Twin boundaries induced by high-temperature shock boost the structural stability of Li-rich layered-oxide†

Zhedong Liu ^a, Cuihua Zeng ^a, Jingchao Zhang ^a, Jiawei Luo ^a, Zhaoxin Guo ^a, Zekun Li ^a, Rui Liu ^b, Wei-Di Liu ^c, Jia Ding ^a, Yanan Chen *^a and Wenbin Hu *^a
aSchool of Materials Science and Engineering, Tianjin University, Tianjin 300072, China
^bSchool of Materials Science and Engineering, Shandong University of Science and Technology, Qingdao 266590, China
^cSchool of Chemistry and Physic, ARC Research Hub in Zero-emission Power Generation for Carbon Neutrality, Center for Materials Science, Queensland University of Technology, Brisbane, QLD 4000, Australia. E-mail: yananchen@tju.edu.cn; wbhu@tju.edu.cn

First published on 29th July 2024

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Yanan Chen

Yanan Chen is Professor in the School of Materials Science and Engineering, Tianjin University. He received his bachelor's degree and joint PhD degree from the University of Science and Technology Beijing/University of Maryland in 2012 and 2017, respectively. He was an advanced innovative fellow at Tsinghua University before joining Tianjin University. His research mainly focuses on nanomaterials, devices, and systems for advanced energy storage and conversion. His research interests include high-temperature shock (HTS), ultrafast nanomanufacturing, metastable high-throughput synthesis, emerging energy storage Li ions and beyond, catalysis, artificial intelligence and interdisciplinary fields. He has published more than 100 research papers in international famous journals, including Nat. Energy, Nat. Sustain., Nat. Commun., Sci. Adv., JACS, PNAS, Adv. Mater., Mater. Today, Nano Lett., Adv. Energy Mater., and ACS Nano.

1 Introduction

Li-ion batteries (LIBs) have revolutionized the production industry and the lifestyle of human beings since their invention and continue to exert important impact.^1–3 With the rapid development of electric vehicles, stringent requirements are raised for high-energy-density secondary batteries, which mostly depend on cathode materials.^4,5 To date, commercialized cathode materials mainly include lithium transition metal oxides and phosphates, such as LiCoO₂ (LCO), LiMn₂O₄ (LMO), LiNi_1−x−yMn_xCo_yO₂ (NMC) and LiFePO₄ (LFP). Li-rich manganese-based layered oxides (LROs) have been regarded as one of the most promising candidates among all cathode materials for next-generation LIBs owing to their high average discharge voltage (>3.5 V) and high discharge specific capacity (>250 mA h g⁻¹).^6–8 The redox of O²⁻ above 4.5 V (vs. Li⁺/Li), accompanied with the electrochemical activation of Li₂MnO₃ components endows LRO materials with extraordinarily high discharge capacity.^9,10 However, reactions at high voltages trigger notorious structural transformation from layered to spinel/rock-salt structures and irreversible oxygen loss, which lead to capacity and voltage fading.^11,12 Additionally, cracks originating from anisotropic shrink and expansion accelerate structural collapse and performance degradation.¹⁰ Tremendous efforts, including element doping,^13,14 surface coating and structural/morphology design have been applied to improve the cycling stability of LROs.^14–16 However, these methods also exert negative effects, such as decreased capacity and coulombic efficiency (CE). Recently, defect engineering has been widely applied to nanomaterials and is considered a significant method to regulate the crystal structure and improve electrochemical performance.^17–19 Defects, including point defects, linear defects and planar defects, can disturb the surrounding atoms and introduce lattice distortion into materials. The type and concentration of defects determine the degree of lattice distortion. Defects can refrain movement of TM ions in materials, decrease strain accumulation, and prevent crack generation, which can effectively regulate the crystal structure of cathodes.²⁰ Pan et al. imported TBs into spinel LiMn₂O₄, which accelerated Li ion diffusion and improved rate performance.¹⁸ Twin boundaries commonly exist in metallic materials and exert positive effects on their mechanical properties.^21–23 TBs with low defect-energy levels act as rapid ion diffusion channels to accelerate ion diffusion.

