Lithium Deficiencies Engineering in Li-Rich Layered Oxide Li1

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Cite This: J. Am. Chem. Soc. 2019, 141, 10876−10882

Lithium Deficiencies Engineering in Li-Rich Layered Oxide Li1.098Mn0.533Ni0.113Co0.138O2 for High-Stability Cathode Pengfei Liu,† Hong Zhang,‡ Wei He,† Tengfei Xiong,§ Yong Cheng,† Qingshui Xie,*,† Yating Ma,† Hongfei Zheng,† Laisen Wang,† Zi-Zhong Zhu,∥ Yong Peng,‡ Liqiang Mai,*,§ and Dong-Liang Peng*,†

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State Key Lab of Physical Chemistry of Solid Surface, Collaborative Innovation Center of Chemistry for Energy Materials, College of Materials, Xiamen University, Xiamen 361005, China ‡ Key Laboratory of Magnetism and Magnetic Materials of the Ministry of Education, School of Physical Science and Technology and Electron Microscopy Centre of Lanzhou University, Lanzhou University, Lanzhou 730000, China § State Key Laboratory of Advanced Technology for Materials, Synthesis and Processing, International School of Materials Science and Engineering, Wuhan University of Technology, Wuhan 430070, China ∥ Collaborative Innovation Centre for Optoelectronic Semiconductors and Efficient Devices, Department of Physics, Xiamen University, Xiamen 361005, China S Supporting Information *

ABSTRACT: Li-rich layered oxides have been in focus because of their high specific capacity. However, they usually suffer from poor kinetics, severe voltage decay, and capacity fading. Herein, a long-neglected Li-deficient method is demonstrated to address these problems by simply reducing the lithium content. Appropriate lithium vacancies can improve dynamics features and induce in situ surface spinel coating and nickel doping in the bulk. Therefore, the elaborately designed Li1.098Mn0.533Ni0.113Co0.138O2 cathode possesses improved initial Coulombic efficiency, excellent rate capability, largely suppressed voltage decay, and outstanding long-term cycling stability. Specifically, it shows a superior capacity retention of 93.1% after 500 cycles at 1 C (250 mA g−1) with respect to the initial discharge capacity (193.9 mA h g−1), and the average voltage still exceeds 3.1 V. In addition, the discharge capacity at 10 C can be as high as 132.9 mA h g−1. More importantly, a Li-deficient cathode can also serve as a prototype for further performance enhancement, as there are plenty of vacancies.

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protection layer to suppress surface lattice oxygen release21−24 and ameliorate the aforementioned issues.8,25 Fortunately, lithium deficiencies can induce the formation of spinel opportunely.2 In addition, nickel ions tend to migrate to lithium slabs under the lithium-deficient conditions because the energy barrier of nickel ions is the lowest among all three transition-metal (TM, Mn, Co, Ni) ions in LLOs and is comparable to that of Li+.26 Hence, nickel ions in lithium slabs can function as doping elements to improve the stability of the crystal structure.11−14,27 Nevertheless, the in situ nickel doping may hinder the migration of Li+ and lead to a poor rate capability.28 But fortunately, the impediment of nickel ions in the migration of Li+ can be compensated by the intentionally preintroduced lithium vacancies. Therefore, the idea that improvement on kinetics by introducing lithium vacancies has a potential to come true. Inspired by the above discussions, a blueprint came to our mind naturally: lithium-deficient lithium-rich layered oxide

