasing the potential by 0.14 V. These calculations agree with experimental reports on the early state charging profile after F and S doping: F increased the voltage of LiNiO2-based NCA and NCM cathode materials by 0.2−0.3 V,21−23 while an S-doping concentration of 1% lowered the voltage of LiNiO2 by 0.1−0.2 V.24 At the stage of 0.5 < x < 0.75, the three dopants increase the potential by 0.15 V. Finally, in the last charging stage, doping effects on the voltage tend to be negligible, something expected considering the reduced amount of LiX bonds. The experimental literature reported similar features: 6943


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our analysis (see Figure S1a−c of the Supporting Information) shows that F could also modify the interactions of neighboring ions with Li by reducing second nearest neighboring Ni ions, thus lowering the average redox potential. Note that Figure 2 also shows that F lowers the redox potential at the earliest and last charging stages. Such effect on the earliest stage can be attributed to the charge balance requirement of a neutral system. Once O2− is replaced with an X− anion, the system is one electron deficient, which facilitates the extraction of positively charged Li+ in order to maintain the overall charge neutrality. A similar mechanism has been identified in multivalent Li-site doped Mn oxides.40 The Bader charge analysis in Table 1 confirms the 1− charge state of F, Cl and S anion dopants (−0.87e−, −0.78e−, and −0.57e−, respectively). With respect to the last stage of charging, it might be related with the screening effect of F in reducing Ni ions, as illustrated in Figure S1c of the Supporting Information. In summary, F doping enhances and Cl and S doping reduces the charging voltage of LiNiO2. However, given the low doping concentration typically considered (∼3%), the effect must be relatively weak: the voltage variation obtained in this work is less than 0.2 V. This result confirms that a 3% anion doping concentration will not be detrimental for the energy density of Ni oxides and will also not show intense effects on the battery voltage. However, further increase in doping concentration might have stronger and more severe consequences. Therefore, F doping is a good candidate for high voltage cathode materials design, while a large doping concentration of S and Cl should be avoided, due to the negative effects on the energy density. 3.2. Rate Performance. The rate performance of a Li-ion battery is a critical concern for large-current applications, such as electrical vehicles, fast charging devices, etc. The rate performance of the battery is mainly determined by the Li ionic conductivity inside the electrodes. For layered oxides, it refers to the intralayer diffusion efficiency. Additionally, LiNiO2 faces one more severe problem: Li/Ni antisite defects. These defects are normally produced at a high ratio of 5−10% in the aftersynthesized powders,18,41 which could ultimately block Li migration paths and increase diffusion resistance.42 Our simulation supports the deteriorating effect of Li/Ni antisite defect on Li ionic conductivity. As shown in Figure S2 and Figure S3 of Supporting Information (SI), the Ni atom in Li layer is very immobile with a migration barrier as high as 1.3 eV. And due to the higher charge state of Ni, the Li migration barrier was also increased. With the above considerations in mind, in this section we estimate the effect of anion doping on both Li ionic conductivity and formation of Li/Ni antisite defects. 3.2.1. Ionic Conductivity. It is widely accepted that the ionic conduction mechanism in layered oxides is Li vacancy (VLi) diffusion,6,43,44 with the diffusivity showing an exponential dependence on the energy barrier for Li hopping between neighboring lattice sites. Therefore, VLi migration energy barriers in pristine and anion doped LiNiO2 must be obtained to examine the doping effect on the ionic conductivity. In most layered structures, including LiMnO2, LiCoO2, LiNi0.5Mn0.5O2, and so forth, VLi hops between neighboring octahedral sites (Oh) through an intermediate tetrahedral site (Td), with the assistance of a neighboring VLi, known as tetrahedral site hop (TSH).6,43,44 The TSH configurations in pristine and doped LiNiO2 are shown in Figure 3a−d, from which we can estimate how a neighboring dopant affects Li migration. The obtained migration energy barriers are shown in Figure 3e. For pristine