In this work, we successfully prepared metastable Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ materials with TBs by an ultrafast HTS strategy. In such a short period of treating time, the reaction temperature instantly reaches a high level, facilitating the rapid transformation of substances. The whole process, featured by extreme rates of temperature rise and fall (10⁵ K s⁻¹), induces unique metastable structures of substances through non-equilibrium thermodynamic synthesis and kinetic modifications. Here, TBs were extensively observed by scanning transmission electron microscopy (STEM) due to ultra-fast cooling. The existence of TBs decreases the barrier of TM ion diffusion from the TM layer to the Li layer, leading to the formation of rock-salt phase around the TBs. The rock-salt phase can also stabilize the crystal structure due to its rigid framework that mitigates the inherent phase transformation and cracks of Li-rich layered-oxide materials, leading to positive influence on the electrochemical performance. The LRO with TBs calcined by HTS deliver a higher reversible capacity of 278 mA h g⁻¹ at 0.1C, an initial CE of 81% and a capacity retention of 89.4% after 100 cycles at 0.5C, compared with the materials synthesized by tube furnace calcination. The structural evolution results indicate that TBs hinder the generation and propagation of cracks by alleviating the anisotropic expansion and contraction of lattice parameters. Simultaneously, the TBs can restrain the irreversible phase transition from the layered to spine phase due to the presence of rock-salt phase. The differences in morphology, structure and the mechanism, in which the LRO electrode with TBs leads to enhanced performances, are shown in Fig. 1. This work extends the methods for synthesizing defective materials and offers profound implications for designing the crystal structure of materials.

2 Experimental section

2.1 Materials and reagents

All reagents were directly used after purchase without further purification in this work. The Li-rich manganese-based precursor Ni_1/6Co_1/6Mn_2/3CO₃ prepared by a co-precipitation method was purchased from Haian Battery Materials Technology Co., Ltd LiCO₃·H₂O (99%) was purchased from Shanghai Aladdin Biochemical Technology Co., Ltd.

2.2 Material synthesis

Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ with TBs prepared by the HTS method is named LRO-TB and Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ prepared by tube furnace calcination is named LRO. To obtain the layered oxide Li_1.2Ni_0.13Co_0.13Mn_0.54O₂, Ni_1/6Co_1/6Mn_2/3CO₃ was first mixed with Li₂CO₃ by ball milling and a nickel foil was used as a heating container. The mixtures were evenly spread onto a nickel foil (2 × 5 cm²), and the nickel foil loaded with mixtures was linked to a direct-current source with the current pulse for 45 s in an HTS setup (Shenzhen Zhongkejingyan Company) and heated in air. The temperature of the heater was tuned by adjusting the current and voltage and monitored using a laser infrared thermometer, and then the cathodes were obtained after thermal shock. The comparison samples were calcined in a tube furnace (TF).

2.3 Material characterization

Transmission electron microscopic (TEM), high-resolution transmission electron microscopic (HRTEM) and scanning transmission electron microscopic (STEM) images of the relevant samples mentioned in this work were obtained using JEM-2100F and JEM-ARM200F instruments (accelerating voltage of 200 kV), respectively. Scanning electron microscopy (SEM) and energy-dispersive spectroscopy (EDS) were performed using a JEOL JSM-7800F microscope to obtain the morphology and elemental distribution of materials. The crystal structure of all materials was evaluated in the 2θ range of 10°–80° with a scan speed of 2° min⁻¹, using a Bruker D8 advance X-ray diffractometer (Cu Kα radiation diffractometer, λ = 1.5406 Å) at a working voltage of 40 kV and a working current of 40 mA. The Rietveld refinement of X-ray diffraction (XRD) was carried out using the EXPGUI-GSAS program. The X-ray photoelectron spectroscopy (XPS) spectrum (Al Kα, hν = 1486.6 eV, working voltage of 14.6 kV) of all materials were recorded using a 250XI X-ray photoelectron spectrometer (US) to analyze the chemical environment of elements. All the bonding energies were calibrated with the C 1s signal at 284.8 eV. The Raman spectra (Xplora, wavelength of 514 nm) of materials were used to identify the micro-zone structure.