INTRODUCTION Lithium-rich layered oxides (LLOs) are among the most promising cathode materials because of their low cost, high voltage, and high specific capacity.1,2 A consensus has been reached that excess lithium ions are a prerequisite for the high specific capacity to take advantage of the extra anionic redox in LLOs.3 However, excess lithium ions in the lattice are kinetically unfavorable because they can prohibit the transformation of Li+ from the hard oxygen dumbbell hop (ODH) to the easy tetrahedral site hop (TSH).4 Notably, the difference in specific capacity can be up to 50 mA h g−1 for the same chemical formula.5,6 So we naturally ask, is it possible to improve the kinetic performance of LLOs and the utilization of lithium by introducing proper lithium vacancies in the premise of no sacrifice in energy density? However, higher utilization of lithium will induce a deeper delithiated state of LLOs, resulting in capacity fading, voltage decay, and poor kinetics.7−10 Various effective strategies have been developed to address the above problems, such as doping11−15 and surface coating.16−20 Spinel phase is one of the most effective coating materials, as it can function as a © 2019 American Chemical Society

Received: May 9, 2019 Published: June 15, 2019 10876

DOI: 10.1021/jacs.9b04974 J. Am. Chem. Soc. 2019, 141, 10876−10882

Article

Journal of the American Chemical Society

Figure 1. (a) Cross section of LLLO. (b) Lithium vacancies. (c) In situ spinel Li4Mn5O12 coating. (d) In situ nickel ion doping. (e) Schematic illustration of the synthesis of LLLO.

(LLLO) where the surface is protected by the spinel phase and the bulk is layer structured. Furthermore, there are a moderate number of lithium vacancies and nickel ions in lithium slabs, as illustrated in Figure 1a−d. Lithium vacancies are preferable on account of a larger diffusion coefficient and lower cost.29 In addition, nickel doping can stabilize the crystal structure, and in situ spinel coating due to lithium deficiencies can protect cathode materials from the erosion of electrolytes and suppress the release of surface lattice oxygen.7,30,31 Under this guidance, LLLOs accounting for only 83.5 at. % of lithium in pristine lithium-rich layered oxide (PLLO) were synthesized as shown in Figure 1e. As expected, high rate capability, suppressed voltage decay, and outstanding long-term cycling stability are realized simultaneously for the LLLO cathode. More importantly, this method is simple, low-cost, repeatable, and very suitable for massive production.

Table 1. Element Ratios in PLLO and LLLO from ICP-MS Tests element PLLO LLLO

Li

Mn

Ni

Co

Li/TM

1.309 1.098

0.533 0.533

0.112 0.113

0.136 0.138

1.676 1.400

which is in agreement with density functional theory (DFT) calculations (Figure 2b, Figure S1, and Table S1).32 Bader charge analysis shows that the charge redistribution mainly happens in TM ions and oxygen ions around the lithium vacancy (Figure S2, Table S2). Changes in TM ions are due to the unclosed shell, and those in lattice oxygen originate from the superexchange interaction in a unique Li−O−Li configuration.33,34 Thus, the c-parameter increases.24 In addition, the preformed lithium vacancies are similar to those induced by the extraction of some lithium ions in the initial stage of charging. Therefore, the vacancies will lead to the oxidization of Ni2+ to Ni3+ first due to its highest electrochemical potential among all the TM ions, while they almost have no influence on the valence states of manganese and cobalt. This was further verified by X-ray photoelectron spectroscopy spectra (XPS, Figure S3). The morphologies and microstructures of PLLO and LLLO were investigated by scanning electron microscopy (SEM), high-resolution transmission electron microscopy (HR-TEM), and scanning transmission electron microscopy (STEM). Element distributions of both of the samples are uniform (Figure S4), and the secondary particle is about 6−7 μm

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RESULTS AND DISCUSSION Sample Synthesis and Structure Characterization. LLLO was synthesized by simply adjusting the ratio between lithium and TM elements. The chemical formula of PLLO is Li 1.309 Mn 0.533 Ni 0.112 Co 0.136 O 2 , while that of LLLO is Li1.098Mn0.533Ni0.113Co0.138O2 by inductively coupled plasma mass spectrometry (ICP-MS, Table 1). Both of them have a typical layered structure (Figure 2a). However, LLLO has extra satellite peaks marked by “*”, which belongs to spinel phase Li4Mn5O12 (Fd3̅m, Figure 2a insets). The slight decrease of a and the increase of c may be caused by lithium vacancies, because the shielding between oxygen slabs will be weakened when some lithium ions are extracted from the lithium slabs, 10877