Figure 2. Charging profiles of pristine and anion doped LiNiO2. The inset shows the configurations of the three intermediate stable states at Li concentrations of 0.25, 0.5, and 0.75. (a: ref 38, b: ref 39). Note that the real voltage profile after anion doping might be slightly different, especially at higher doping concentrations, since the doping effect on Li/vacancy ordering may be not negligible.

the doping effect on charging profiles at the late stage tends to be substantially weaker.21−24 In order to develop a deeper understanding of the doping effects, Table 1 gives the average redox potential for pristine Table 1. Average Redox Potential, Li-Dopant Bonding Length (dLi‑dopant), Bader Charge State of Anion Dopant (Qdopant) and Electronegativity (Pauling) of Anion Dopant in Pristine and Doped LiNiO2 pristine LiNiO2 average redox potential (V) dLi‑dopant (Å) Qdopant (e−) electronegativity

F-doped LiNiO2

Cl-doped LiNiO2

S-doped LiNiO2

3.92

3.97

3.87

3.82

2.12 −1.20 3.44

2.05 −0.88 3.98

2.23 −0.78 3.16

2.39 −0.57 2.58

and doped LiNiO2. The table clearly shows that the overall effect of F is to increase the battery voltage whereas Cl and S lower it. Since voltage reflects the energetic effort required to extract Li ion from the electrode materials, Li-anion bonding strength plays a determinant role. To analyze the Li-dopant bonding strength, the corresponding bonding length and Bader charge states of anion dopants were obtained (See Table 1). The bonding length of LiCl and LiS are 0.11 and 0.27 Å longer than that of LiO in pristine LiNiO2, respectively. Owing to the lower electronegativity, Cl and S ions also have higher Bader charge states (−0.78e− and −0.57e− compared with −1.20e− for O). Considering the ionic nature of Li-anion bonds, an elongated bond length and increased anion Bader charge state lead to relatively weaker Li-dopant bonding strength, which lowers the redox potential. However, larger Bader charge state of F and shorter LiF bonds reveal a similar bonding strength between LiO and LiF. In fact, 6944


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Figure 4. (a) Li ion migration path, with a distance between the migrating ion and the anion dopant of ∼5 Å, and (b) the corresponding migration energy barriers.

these results imply that S and Cl-doping show conflicting effects on the ionic conductivity depending on their relative distances to migrating Li ions: diffusion of neighboring Li ions shows an increased energy barrier, whereas that of distant Li ions is eased by lowering the energy barrier. To understand the underlying mechanism that determines such distance-dependent effects on Li migration, we have investigated the doping effect on the area of Oh−Td shared plane during TSH migration (Figure 5a). The plane area estimates the available spacing for TSH migration and has been proved key in determining the migration energy barrier of many layered oxides.42 The areas of this Oh−Td shared plane neighboring and distant to Cl or S dopants are very different. A narrowed neighboring plane indicates higher spacial resistance for Li migration, which increases the corresponding energy barrier. On the contrary, an expanded plane distant to the dopants confirms the weakened spacial resistance for Li migration and the expected lower energy barrier. Although F also narrows the Oh−Td shared plane, it might be partially offset by the relative smaller radius of F− (1.36 Å), as compared to O2− (1.40 Å). By analyzing the saddle point structure of the Li diffusion path in F-doped LiNiO2, we found that the lower bader charge of F ion with respect to O might also contribute to the lower Li migration barrier (see Figure S4). Additionally, the F effect on the plane areas is distance-independent, which is consistent with the migration energy barriers shown in Figures 3 and 4. 3.2.2. Li/Ni Antisite Defects. To investigate the anion doping effects on Li/Ni antisite defects, we have used the model shown in Figure 6a. One of the 18 Ni atoms of the supercell is exchanged with a neighboring Li ion to form a Li/Ni antisite defect, which corresponds to an exchange ratio of ∼5.56%, consistent with the obtained experimental range of 5−10%.18,41

Figure 3. Illustration of Li migration paths in (a) pristine LiNiO2, (b) F-doped LiNiO2, (c) S-doped LiNiO2, (d) Cl-doped LiNiO2, and (e) the corresponding migration barriers.