2.4 Electrochemical test

The electrochemical performance was evaluated using coin cells (CR2032) assembled in an argon-filled glove box (H₂O < 0.1 ppm and O₂ < 0.1 ppm). The active cathode materials, Super P (conductive agent) and binder (polyvinylidene fluoride, PVDF) were dissolved into N-methyl-2-pyrrolidone (NMP), at a mass ratio of 8/1/1, to prepare a cathode slurry. Then, the slurry was coated onto an Al current collector and dried in a vacuum oven at 100 °C for 10 h. Then, the Al foil collector was punched into disks with a diameter of 9 mm. A lithium metal foil was used as the anode and a polypropylene membrane (Celgard 2500) as the separator. The mass loading of LRO-TB and LRO cathodes was about 2 mg cm⁻². A high-voltage electrolyte was purchased from Dodo-Chem. The electrochemical performances were tested from 2 V to 4.8 V at different C-rates (1C corresponding to 200 mA g⁻¹) at room temperature (25 °C). All cells were charged and discharged for one cycle at 0.1C for activation before cycling. The cycling performance was conducted at 0.5C and 2C. The rate is tested from 0.1C to 10.0C. For the Galvanostatic Intermittent Titration Technique (GITT) measurement, batteries were first charged at a constant current density (0.05C) for 30 min, and then left at the open circuit voltage for 5 h. The batteries were discharged at 0.05C for 30 min and then left at the open circuit voltage for 5 h, and the above-mentioned processes were repeating until the voltage drops to 2.0 V. Cyclic voltammetry (CV) was performed in the range of 2.0–5.0 V using an electrochemical workstation (Chen Hua), at scan rates of 0.1–1 mV s⁻¹. Electrochemical impedance spectroscopy (EIS) was performed at a frequency in the range of 10⁵–0.01 Hz and a test voltage of 3.0 V with a perturbation amplitude of 5 mV.

3 Results and discussion

3.1 Characterization of LRO and LRO-TB

The crystal structure of LRO-TB and LRO was characterized by X-ray diffraction (XRD) with the Rietveld refinement. As shown in Fig. 2a and b, most peaks can well match with the hexagonal layered α-NaFeO₂ structure of the Rm space group (JCPDS No. 89-3601). Another superlattice peak appearing at 20–23° can be indexed to the space group of C2/m (JCPDS No. 84-1634), resulting from the short-range ordering of Li⁺ and Mn⁴⁺ of Li₂MnO₃ to form LiMn₆.⁷ The obvious splitting of (006)/(102) and (018)/(110) peaks, and c/a beyond 4.9, indicates that LRO-TB and LRO exhibit a good layered structure.²⁴ From Table S1 (ESI),† it can be observed that LRO-TB delivers a larger interplanar spacing of c-axis, which is beneficial to accelerate Li ion diffusion and relax strain generated by anisotropic lattice expansion and contraction.²⁵ The atom occupancy of samples is shown in Tables S2–3.† The intensity ratio, I₍₀₀₃₎/I₍₁₀₄₎, is an indicator of ordering degree for Li⁺ and Ni²⁺ in layered cathodes. The ratio beyond 1.2 indicates that LRO-TB and LRO deliver a highly ordered structure. Ni elements play a role of “support column” when Ni²⁺ migrates to the Li layer and replace Li⁺. Moderate Li⁺/Ni²⁺ intermixing can weaken the electrostatic repulsion of oxygen layer, which reduces structural damage during Li ion (de)-intercalation processes.^25–28 For comparison, the higher Li⁺/Ni²⁺ intermixing of LRO-TB (1.89%) than LRO (0.88%) is beneficial to improve the structural and cycling stability.

