A Guideline for Tailoring Lattice Oxygen Activity in Lithium-Rich

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A Guideline for Tailoring Lattice Oxygen Activity in Lithium-Rich Layered Cathodes by Strain Pengfei Liu, Wei He, Qingshui Xie, Yong Cheng, Wanjie Xu, Zhensong Qiao, Laisen Wang, Baihua Qu, Zi-Zhong Zhu, and Dong-Liang Peng J. Phys. Chem. Lett., Just Accepted Manuscript • DOI: 10.1021/acs.jpclett.9b00419 • Publication Date (Web): 01 Apr 2019 Downloaded from http://pubs.acs.org on April 9, 2019

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The Journal of Physical Chemistry Letters

A Guideline for Tailoring Lattice Oxygen Activity in Lithium-Rich Layered Cathodes by Strain Pengfei Liu1, Wei He1, Qingshui Xie1,* Yong Cheng1, Wanjie Xu1, Zhensong Qiao1, Laisen Wang1, Baihua Qu2, Zi-Zhong Zhu3, Dong-Liang Peng1,*

1Department

of Materials Science and Engineering, 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.

2Pen-Tung

Sah Institute of Micro-Nano Science and Technology, Xiamen University, Xiamen

361005, PR China 3Collaborative

Innovation Centre for Optoelectronic Semiconductors and Efficient Devices,

Department of Physics, Xiamen University, Xiamen 361005, China AUTHOR INFORMATION Corresponding Authors: [email protected] (Q. Xie), [email protected] (D. Peng)

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ABSTRACT: Lattice oxygen activity plays a dominant role in balancing discharge capacity and performance decay of lithium-rich layered oxide cathodes (LLOs). Based on density functional theory (DFT) and tight-binding theory, the activity of lattice oxygen can be improved by tensile strain, while suppressed by compressive strain. In order to verify the conclusion, LLOs with large lattice parameters (L-LLOs) were synthesized taking advantage of the lattice expansion effect in nanomaterials. Compared with conventional LLOs with small lattice parameters (S-LLOs), particles in L-LLOs are imposed by tensile strain. L-LLOs show a larger initial discharge capacity while decay faster in the prolonged cycles than S-LLOs. Actually, most of modified methods in LLOs can come down to strain-induced changes in lattice parameters. We believe this conclusion is a useful guideline to understand and tailor the lattice oxygen activity and may be generalized to other layered oxide cathodes involving anionic redox.

TOC GRAPHICS

KEYWORDS: lithium-rich; cathode materials; lattice oxygen; anionic redox; strain


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Lithium-rich layered oxides (LLOs) are among the most promising cathodes due to their higher reversible capacity than other conventional cathodes,1-3 such as LiCoO2,4 LiFePO4.5 However, they always suffer from poor rate capability, cycling fading and voltage decay, mainly originated from unstable lattice oxygen which also provides extraordinary capacity at the same time.6-9 Encouragingly, great improvements have been achieved by numerous methods like morphology control,10-11 surface coating12-13 and doping.14-16 For example, discharge capacity and rate capability can be largely promoted by carefully synthesizing bowl-like LLOs.11 Voltage decay can be suppressed by surface spinel coating.12 Besides, overall performance of LLOs including discharge capacity, voltage decay and rate capability can be improved by gradient Na+ doping.15 Whereas, mechanisms behind are still unclear. Fortunately, a common sense has been reached that all the problems and the corresponding solutions are closely related to lattice oxygen redox.17-19 Thus, a deep understanding of the lattice oxygen redox is urgently needed. Considering high crystallinity of the synthesized LLOs because of calcination at high temperature for hours,20-23 seeking an answer in solid physics is natural. According to the tight-binding theory that energy level splitting will intensify by reduced atomic distance,24 the activity of lattice oxygen is closely related to changes in lattice parameters, which can be seen as strain-induced and may be the operation rules behind these numerous modification methods.