DOI: 10.1021/jacs.9b04974 J. Am. Chem. Soc. 2019, 141, 10876−10882

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Journal of the American Chemical Society

Figure 2. (a) XRD patterns of PLLO and LLLO. Insets are the enlarged patterns. (b) Rietveld refinement results of LLLO. SEM (c) and HR-TEM images (d) of PLLO. SEM (e), HR-TEM (f, g), and HAADF (h) images of LLLO. Red arrows in (h) are reflections of lithium slabs partly occupied by nickel ions.

concluded that some TM ions have occupied the lithium slabs, in consideration of the fact that nickel ions have the best mobility among all the TM ions in both intra- and interlayers in LLOs.26 In addition, the energy barrier of nickel ions is comparable to that of lithium ions. Therefore, we believe that nickel ions have occupied the lithium slabs. It is further verified by the ratios of (003)/104, which are 1.66 and 1.98 for LLLO and PLLO, respectively. This means that LLLO has a higher degree of Li/Ni mixing and more nickel ions occupy the lithium slabs, which can be regarded as an indirect evidence of lithium vacancies. Thus, the blueprint has been realized. Electrochemical Properties. Electrochemical performances of the two samples were evaluated in coin cells using Li as counter electrode. They were activated at 0.1 C (2.0−4.6 V) first and then used for other tests in the voltage range of 2.0− 4.8 V. The activating charge and discharge specific capacities of LLLO are 295.3 and 247.9 mA h g−1 with a Coulombic efficiency (CE) of 83.9% (Figure 3a). Due to the unsuppressed activity of lattice oxygen and poor kinetics, PLLO shows a higher charge capacity yet a lower discharge capacity. Thus, the utilization ratio of lithium in LLLO is improved greatly in consideration of the fact that lithium in LLLO only accounts for 83.5 at. % of that in PLLO. In addition, the spinel-related characteristic plateau is observed in the discharge profile of

(Figure 3c,e). However, the primary particle size of PLLO is about 200−500 nm, and that of LLLO is about 50−150 nm. The decreased primary particle size in LLLO leads to an increase of the specific surface area to 6.33 m2 g−1, which is larger than 1.50 m2 g−1 in PLLO (Figure S5). However, the influence of lithium vacancies on tap density is minor, as those of PLLO and LLLO are 1.75 and 1.66 g cm−3, respectively. The d-spacings of PLLO in the upper and lower insets of Figure 3d are 0.472 and 0.427 nm, belonging to the (003) planes of R3̅m and the (020) planes of C2/m structures, respectively. For LLLO, the d-spacing of 0.476 nm corresponds to (003) planes of the layered structure (Figure 3f). Meanwhile, fast Fourier transformation (FFT) of the selected area in red dashed lines in Figure S6 identifies the formation of spinel phase Li4Mn5O12. The d-spacing of 0.470 nm in Figure 3g corresponds to the (111) planes of the spinel phase, which can further be verified by the high-angle annular dark field (HAADF) image in Figure S7 and blue shift of the Raman spectra (Figure S8) because of the shortened bonds and internal stresses.8 The HAADF image of the bulk LLLO shows a typical layered structure (Figure 2h). Interestingly, a row of dark but visible dots indicated by red arrows is abnormal and noteworthy because they locate in lithium slabs which should be invisible in the HAADF image. Therefore, it can be 10878

DOI: 10.1021/jacs.9b04974 J. Am. Chem. Soc. 2019, 141, 10876−10882

Article

Journal of the American Chemical Society

Figure 3. (a) Activating charge−discharge curves at 0.1 C (1 C = 250 mA g−1). (b) CV plots of LLLO at a scan rate of 0.05 mV s−1. (c) Rate performances. (d) Cycling performances at 2 C. (e) Stabilities of specific capacity, voltage, and energy density of LLLO at 1 C. All cells were first activated at 0.1 C.