LiNiO2, the barrier is 0.33 eV, relatively close to the previously reported value of 0.31 eV.42 F-doping lowers the barrier by 0.08 eV, which would correspond to an increased hopping rate by a factor of 23 at room temperature [∼exp(0.08 eV/kT)]. On the contrary, the TSH path could be ultimately blocked by S and Cl dopants. Figure 3c,d implies that the migrating Li ion does not pass through the Td site, but through the neighboring “oxygen dumbbell”43 to the VLi site. This site introduces an intermediate stable state with the migrating Li ion occupying the VLi site (Figure 3e). In fact, this result originates from the weaker LiS and LiCl bonds: Li forms five LiO bonds and one Lidopant bond at the initial Oh site, but six LiO bonds at the intermediate VLi site. Figure 3e shows that the “oxygen dumbbell” intermediate state increases the barrier to ∼0.55 eV for hopping between Oh sites. To determine the range of the doping effect on Li migration (specially considering the low doping concentration in the present model, 2.78%), we also obtained the migration barriers of non-neighboring Li ions. The configurations and corresponding energy barriers are given in Figure 4, in which the migrating Li ion and the dopant X are located at different layers, with a relative distance of ∼5 Å. It shows that F-doping has a similar effect than neighboring doping, lowering the barrier by only 0.06 eV. However, instead of impairing Li ion migration, S and Cl unexpectedly facilitate ionic diffusion, with a decrease in the energy barrier by 0.17 and 0.23 eV, respectively. The hopping rate could then be increased by a factor of as high as 102−103 at room temperature. Therefore, 6945


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As illustrated in Figure 6a, the Ni ion occupying the Li site forms six new 180° NiNi pairs with Ni ions in neighboring Ni sites. An oxygen ion or the anion dopant resides in the middle of the pair and acts as the bridging atom. NiNi pairs have been reported to stabilize Li/Ni antisite defects due to the strong 180° NiONi superexchange interaction.45 Here, we also find that the strength of NiONi superexchange interactions strongly depends on the spin state of the two bridging Ni ions. As shown in Figure 6b, if the Ni ions are in a parallel spin state, then the formation energy is 0.70 eV, consistent with a previous work.45 However, when the Ni ions are in antiparallel spin state, the Li/Ni antisite defect shows a decreased formation energy of 0.52 eV. This reduction originates from the nature of the superexchange interaction. Our Bader charge analysis shows that both Ni ions in the pair are in a 2+ charge state, which corresponds to an electron configuration of t2g6eg2. (The charge compensation is achieved through charge-disproportionation of neighboring Ni ions, see SI for details.) On the basis of the semiempirical rules set by Goodenough and Kanamori, half-filled TM layers favor such 180° superexchange configurations.46,47 Figure 7 illustrates that

Figure 5. (a) Illustration of the Oh−Td shared plane along Li migration path. (b) The area of this plane under different doping situations.

Figure 7. Schematic representation of the 180° superexchange interaction (Ni2+−pz−Ni2+).

the pz orbital of anion has a superexchange interaction with the dz2 orbital of Ni ions along the y axis. Since the two electrons occupying the pz orbital are in opposite spin states, neighboring Ni ions with opposite spin states will be thermodynamically more favorable. In order to evaluate the doping effect on Ef (antisite), the bridging oxygen atom of the NiNi pair is replaced with an anion dopant X (X = F, Cl, or S). Indeed, Figure 6b indicates that both Cl and F favor Ni ions in opposite spin states, but S favors the parallel spin state. Table 1 has shown that the charge state of S is 1e−, due to its low electronegativity. This indicates that the pz orbital of S is not fully filled, thus impairing the superexchange interaction with neighboring dz2 orbitals and leading to higher antisite defect formation energy in S-doped LiNiO2. Figure 6b also reveals that F greatly reduces the formation energy of Li/Ni antisite defects by 0.40 eV, while Cl slightly increases the energy by 0.05 eV. This result agrees well with the experimental observation that, by increasing F doping concentration, the Li/Ni exchange ratio of after-synthesized LiNiO2 and LiNi0.8Co0.15Al0.05O2 is also increased.20,21 The analysis of the electronic structure shows that anion doping effects on Li/Ni exchange seem to be related to the energy states of dopant electrons. Figure 8 shows the density of states (DOS) of the pz electrons of the bridging anion in both pristine and doped LiNiO2. By comparing the DOS features of anion dopants with oxygen in pristine LiNiO2, one can contend