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Therefore, density of states (DOS) of lattice oxygen in LLOs were carefully investigated by density functional theory (DFT) under compressive, zeroth and tensile strains. The activity of lattice oxygen is less active under compressive stain than that under tensile strain. And there is a positive correlation between the conductivity and lattice parameters of LLOs. Experiments were conducted to verify our theoretical conclusion using the lattice expansion effect in nanomaterials to introduce tensile stain.25-28 Compared with LLOs with small lattice parameters (S-LLOs), LLOs with large lattice parameters (L-LLOs) have a larger discharge capacity but a worse cycling ability, which coincides with our calculations. We believe this conclusion will be a useful guideline for understanding and tailing the activity of lattice oxygen not only in LLOs but also in other cathodes involving anionic redox. A previously published model of lithium-rich cathodes and Hubbard U values were adopted for all the calculations (Figure S1).14 Pseudopotentials were described by projector augmented wave (PAW) potentials29-30 in the Perdew−Burke−Ernzerhof (PBE) form31 within the frame of Vienna ab initio simulation package (VASP) .32 Energy cutoff was 600 eV. Converge criteria for electron and ion step were 10-5 eV and 0.05 eV/Å, respectively. A 2×2×1 Monkhorst-Pack grids were used for relaxation and a 4×4×2 grids for DOS calculations.33 Strain is defined as ε = (a-a0)/a0×100%, where a0 is the lattice parameter in equilibrium state and a is that in the strain-applied state.

Experiments were conducted to verify our aforementioned conclusion. An optimized ratio between transition metal elements and lithium element was adopted by a previous work during the 4 ACS Paragon Plus Environment

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calcination process.3 For the synthesis of S-LLOs, commercial carbonate precursors (Mn:Co:Ni=4:1:1) were used. For the synthesis of L-LLOs, stoichiometric acetates (Mn:Co:Ni=4:1:1) were dissolved in ethanol A and full dose oxalic acid was dissolved in ethanol B. Then solution A was pumped into B with continuous stirring. The mixed solution were aged for a whole night before washing and drying. Carbonate and oxalate precursors were firstly calcined at 500 C for 5 hours to get oxide precursors, then mixed with lithium carbonate and finally continuously calcined at 500 C for 2 hours and 900 C for 12 hours. The heating rate was 2 C per minute. Details of characterization techniques and half-cells assembly can be found in Note S1.

Figure 1. (a-c) Charge densities of selected transition metal slabs under different strains. Isosurface is 0.4 e Å-3. (d) PDOS of selected oxygen p orbital under different strains. 5 ACS Paragon Plus Environment


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Calculations were conducted to study the influence of strains on the activity of lattice oxygen. Charge densities of selected transition metal slabs are plotted in Figure 1a-c, where the isosurface is 0.4 e Å-3. As shown in the dashed boxes in Figure 1a-c, hybridizations between manganese 3d/4s orbitals and oxygen 2p orbital are strong along the linear O-Mn-O configuration, but those between lithium 2s orbital and oxygen 2p orbital are weak along the nearby linear O-Li-O configuration because of the huge energy difference between lithium 2s orbital and oxygen 2p orbital. However, the overlap between manganese and oxygen becomes less as with the increase of lattice parameters. This trend is in agreement with the tight-binding theory. Therefore, the binding energy will become smaller and lattice oxygen will be more active because electrons will be freer as a result of the increase of lattice parameters. Compared with no strain-applied situation, partial density of states (PDOS) of selected oxygen p orbital (Figure 1, Figure S2) show that both the spin up and spin down electrons in valence bands tend to move up toward Fermi level under tensile strain ε=+5%, while they move down to keep away from Fermi level under compressive strain ε=-5% (Figure 1d). Aggregation of electrons in valence bands near Fermi level means lattice oxygen will be more active in the electrochemical process, which is beneficial for discharging. Besides, oxygen vacancy formation energy does not change much in these systems (Table 1). However, there is positive correlation between lattice parameters and energy gap (Table S1), which also has a positive correlation with conductivity. Considering the applied external electric field during charge/discharge process, local electric field in crystal with larger parameters will be stronger.