Figure 4. (a, b) In situ peak (003) and (101) evolution of PLLO (a) and LLLO (b) and the corresponding charge−discharge curves at 0.5 C (2.0− 4.8 V), where the value of intensity changes from 0 (red) to 900 (blue). (c) Profiles of Zr vs ω−1/2 of fresh half-cells. (d) Top view of the diffusion path. (e, f) Influence of a lithium vacancy (e) and nickel ion doping (f) on the diffusion energy barrier of Li+. Insets: Diagram of the diffusion path.

LLLO around 2.6 V24,30 and further verified in cyclic voltammetry tests (CV, Figure 3b). However, PLLO does not have spinel-like peaks (Figure S9). The rate performance

of LLLO is also much better than that of PLLO (Figure 3c, Figure S10). It has a discharge capacity of 132.9 mA h g−1 even at 10 C. When cycled at 0.2 C, LLLO delivers a discharge 10879

DOI: 10.1021/jacs.9b04974 J. Am. Chem. Soc. 2019, 141, 10876−10882

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Journal of the American Chemical Society capacity of 241.9 mA h g−1 in the 80th cycle, and PLLO has a discharge capacity of only 218.5 mA h g−1, confirming the higher lithium utilization in LLLO (Figure S11). Cyclic lifespan at high current density is used for further investigation on the influence of lithium vacancies. The first discharge capacity of LLLO at 2 C is 194.1 mA h g−1 and only declines to 192.3 mA h g−1 after 300 cycles with a capacity retention of 99.1% (Figure 3d). However, PLLO has a discharge capacity of only 141.8 mA h g−1 after the same cycles. When cycled at 1 C, the initial discharge capacity of LLLO is 193.9 mA h g−1 and the energy density is 663.6 Wh kg−1 (Figure 3e). Surprisingly, it retains a discharge capacity of 180.6 mA h g−1 at the 500th cycle with a capacity retention of 93.1%. However, the initial discharge capacity of PLLO is only 180.0 mA h g−1 and declines to 142.7 mA h g−1 only after 400 cycles (Figures S12, S13). dQ/dV plots were used to study the electrochemical reactions of PLLO and LLLO during longterm cycling at 1 C (Figure S14). Overlap of dQ/dV curves among different cycles for PLLO is bad, indicating its poor electrochemical reaction reversibility and poor capacity retention. On the contrary, overlap of dQ/dV curves among different cycles for LLLO is good, indicating its obviously enhanced capacity and voltage stability. Therefore, LLLO has an average voltage above 3.1 V in the entire electrochemical process, and the energy density of LLLO is still 560.6 Wh kg−1 even after 500 cycles with a retention of 84.5% (Figures 3e, S12). In addition, the first CEs of the LLLO cathode at 1 and 2 C are both about 83% (Figures 3d, S15). Then they are almost 100% in the following cycles. A similar trend is also observed for PLLO except that CEs in the first cycles at 1 and 2 C are both below 80%. Furthermore, CEs of LLLO during long-term cycling are more stable than those of PLLO. In conclusion, the designed LLLO cathode material delivers a higher and more stable Coulombic efficiency than PLLO, which is beneficial for its practical application. The above improvements may be caused by a combination of lithium vacancies in the bulk, spinel coating on the surface, and nickel doping in lithium slabs. Structure Evolution and Mechanism. In situ XRD measurements in the first three cycles at 0.5 C were conducted to verify the enhanced structural stability and the improved utilization of lithium of LLLO. Although similar behaviors are observed with previous reports in Figure 4a,b,8,35,36 an abnormal phenomenon contradicting our common sense is worthy of attention. That is, the peak shift of LLLO is larger than PLLO especially at the end of the charging process, which disagrees with our conventional understanding that stable electrode materials mean few changes in lattice parameters during cycling.8,36 Considering the fact that LLLO has much better electrochemical performances than PLLO and the lithium content accounts for only 83.5 at. % of that in PLLO, the utilization of lithium of LLLO is much higher. This will lead to a higher delithiated state and a higher oxidation state of lattice oxygen as shown in the analysis of Bader charge (Table S2).3,34 Therefore, the decrease of the c-parameter is more serious. Due to the in situ surface spinel Li4Mn5O12 coating and the in situ nickel doping in LLLO, the release of surface lattice oxygen is suppressed and the stability of the whole structure can be maintained at the same time in spite of a higher utilization of lithium. In order to figure out the enhancement mechanism of lithium vacancies on kinetics, electrochemical impedance spectroscopy (EIS) investigations were carried out (Note S1,