Figure 6. (a) Schematic illustration of a Li/Ni antisite defect and a NiNi pair in doped LiNiO2. (b) Li/Ni antisite defect formation energy in pristine and anion doped LiNiO2, when the Ni ions in the NiNi pair have parallel and opposite spin.

The formation energy of a Li/Ni antisite defect is calculated according to the following formula: Ef (anti‐site) = E(defect) − E(perfect)

in which E(defect) and E(perfect) denote the total energy of LiNiO2 with and without the Li/Ni antisite defect, respectively. 6946


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The results of this section can provide useful insights for many experimentally observed anion doping effects on LiNiO2 rate performance. F promotes fast Li ion migration over the whole structure, thus contributing to the observed enhanced rate performance. However, since it largely reduces the formation energy of Li/Ni antisite defects, more blocking Ni ions in the Li layer are expected upon increased F concentration, and the subsequent lowering of the rate performance. These two conflicting effects of F doping help to understand several experimental observations: Woo et al. performed EIS measurements on pristine and F-doped LiNi0.8Co0.1Mn0.1O2 with different doping concentrations.23 The electrode particles showed decreased charge-transfer resistance from 141 Ω to 60 Ω when F concentration increases from 1% to 2%. However, when the concentration further increases to 3%, the resistance goes back to 132 Ω again. We expect that, upon increased F doping concentration, the negative effect of having more Li/Ni antisite defects finally overwhelms the positive effect of facilitating Li migration. As mentioned before, X. Li et al. also reported an unusual F doping effect in LiNi0.8Co0.15Al0.05O2: F-doping decreased the capacity with current rate ranging from 0.1C to 2C, but increased the capacity at the high rate of 5C.21 Their EDS analysis showed that most of the F ions are distributed in the surface region, which may also be more abundant in Li/Ni antisite defects, based on the present results. Therefore, the surface acts as a “blocking” layer for Li diffusion into the electrolyte, leading to the reduced capacity. However, at a higher current rate of 5C, Li migration energy barrier might become the determining factor for rate performance, thus the F effect of facilitating Li migration increases the high rate capacity. Although no experimental results about Cl- and Sdoping effects on rate performance have been reported, our results suggest that they could still enhance the rate performance, owing to their effect of expanding the Li migration spacing and inhibiting the formation of Li/Ni antisite defects (although this effect is not very remarkable). However, the doping concentration should still be optimized,

Figure 8. Density of states (DOS) of pz electrons of (a) oxygen in pristine LiNiO2, (b) Chlorine in Cl-doped LiNiO2, (c) sulfur in Sdoped LiNiO2, and (d) fluorine in F-doped LiNiO2. The main DOS contribution regions of the anions are highlighted with yellow shades.

that the energy state of the Cl electron is similar to that of O, which is consistent with the similar Li/Ni antisite defect formation energy between pristine LiNiO2 and Cl-doped LiNiO2. However, the DOS of the S atom is shifted to a larger energy state, indicating a higher energy level of the Ni Ni pair. For F doping, Figure 8 shows that the DOS of F is mainly distributed in the lower energy range between −7 and −4 eV. This can provide a reasonable interpretation of the low antisite defect formation energy of 0.12 eV shown in Figure 6b.