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Thus, lattice oxygen in tensile system with enlarged lattice parameters will provide more capacity but be easy access to the formation of oxygen and vice versa in compressive system. Table 1. Oxygen vacancy formation energy34 (Evac) and band gap (Eg) of LLOs under compressive, zeroth and tensile strains. According to tight-binding theory, there is a negative correlation between band gap and lattice parameters. However, extra hybridizations will occur when ions are getting closely under compressive strains. So the results are not contrary to the predication.

ε (%)

-5

0

5

Eg (eV)

0.941

1.207

0.954

Evac (eV)

2.442

3.095

2.474

Figure 2. (a) XRD patterns. (b, c) Enlarged peaks of (003) (b) and (101) (c). (d) N2 adsorption/desorption isotherms. (e) Distributions of pore width.



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Experiments were conducted to verify our theoretical predication. Both the final products have the same chemical formula of Li1.08Mn0.54Co0.13Ni0.13O2. X-ray diffraction (XRD) patterns are shown in Figure 2. The main peaks can be indexed into α-NaFeO2, a typical trigonal structure with a space group R-3m, and peaks in the range of 20-23°correspond to the superlattice structure in C2/m Li2MnO3. (003) and (101) peaks of acetate-based products move toward low angle compared with carbonate-based products (Figure 2b,c), which means that S-LLOs and L-LLOs have been successfully synthesized according to our design. The enlarged parameters in L-LLOs were ascribed to the special synthesis process. Strong chelation between the transition metal ions and the oxalic ions promoted the formation of 3D network structure not the conventional globular structure during aging. When calcined at high temperature, the 3D secondary structure would be maintained, contributing to plenty of pores and the primary particle size would be small. Lattice parameters in small particles will be enlarged because of the lattice expansion effect in nanomaterials.25-28 Brunauer−Emmett−Teller (BET) method was used to further confirm the situation. The specific surface areas of S-LLOs and L-LLOs are 1.78 and 5.11 m2/g (Figure 2d), respectively. Both the samples have a broad pore size distribution range of 1.5-10 nm and pores in the range of 1.5-4 nm contribute to most of the total volume (Figure 2e). Obviously, L-LLOs have advantages over S-LLOs both in specific surface areas and pore volume in the range of 1.5-4 nm, which is accordance with our design. Therefore, lattice expansion will be more effective in LLLOs and larger lattice parameters are equivalent to withstand a tensile strain compared with SLLOs. 8 ACS Paragon Plus Environment

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Figure 3. (a, b) Schematic diagram of lattice expansion effect in nanomaterials. (c) Particle size distributions. (d-f) SEM image (d), HR-TEM image (e) and SEAD pattern (f) of S-LLOs. (g-i) SEM image (g), HR-TEM image (h) and SEAD pattern (i) of L-LLOs.

Due to the tiny difference between the XRD patterns of the two samples, DFT calculations were conducted to identify the changes of lattice parameters aroused by strains, which is introduced by surface reconstruction according to our design. D-spacing of (003) facet expands from 0.479 nm of the bulk to 0.499 nm of surface and the schematic diagram is shown in Figure 3a,b. However, this difference is between the ideal bulk and the ideal surface. In reality, both the S-LLOs and LLLOs have plenty of exposed surface. Thus, the difference shrinks to 0.0001 nm (Note S2) and may even be hard to identify in high-resolution transmission electron microscopy (HR-TEM). As 9 ACS Paragon Plus Environment