As shown in Fig. 2c, the LRO-TB electrode exhibits a microsphere-like morphology with a diameter of ∼10 μm, which consisted of primary nanoparticles with a size of 100 nm. The diameter of primary particles of LRO (Fig. 2d) is slightly larger, 150–300 nm, which prolongs the distance for Li ion diffusion and is unfavorable for rate performance.¹⁴ Abundant grain boundaries (GBs) between primary nanoparticles can be clearly observed on the surface of LRO-TB, which can impede the formation and propagation of cracks along and across the GBs.^2,29 In contrast, LRO is heated for a long time at high temperatures, leading to more intense melting and recrystallization processes of primary particles, leading to more blurred GBs of primary particles. Compared with LRO, the more compacted surface of LRO-TB strengthens the bonding force of primary particles, which is beneficial to decrease the side reaction between the active material and the electrolyte. The particle diameter and morphology of secondary particles for LRO are similar to those of LRO-TB, while the surface of LRO-TB is smoother than that of LRO (Fig. S1, ESI†). The energy-dispersive X-ray spectroscopy (EDS) elemental mapping results (Fig. S2, ESI†) indicate that the elements of Ni, Co, Mn and O are uniformly distributed on the surface of LRO-TB and LRO. As shown in Fig. 2e, LRO-TB delivers a typical interplanar spacing of 0.47 nm, corresponding to the (001) plane of Li₂MnO₃ (C2/m) and the (003) plane of LiMO₂ (Rm). Moreover, the lattice fringes of LRO-TB are partially mismatched, shown with red ellipses, which can be attributed to lattice distortion caused by twin boundaries. LRO delivers an interplanar spacing of 0.426 nm in Fig. 2f, corresponding to the (020) plane of Li₂MnO₃ (C2/m). As shown in Fig. S3 (ESI),† the increased strain triggered by defects from 0.208% of LRO to 0.25% of LRO-TB indicates the growth of defects in LRO-TB.³⁰ The presence of twinning defects tunes the crystal structure of LRO-TB and introduces faster ion diffusion channels, further improving the electrochemical performance of LRO-TB.

The Raman spectrum can characterize the overall structure of materials and is sensitive to changes with the TM–O bond between layered and spinel-like phase, which is applied to further explore the micro-zone phase structure of materials. As shown in Fig. S4 (ESI),† there are four main peaks in the Raman spectra of LRO-TB and LRO. The peaks at ∼470 cm⁻¹ and 599 cm⁻¹ match with the E_g vibration (O-TM-O bending mode) and A_1g vibration (TM-O stretching mode) of the LiMO₂ (Rm) component, respectively. The peaks at ∼420 cm⁻¹ and 540 cm⁻¹ matched with the phonon vibration of the Li₂MnO₃ (C/2m) component.³¹ The A_1g peak of the spinel-like phase appears at a higher position at ∼645 cm⁻¹ compared with the layered phase. The A_1g peak of LRO-TB causes a small blue shift, which indicates that LRO-TB synthesized by the HTS is the main layered structure containing spinel-like phase. LRO synthesized by tube furnace calcination is a single layered structure, as indicated by its A_1g peak which is sharp and located at ∼599 cm⁻¹.