Figure 9. (a) Comparison between DFT+U, DFT+U+D3, and experimental data of the obtained lattice parameter c. Evolution of lattice constants (b) a, (c) c, and (d) volume of pristine and anion-doped LiNiO2 during delithiation. (a: ref 38, b: ref 39). 6947


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Figure 10. Configurations of the interlayer Ni migration in a position (a) neighboring to a dopant X, and (c) distant to a dopant X. Parts (b) and (d) plot the corresponding energy barriers for Ni migration, neighboring and distant to the dopant, respectively. The yellow spheres correspond to the migrating Ni ion.

measurements, Figure 9a shows the necessity of including vdW dispersion corrections to calculate the lattice parameters of highly delithiated LiNiO2 accurately. By doing that, the difference on the obtained lattice parameter c and the available experimental data is within 2%. Upon delithiation, a decreases continuously from 2.89 to2.77 Å, but c reveals a turning point: it increases gradually from x = 0 to x = 0.75, but suddenly decreases when x is above 0.75. Consequently, the cell volume is almost constant during delithiation between 0 < x < 0.75, but a sharp shrinkage is observed for 0.75 < x < 1.00. This sharp change is the result of the transformation of larger Ni3+ octahedron (with a volume of 10.14 Å3) to the shrunken Ni4+ octahedron (with a volume of 8.49 Å3). These results agree with the experimental evidence, in which a discontinuous interlayer distance drop at the end of charging was observed.38,39 Within the LiNiO2 powders, a sharp lattice distortion will introduce high strain among primary particles with different crystal orientations, which could lead to the formation of micro cracks.51 Figure 9 also shows that the effects of anion doping on the lattice constant a are almost negligible, but more noteworthy on the lattice constant c and the cell volume. Among the considered dopants, both S and Cl show stronger effects than F in enlarging the structural parameters, owing to their larger atomic radius. Such expanding effect on the interlayer distance could enlarge Li ion migration space, lowering the migration energy barrier,42 in accordance with previously reported results on ionic conductivity. Additionally, Cl seems to alleviate the sharp lattice constant c and volume change at the final charging stage. Such c shrinkage is reduced from 1.08 to0.73 Å with Cl doping. Note that this result indicates that Cl cannot totally suppress LiNiO2 lattice distortion, although it still might improve the cycling retention by partially inhibiting micro crack formation.

since a high concentration might amplify their blocking effect on the diffusion of neighboring Li ions. Additionally, initial capacity loss has been widely observed in both F- and S-doped oxides, in an amount proportional to the doping concentration.21,22,24 On the basis of our results, doping-induced capacity losses might originate from different mechanisms: F increases the Li/Ni exchange ratio, thus occupies active Li sites and blocks Li ion migration, whereas S increases the migration energy barrier of neighboring Li ions. 3.3. Structural Stability. LiNiO2 shows poor structural stability during cycling, especially at high operating temperatures.48−50 On the basis of the available experimental and theoretical reports, the problems are intimately related with cycling induced secondary particle micro cracks and phase transformations into the NiO rock-salt structure. The former is typically ascribed to the sharp lattice expansion/shrinkage during the delithiation/lithiation cycling,51−53 and the latter involves the occupation of Li sites by Ni ions through interlayer Ni migration at the delithiated state. This phase transition from layered structure (R3̅m) to spinel (Fd3̅m) and subsequently to rock-salt (Fm3̅m) was reported to happen at a low temperature of 200 °C.54,55 Therefore, to evaluate the doping effect on the structural stability, the doping effects on both lattice distortion and interlayer Ni migration during delithiation must be carefully examined. 3.3.1. Lattice Distortion. Figure 9 shows the lattice parameters as a function of the delithiation ratio x in pristine and anion doped Li1−xNiO2. As stated in previous sections, the delithiation of Li1−xNiO2 can be characterized by four different stages, separated by x = 0.25, 0.5, and 0.75, in terms of changes in lattice parameters, interlayer distance, and open circuit voltage.38 The obtained lattice constants a and c of pristine LiNiO2 are 2.92 and 14.29 Å, respectively, in agreement with the experimental values of 2.89−2.90 Å and 14.18−14.24 Å.38,56 By comparing DFT+U, DFT+U+D3, and experimental 6948


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Figure 11. Averaged charge density (e·Å−3) on the (210) lattice plane for the final state of interlayer Ni migration in (a) pristine LiNiO2, (b) S doped LiNiO2, (c) F doped LiNiO2 and (d) Cl doped LiNiO2. (X axis: [120]; Y axis: [001]).