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showed in Figure 3c, L-LLOs has a wider particle size distribution and more content of particles below 5 nm than that of S-LLOs, which is consistent with our design. Then, energy disperse spectroscopy (EDS) mappings (Figure S3) of S-LLOs and L-LLOs evidence that elements distributions are uniform in both two samples. Scanning electron microscopy (SEM) images show that spherical secondary particles of S-LLOs are about 5-7 μm in diameter, while irregular secondary particles of L-LLOs can be more than 20 μm (Figure 3d,g). The primary particles of the former are 250-500 nm in size while those of the latter are 50-250 nm in size. Moreover, the dense degree of L-LLOs is higher than that of S-LLOs according to our design. The d-spacing of S-LLOs in HR-TEM image is 0.47 nm (Figure 3e), which can be indexed to (003) plane of layered lithiumrich oxides. The d-spacing of L-LLOs in Figure 3h is also 0.47 nm. The above situation seems to be conflict to our design that L-LLOs have larger lattice parameters than S-LLOs. However, this confusion can be easily understood when the differences between experiments and calculations are considered. One is that lattice parameters in these two situations are different. Besides, lattice parameters under strains in calculations are enlarged more intentionally to make the differences in PDOS of lattice oxygen distinct. Nevertheless, changes of the experimental lattice parameters are minor which cannot be distinguished in d-spacing even under the effect of lattice expansion in nanomaterials. But changes in XRD patterns will do. Furthermore, analysis shows that selected area electron diffraction (SAED) patterns of S-LLOs and L-LLOs are belong to {010} and {-110} planes, respectively (Figure 3f,i). Therefore, the influence of exposed surface on the charge/discharge capacity can be ignored. 10 ACS Paragon Plus Environment

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Figure 4. (a, b) Cyclic voltammograms (CV) of S-LLOs (a) and L-LLOs (b) at a scan rate 0.1 mV/s in the voltage range of 2.0-4.8 V. (c) Cycling performances at 0.2C in the voltage of 2.04.8 V after being activated at 0.1C (2.0-4.6 V). (d) The corresponding average discharge voltage.

Based on the above results, electrochemical tests of the two samples were conducted to verify our theoretical predication that L-LLOs would have a larger discharge specific capacity but a worse cycling stability than S-LLOs. Firstly, CV investigations of the two samples have been tested at a scan rate of 0.1 mV/s. As shown in Figure 4a,b both two samples have similar shapes in CV 11 ACS Paragon Plus Environment


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profiles. In the first cycle, the first oxidation peak moves from 4.11 V for S-LLOs to 4.03 V for LLLOs, demonstrating easier extraction of lithium ions in L-LLOs.16 The second oxidation peak around 4.6 V, which is a sign of irreversible lattice oxygen oxidization, is intensified in L-LLOs compared with that in S-LLOs. In the following cycles, the reduction peaks corresponding to the reduction of Mn4+ in the range of 3.0-3.6 V are more obvious in L-LLOs, which can be ascribed to the intercalation of more content of lithium ions. Thus, the introduced tensile strains can improve the activity of lattice oxygen and accelerate the electrochemical process. As for cycling performances, the first activation charge/discharge specific capacities of S-LLOs and L-LLOs are 311.4/232.7 mAh g-1 and 312.9/250.4 mAh g-1, corresponding to initial Columbic efficiencies 74.7% and 80.1%, respectively. The improved CE of L-LLOs can be ascribed to easy migration of lithium ions and decreased band gap introduced by tensile strains via surface reconstruction. However, there is no obvious difference in the flowing cycles. The operation voltage decrease of L-LLOs in the charge range of 3.5-4.5 V can be ascribed to the rising of fermi level (Figure S4).35 When cycled at 0.2 C in the range of 2.0-4.8 V, the first charge/discharge capacities of S-LLOs and LLLOs are 275.9/249.8 mAh g-1 and 293.1/267.7 mAh g-1 (Figure 4c). Although L-LLOs have a larger discharge capacity than S-LLOs in the first few cycles, capacity fading in L-LLOs is faster. After 50 cycles at 0.2 C, the discharge capacities of S-LLOs and L-LLOs are 240.7/211.4 mAh g1

with capacity retentions of 96.4% and 79.0%. As for average voltage, S-LLOs have advantages

over L-LLOs before 35 cycles and nearly become the same thereafter (Figure 4d). So the energy density of S-LLOs is much better than that of L-LLOs. The experimental results of the two samples 12 ACS Paragon Plus Environment

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are quite in agreement with our theoretical predication that L-LLOs under tensile strains will have a larger discharge specific capacity but a worse cycling stability than S-LLOs under compressive stains due to the closely related lattice oxygen activity.