The valence of the Ni, Co, Mn and O elements for LRO-TB and LRO was determined by X-ray photoelectron spectroscopy (XPS). All spectra were calibrated using C 1s 284.8 eV. As shown in Fig. 2g, the fitting and deconvolution of Ni 2p in LRO-TB reveals two main spin–orbit lines with a spin–orbit separation energy of 17.7 eV. The Ni 2p_3/2 peak is located at 853.8 eV with a satellite peak at 860.6 eV and the Ni 2p_1/2 peak is located at 871.5 eV with a satellite peak at 878.3 eV, which indicate that the main valence state of the Ni cation is +2.³² In addition, two weak peaks with binding energies of 855.4 eV and 873.1 eV are observed by the Gaussian fitting method, indicating the existence of Ni³⁺. The Co 2p (Fig. 2h) spectra show two main peaks around 779.3 eV (Co 2p_3/2) and 794.3 eV (Co 2p_1/2) with a spin–orbit separation energy of 15.0 eV. Two satellite peaks at 788.8 eV and 803.8 eV can be attributed to Co³⁺. The two weak peaks can be further detected around 781.4 eV and 796.4 eV, indicating the existence of a small amount of Co²⁺.^33,34 The Mn 2p (Fig. 2i) reveals two main spin-splitting peaks, Mn 2p_3/2 at 641.1 eV and Mn 2p_1/2 at 652.7 eV with a spin–orbit separation energy of 11.6 eV, which can be ascribed to Mn⁴⁺. There are also two Mn³⁺ peaks at around 642.1 eV and 653.7 eV.^35,36 The appearance of Ni³⁺ and Mn³⁺ is attributed to electron transfer between Ni²⁺ and Mn⁴⁺.³⁷ The O 1s spectrum of LRO-TB contains three peaks in Fig. 2j. The peak at 529.5 eV is the lattice oxygen between metal and oxygen. The peak at 531.2 eV is related to the oxygen vacancies in the bulk phase. The peak at 534.0 eV is related to the adsorbed oxygen species on the surface and oxidation products of lattice O²⁻.^38,39 As shown in Fig. S5 (ESI),† the peak positions for LRO after fitting are close to those of LRO-TB. Differently, LRO-TB exhibits a higher amount of Mn⁴⁺ (64.2%), Co³⁺ (70.6%) and oxygen vacancies (41.0%) compared with Mn⁴⁺ (56.9%), Co³⁺ (64%) and oxygen vacancies (36.2%) of LRO. The Jahn–Teller distortion of Mn³⁺ leads to dissolution of manganese, which will destroy structural stability and cause capacity attenuation.⁴⁰ The higher ratio of Mn⁴⁺/Mn³⁺ for LRO-TB is beneficial for enhanced structural stability.