3.3.2. Ni Migration. To evaluate the doping effect on the interlayer Ni migration during charging, a Li11Ni18O34X1 (X = O, F, Cl, or S) supercell in which ∼40% Li ions have been extracted was used, thus allowing sufficient space for a Ni ion in a Oh site of the Ni layer to migrate to the neighboring Td site of the Li layer. An oxygen vacancy is also located at the migration path, which is typically considered in order to assist TM migration in layered structures.57−59 The Ni migration can be regarded as the initial step of a phase transition from the layered to the spinel and the rock-salt structures.57 The two models considered, corresponding to a migrating Ni ion neighboring and distant to the dopant, and the obtained Ni migration energy barriers are shown in Figure 10. Interlayer Ni migration in pristine LiNiO2 shows a barrier of ∼0.78 eV, with a structurally stable final state, indicating the thermodynamic stability of a Ni ion located at the Td sites of the Li layer. Figure 10a,b implies that the three anion dopants considered show a possible blocking effect on the migration of neighboring ions. Cl and S increase the migration energy barrier by 0.2 and 0.6 eV, respectively, whereas F decreases the barrier by 0.05 eV. As discussed in the previous section, owing to the atomic radius sequence of S > Cl > O > F, S and Cl might squeeze the migration path, while the smaller anion F enlarges it. Besides the effect on the migration energy barriers, the three dopants increase the energy of the final state by 0.65 eV, thus indicating an endothermic character of the reaction leading the Ni ions to the Li layer. Furthermore, since the reverse energy barriers are as low as 0.3−0.6 eV, the migrated Ni ion might show large probabilities to hop back to the initial Oh site. To examine the electronic origin of the endothermic character of the Ni migration induced by the dopants, the charge density on the (210) lattice plane is plotted in Figure 11. It shows the electron density distribution around the Ni-dopant bonds at the final state of the interlayer Ni migration (Image #4 in Figure 10b). For pristine LiNiO2, the lengths of the four

NiO bonds at the Td site are 1.80 Å. The significant charge density overlap between Ni cation and O anion indicates a robust NiO bonding interaction. However, the lengths of Nidopant bonds in doped LiNiO2 increase to 1.98−2.21 Å, which combined with the reduced charge density overlap between Ni and the corresponding dopant anion shown in Figure 11b−d, imply a weaker Ni-dopant interaction. Therefore, extrinsic anion doping does not facilitate the Ni occupation of Td sites in the Li layer, thus impairing Ni migration-induced phase transitions. Ni migration to Td sites distant to the corresponding dopant (∼5 Å in the present model, see Figure 10c) shows opposite results. As can be seen in Figure 10d, Cl and S lower the energy barrier to 0.37 and 0.08 eV, respectively. Both Figures 5 and 9 prove that Cl and S remarkably expand the interlayer distance and the cell volume. Since the anion dopant is distant to the migrating Ni ion, the dopant-induced narrowed migration path does not affect Ni diffusion, and the dopant-induced enlarged interlayer space effectively facilitates Ni migration. Figure 10d also shows that F-doping has a negligible effect on the migration energy barrier, something expected because of its small effect on the structural lattice parameters. Then, in terms of the anion doping effects on both lattice distortion and Ni migration, F seems to be the best choice in order to maintain the structural stability of LiNiO2-based oxides: it can suppress interlayer Ni migration and will not deteriorate lattice distortion. This is the obvious reason why F is the most widely reported anion dopant employed to alleviate the structural instability. For instance, the decomposition temperature of LiNi0.8Co0.1Mn0.1O2 was increased from 190 to 245 °C with a F doping concentration of 3%,23 and many Fdoped LiNiO2-based oxides show an improved cycling stability.20−23 Finally, a priori Cl would also be a good candidate to improve the structural stability, given its effect in suppressing lattice distortion and impairing neighboring Ni 6949


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Chemistry of Materials Table 2. Summary of Anion Doping Effects on the Battery Performancea

a

“L” and “H” represent low doping concentration and high doping concentration, respectively.