Rational explanations of the above phenomena can be found if oxygen redox, lattice expansion effect and tight-binding theory are considered together. Tight-binding theory tells that the splitting of energy level will intensify with the decreasing of interatomic spacing. And lattice expansion effect thinks that lattice parameters will be enlarged as a result of surface reconstruction. Increase of the lattice parameters would weaken the splitting of energy level in LLOs and induce the upward movement of top valence bands of lattice oxygen. Thus, the activity of lattice oxygen in L-LLOs would be strengthen yet that in S-LLOs would be weaken. The conclusion was verified by the PDOS of lattice oxygen p orbital. Besides, energy gap was reduced with the increase of lattice parameters, leading to the increase of conductivity. Based on the above analysis, it can be predicted that L-LLOs would have a larger discharge capacity but a worse cycling stability than S-LLOs. The electrochemical performances of S-LLOs and L-LLOs are just the same with our theoretical prediction. Another interesting fact is that the increase of discharge capacity will induce the decrease of average voltage, which is contradict to what we want. In fact, the unit of discharge capacity, a sign of the ability to store charge, is Coulomb. The energy, product of discharge capacity and operation voltage, is exactly what we want. According to energy conservation law, chemical energy in a specific material has its threshold value and there is always a trade-off 13 ACS Paragon Plus Environment


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between discharge capacity and average voltage. Therefore, the one-sided pursuit of discharge capacity may be not the right direction. But the energy density of one specific material can be improved within the limit of its threshold value.

In summary, the conclusion that the activity of lattice oxygen can be improved by tensile strain while suppressed by compressive strain is predicted by DFT calculations and verified by experiments. We believe the above conclusion is a useful guideline for tailoring lattice oxygen activity in lithium-rich layered cathodes in experiments and can also be used for enhancing the understanding of mechanisms behind such as anionic redox, surfacing coating, doping and lattice oxygen evolution during charge/discharge process.

ASSOCIATED CONTENT

Supporting Information. The following files are available free of charge

Calculation model, PDOS of other selected lattice oxygens under strains, EDS mapping of the two final products, initial charge/discharge curves of the two sample at 0.1 C in the voltage range of 2.0-4.6 V, band gap information under different strains, details of characterization techniques and half-cells assembly, difference between the d-spacing of (003) facet of S-LLOs and L-LLOs deduced from XRD patterns AUTHOR INFORMATION


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Corresponding Author: [email protected] (Q. Xie), [email protected] (D. Peng) Notes The authors declare no competing financial interests.

ACKNOWLEDGMENT This study was financially supported by the National Key R&D Program of China (No. 2016YFA0202602), 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 R.J. Xie, Y.X. Zhuang, T.L. Zhou, L. Zhang and Y. Lv for the support in materials synthesis, F. Zheng for the discussion of DFT calculations. Jinming Wang for the support in TEM imaging.

REFERENCES (1) Whittingham, M. S., Ultimate Limits to Intercalation Reactions for Lithium Batteries. Chem. Rev. 2014, 114, 11414-11443.



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(2) Lee, J.; Kitchaev, D. A.; Kwon, D.-H.; Lee, C.-W.; Papp, J. K.; Liu, Y.-S.; Lun, Z.; Clément, R. J.; Shi, T.; McCloskey, B. D.; Guo, J.; Balasubramanian, M.; Ceder, G., Reversible Mn2+/Mn4+ Double Redox in Lithium-Excess Cathode Materials. Nature 2018, 556, 185-190. (3) Pimenta, V.; Sathiya, M.; Batuk, D.; Abakumov, A. M.; Giaume, D.; Cassaignon, S.; Larcher, D.; Tarascon, J.-M., Synthesis of Li-Rich NMC: A Comprehensive Study. Chem. Mater. 2017, 29, 9923-9936. (4) Mizushima, K.; Jones, P. C.; Wiseman, P. J.; Goodenough, J. B., LixCoO2 (0

A Guideline for Tailoring Lattice Oxygen Activity in Lithium-Rich

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