Figure S16, and Table S3). The Warburg factor was calculated using the linear fitting of the Zr vs ω−1/2 profiles from 0.06 to 0.01 Hz (Figure 4c). Accordingly, the diffusion coefficients of Li+ in LLLO and PLLO are 7.81 × 10−18 and 1.14 × 10−18 cm2 s−1, respectively. The Warburg factor and lithium concentration are closely related to the diffusion energy barrier of Li+ according to eq S1. Therefore, DFT calculations were performed to figure out the effect of lithium vacancies and nickel doping (Note S2). The end points of the diffusion path are in purple, the lithium vacant site is in black, and the nickel ion is in green (Figure 4d). In addition, the Li+ in black is taken away to create a lithium vacancy for LLLO. The local lithium-vacant arrangement around the end points of the hopping path are closely related to the two qualitatively different migration paths for Li+. One is the hard ODH when the two adjacent sites of the end points are both occupied by Li+. The other is the easy TSH when there is at least one lithium vacancy in the two adjacent sites of the end points.4 In the ODH path, Li+ migrates directly along the shortest line between the origin and the end point of the hopping path (Figure 4e). The energy differences between the end points are both 75.3 meV for PLLO and LLLO. Choosing site 1 as the reference point, the energy barrier of PLLO is 447.9 meV, whereas the energy barrier of LLLO drops to 422.1 meV when there is a lithium vacancy in the next nearest site. More critically, Li+ in LLLO will migrate in the TSH mode in advance during the charging process because of the preintroduced lithium vacancies. Moreover, the influence of nickel doping on the migration of Li+ is shown in Figure 4f. For LLLO, the energy difference between the diffused sites is 53.8 meV and the energy barrier is 510.3 meV. For PLLO, the energy difference is 44.6 meV, and the energy barrier is 523.5 meV. So nickel doping will enlarge the energy difference between diffused sites and reduce the energy barrier slightly. Furthermore, the improvements of the lithium-deficient strategy on the kinetics can be maintained during cycling, as indicated by the EIS after charging−discharging at 2 C for 50 cycles in Figure S17. Furthermore, morphologies and crystal structures of PLLO and LLLO after cycling for 300 times at 2 C were carefully investigated. As shown in Figure 5a,b, there are obvious crystal lattice distortions and plenty of stack faults in the HR-TEM image of PLLO. However, the crystal lattice of LLLO is integrated. The morphological maintenance of secondary particles for LLLOs is also better than that of PLLO (Figure S18). The good performance of LLLO may originate from the synergistic effect of the in situ surface spinel coating and in situ nickel doping. Nickel ions in lithium slabs can form strong Ni− O bonds, prevent the migration of manganese ions from transition metal slabs to adjacent lithium slabs, and prevent the collapse of the crystal structure especially at the high delithiated state.37 The stabilization of LLLO can be further strengthened with the addition of the protection from surface spinel coating. On the basis of the above discussion, ingeniously introduced multiple effects by lithium deficiencies on LLO during the whole electrochemical process are schematically illustrated in Figure 5c. The preintroduced lithium vacancies can promote the migration of Li+, the asformed in situ spinel coating can suppress the release of lattice oxygen and prevent the corrosion from the electrolyte, and the as-formed in situ nickel doping can stabilize the crystal structure in the deep-delithiated state. Therefore, LLLO has excellent electrochemical performances including improved 10880