As shown in Fig. 3, spherical aberration-corrected STEM investigations were conducted to further observe the detailed microstructure of samples. Fig. 3a shows the STEM image of the primary particle at a low magnification for LRO-TB. Fig. 3b shows the enlarged image of the red square region displayed in Fig. 3a. The zigzag-shaped atomic ladder structure is observed at the edge of LRO-TB due to the existence of oxygen vacancies. By contrast, LRO (Fig. 3c) delivers a flat surface. The existence of the special structure for LRO-TB can provide more active sites for Li ion (de)-intercalation and decrease the thickness of cathode electrolyte interphase (CEI) during cycling.⁴¹ The atomic-resolution high-angle annular dark-field (HAADF) STEM images of LRO-TB and LRO were acquired along the 〈010〉 direction. In Fig. 3d, the alternate arrangement of the Li layer and TM layer along the c axis is clearly observed in layered LRO with the Rm structure. The contrast of atomic columns is roughly proportional to the square of the atomic number (Z-contrast) for which reason it is only sensitive to heavy elements. Correspondingly, the bright spots stand for Ni, Co and Mn atoms, while the Li and O atoms are virtually indiscernible under this imaging condition.⁴² The green and purple arrows reveal the Li layer and TM layer, respectively. As shown in Fig. 3e, LRO-TB exhibits a twinning structure which consisted of a main layered structure with the well-grown rock-salt (Fmm) phase. The STEM image delivers a typical layered arrangement of TM in one domain (down) and rock-salt phase in the other domain (up). The compatibility intergrown structure solidly anchors the rock-salt domain in the main layered domain. The domain with additional images from different regions is presented in Fig. S6 (ESI†) and also shows the twinning defect. The Fast Fourier Transform (FFT) pattern (Fig. 3f) also confirms the existence of twinning structure for two sets of diffraction patterns, and another diffraction spot marked by purple circle is indexed to the (00) plane of rock-salt with Fmm. The calcination methods for LRO-TB and LRO are different which should be the origin of TBs in LRO-TB. The amounts of Li ions are embedded into the TM oxide and uneven distributed Li-rich areas cannot diffuse to the normal sites in the surface of large particles due to short duration in HTS, resulting in the accumulation of stress. The appearance of TBs can relieve stress caused by lattice expansion.⁴³ Molecular dynamics (MD) simulations reveal that the rapid cooling process controlled by dynamics under non-equilibrium conditions can promote the formation of Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ with abundant TBs, and the ultrafast cooling rate can well retain such TBs to materials. The twinning region is enlarged to further verify these features and the atoms in TB are marked in orange in Fig. 3g. The STEM image shows that the TM layer is closely arranged and TM ions occupy the Li sites to form rock-salt phase. The atomic intensity profile also indicates the cation disorder. Specifically, the signal intensity of Li atoms is very weak so only signal peaks of TM atoms can be observed for Li_1.2Ni_0.13Co_0.13Mn_0.54O₂, consistent with LRO (Fig. 3h). However, The Li sites of LRO-TB present a stronger signal intensity (Fig. 3i), indicating that TM ions migrate and occupy the Li sites. The Inverse Fast Fourier Transform (IFFT) pattern of LRO-TB (Fig. 3j) clearly shows the twinning structure with symmetry atom distribution where TB is marked with a blue arrow. The first-principles calculations of the kinetics of a series of reaction steps indicate that TBs can decrease the energetic barriers of TM ion migration and facilitate TM ion migration from TM sites to Li sites, so that they promote gradual growth of the rock-salt phase around TBs.⁴⁴ In addition, the formation of rock-salt is related to oxygen vacancies which reduce the TM–O binding energy and promote the TM ion migration into the Li layer. Furthermore, the rock-salt phase could be NiO because the Ni element is more mobile than Co and Mn.⁴⁵ The rock-salt phase has positive effects on cycling performance due to its stable structure.^46,47Fig. 3k and l show the FFT and IFFT patterns of LRO, corresponding to Fig. 3d. It is concluded from the above-mentioned results that LRO delivers a typical layered structure without TBs. The long reaction time in tube furnace calcination leads to sufficient migration of atoms and crystal growth, which eliminate lattice distortion during sluggish heating and cooling processes and lead to the formation of a perfect crystal structure.

3.2 Electrochemical performance

To analyze the effect of calcination temperature and time on crystal structure and electrochemical performances, a series of Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ by the HTS method at different temperatures for different time durations were synthesized. Fig. S7 (ESI†) shows the XRD patterns of samples calcined by HTS. The optimal calcination temperature was determined to be 850 °C to endow a desirable electrochemical performance (Fig. S9a and b, ESI†). Too low temperature is unfavorable for the growth of ideal crystal structure and too high temperature leads to severe loss of Li element, which deteriorate the electrochemical performance. Based on the single variate, the sample calcined at 850 °C for 45 s delivers the most excellent electrochemical performances including initial discharge specific capacity, CE and cycling performance (Fig. S9c and d, ESI†). The changes in crystal structure with calcination time attract our attention and Rietveld refinements are shown in Fig. S8 and Table S4 (ESI).† The phase content of C/2m gradually decreases and Rm increases as the calcination time prolongs, attributed to the increased Li loss. All materials possess a good layered structure indicated by c/a close to 5.0, and the decreased value of LRO-TB originates from abundant TBs. Moreover, long calcination time reduces the cation disorder of LRO-TB. In addition, the higher Li⁺/Ni²⁺ mixing ∼2% can improve structural stability in a high-order material. In addition, the interplanar spacing and volume first increase and then decrease as time increases, related to changes in repulsive force in the O layer.⁴⁸ Appropriate lattice parameters of LRO-TB not only reduce structural destruction but also accelerate Li ion diffusion, contributing to improved cycling stability. The above-mentioned results provide evidences for the outstanding electrochemical performance of LRO-TB.