wider scope of possibilities to optimize the overall electrochemical performance. However, due to possible dopant− dopant interactions, a realistic analysis would be substantially more complex and it is beyond the scope of the present work. But some useful guiding rules for designing codoping strategies can still be extracted. For instance, cation dopants with the capability to suppress Li/Ni exchange, such as Co,45 might be likely candidates to form a codoping pair with F. Similarly, cation dopants that can improve the redox potential or the structural stability could be coupled with Cl or S dopants. We hope that this work will stimulate further research along that direction. Moreover, although the present work is based on the LiNiO2 model system, which can represent a wide range of Ni-rich layered oxides, some of the conclusions can be extended to other class of oxide materials. Taking the spinel Li2FeMn3O8 as an example, Cl doping improved the rate performance by expanding the interlayer distance. However, for large Cl concentrations, the effect is the opposite, and the rate capacity is lowered.62 Following the conclusions of this work, localized blocking effect of Cl atoms on the migration of neighboring Li ions might indicate conflicting effects of Cl doping in spinel cathode materials. Therefore, it would be useful to examine similar conflicting roles of anion doping in other classes of oxide electrode materials.

migration. However, similar to S, it also expands the cell volume and facilitates Ni migration to Td sites distant to the dopants. 3.4. Proposed Doping Strategy. On the basis of the results reported so far, the three anion dopants considered seem to enhance certain properties while worsen others, also showing distance-dependent effects on the rate performance and the structural stability. These effects have been summarized in Table 2, based on which we propose several doping strategies that can take advantage of the positive doping roles and suppress undesirable side effects. 3.4.1. Optimized F-Doping. Table 2 shows that F could be the ideal candidate to improve almost all the considered properties. However, it facilitates the formation of Li/Ni antisite defects, which might impede fast Li ion diffusion. Therefore, suppressing Li/Ni exchange in F-doped LiNiO2 would be the most critical concern. Ion exchange is a very promising approach, in which a well-layered Na compound is synthesized and latter Na+ ions are replaced with Li+. It has been successfully applied to synthesize layered LiNi0.5Mn0.5O2 with low exchange ratio.42 3.4.2. Optimized Cl- and S-Doping. Cl and S dopants show similar effects on the properties of LiNiO2. They suppress the formation of Li/Ni antisite defects and the lattice distortion (especially Cl), but also reveal conflicting roles on the structural stability and rate performance, depending on the dopant distance to the migrating cation. Consequently, a low doping concentration can benefit the rate performance whereas a high concentration benefits the structural stability. Therefore, solely tuning the doping concentration would be insufficient to optimize the overall performance. It has been reported that concentration gradient or core−shell particles with resistant films to oxygen evolution on the surface can produce highly stable Ni-rich layered oxides.60,61 However, these methods cannot impede Li/Ni exchange and crack formation. Then, to best optimize the overall performance, they should be coupled with a low concentration of Cl or S doping, in order to improve the rate performance without lowering the voltage. Additionally, benefiting from the large family of candidates for both cation and anion doping, codoping could provide a

4. CONCLUSIONS In this work, we have examined in detail the effects of three common anion dopants (F, Cl, and S) on the electrochemical performance of LiNiO2 cathode materials. Our results unveil the conflicting nature of the roles played by anion dopants: they enhance some properties while worsen others, also showing distance-dependent features (especially Cl and S). The results agree well with available experimental observations and provide the necessary understanding at the atomic scale. On the basis of the obtained data, we have proposed several promising directions that might help to optimize the anion doping strategy. We expect that this work can inspire further anion doping studies and promote related research and materials doping design for other relevant cathode materials. 6950


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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.6b02627. Bader charge analysis in F doped LiNiO2, Ni migration in Li layer, Li migration neighboring to Li/Ni antisite defect, Saddle point structure during Li migration in F doped LiNiO2, Charge distribution of Ni ions neighboring to Li/Ni antisite defect (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (S.-G.D.). *E-mail: [email protected] (K.J.C.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by Samsung GRO Project. The authors also acknowledge the Texas Advanced Computing Center (TACC) for providing computational resources.

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REFERENCES

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