DOI: 10.1021/jacs.9b04974 J. Am. Chem. Soc. 2019, 141, 10876−10882

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cycling performance of PLLO at 1 C, dQ/dV plots, CEs at 1C, EIS fitting model, EIS after cycling, SEM images after cycling, changes in lattice parameters of PLLO and LLLO, Bader charge analysis, and EIS fitting results (PDF)

AUTHOR INFORMATION

Corresponding Authors

*[email protected] (Q. Xie) *[email protected] (L. Mai) *[email protected] (D. Peng) ORCID

Qingshui Xie: 0000-0003-2105-6962 Laisen Wang: 0000-0001-9531-4480 Zi-Zhong Zhu: 0000-0001-5353-4418 Liqiang Mai: 0000-0003-4259-7725 Dong-Liang Peng: 0000-0003-4155-4766 Notes

Figure 5. (a, b) HR-TEM images of PLLO (a) and LLLO (b) after cycling for 300 times at 2 C. (c) Schematic illustration of the charging−discharging process of LLLO.

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This study was financially supported by the National Key R&D Program of China (No. 2016YFA0202600), the National Natural Science Foundation of China (Nos. 51701169 and 51871188), the Natural Science Foundation of Fujian Province of China (No. 2017J05087), the Key Projects of Youth Natural Foundation for the Universities of Fujian Province of China (No. JZ160397), and the “Double-First Class” Foundation of Materials and Intelligent Manufacturing Discipline of Xiamen University. The authors thank the Electron Microscopy Centre of Lanzhou University for the microscopy and microanalysis of our specimens.

CEs, fast rate capability, extraordinary voltage stabilization, and long lifespan.

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CONCLUSION In conclusion, lithium-deficiency-induced Li1.098Mn0.533Ni0.113Co0.138O2, accounting for only 83.5 at. % of lithium in pristine lithium-rich layered oxide, has been successfully fabricated using a facile and effective strategy. The lithium deficiencies would induce an in situ surface spinel phase coating and in situ nickel doping in lithium slabs. Preintroduced lithium vacancies could reduce the energy barrier of Li+ in the initial ODH diffusion path, accelerate the transformation from the hard ODH to the easy TSH, and effectively improve the utilization of lithium as verified by the combination of in situ XRD patterns and DFT calculations. Moreover, in situ surface spinel coating could suppress the release of lattice oxygen, and in situ nickel doping could stabilize the whole crystal structure. Thus, LLLO can cycle for over 500 cycles stably with a high capacity retention of 93.1%, and the average voltage is still over 3.1 V. In addition, LLLO shows an excellent rate capability with a discharge capacity of 132.9 mA h g−1 even at 10 C. Also, it can be used as a prototype for further modification, as there are plenty of host sites. Finally, the Li-deficient method has the potential to promote massive production of LLOs because of its simple, low-cost, and repeatable process. We believe that this design philosophy can be generalized to other layered oxide cathodes to boost their overall performance economically and succinctly.

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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/jacs.9b04974. Experimental Section, calculation details, Rietveld refinement of PLLO, calculation model, XPS, EDS mapping, N2 adsorption−desorption isotherms, FFT patterns of spinel phase, HAADF image on the edge of LLLO, Raman spectra, CV curves of PLLO, rate tests of LLOs with different contents of lithium vacancies, cycling performances at 0.2C, charge−discharge profiles at 1 C, 10881

DOI: 10.1021/jacs.9b04974 J. Am. Chem. Soc. 2019, 141, 10876−10882

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DOI: 10.1021/jacs.9b04974 J. Am. Chem. Soc. 2019, 141, 10876−10882