The electrochemical performances of LRO-TB by HTS and LRO by tube furnace calcination are compared in detail. The initial charge–discharge curves of LRO-TB and LRO at 0.1C show two charge–discharge plateaus in Fig. 4a, with a slope below 4.5 V and a long plateau at 4.5 V. The first plateau is attributed to Li delithiation from the Li layer accompanied by the oxidation of Ni²⁺ and Co³⁺. In addition, the second plateau stands for further Li delithiation from the TM layer accompanied by oxidation of lattice oxygen and the electrochemical activation of Li₂MnO₃ contents.⁴⁹ The LRO-TB delivers a charge capacity of 343 mA h g⁻¹, a discharge capacity of 278 mA h g⁻¹ and an initial CE of 81.0%. By contrast, the LRO electrode shows a charge capacity of 351 mA h g⁻¹, a discharge capacity of 271 mA h g⁻¹ and an initial CE of 77.2%. The charge capacity of LRO-TB is lower than that of the LRO electrode, which is consistent with the shortened activation platform at 4.5 V and can be attributed to the small amount of Li₂MnO₃ being consumed to form the layered/rocksalt phase.⁴² The discharge capacity of LRO-TB is higher than that of LRO, contributing to a higher initial CE. In addition, the irreversible capacity is reduced from 80 mA h g⁻¹ to 65 mA h g⁻¹. These phenomena can be attributed to the enlarged d-spacing and the reduction of side reactions between the electrode material and the electrolyte, which reduce Li loss, cause irreversible release of oxygen and increase the reversible capacity of LRO-TB. The first cyclic voltammetry (CV) curves at 0.1 mV s⁻¹ of LRO-TB and LRO (Fig. S11a, ESI†) display two sets of redox peaks, which is consistent with the initial charge–discharge curves. The peak intensity related to lattice oxygen slightly decreases for LRO-TB, indicating slightly less irreversible oxygen release.⁵⁰

The cycling performances of LRO-TB and LRO electrodes at 0.5C and 2C are displayed in Fig. 4b and c. All batteries are activated at 0.1C for one cycle before cycling in order to prompt (de)-intercalation of Li ions, which is conducive to the development of capacity. As shown in Fig. 4b, LRO-TB and LRO deliver discharge capacities of 210 mA h g⁻¹ and 187 mA h g⁻¹ with capacity retention values of 89.4% and 81.0%, respectively, after 100 cycles. Furthermore, LRO-TB maintains 72.6% of initial capacity, while LRO merely delivers 56% of initial capacity after 200 cycles. From Fig. 4c, the maintaining capacities of LRO-TB and LRO are 150 mA h g⁻¹ and 126 mA h g⁻¹ after 200 cycles, with capacity retention values of 76.8% and 66.3%, respectively. As shown, LRO-TB exhibits a much better cycling stability, originating from the increased ratio of Mn⁴⁺/Mn³⁺ and the occurrence of stable rock-salt phase generated by TB regulation. Fig. S11b (ESI†) discloses that the voltage decay of LRO-TB (0.475 V) is smaller than that of LRO (0.611 V) after 200 cycles at 2C.

Rate capability is also one of the key performance indicators of power battery cathodes. The rate performances of LRO-TB and LRO are shown in Fig. S10.† The LRO-TB cathode provides reversible discharge capacities of 277, 223, 203, 175, 119 and 76 mA h g⁻¹ at 0.1, 0.5, 1, 2, 5 and 10C, respectively, as well as recovers to 254 mA h g⁻¹ at 0.1C. The rate property of LRO is slightly worse than LRO-TB. As shown in Fig. S11c (ESI),† the initial charge–discharge curves of LRO-TB at different rates show a smaller gap than that of LRO, indicating that LRO-TB possesses a smaller polarization. The Galvanostatic Intermittent Titration Technique (GITT) was performed to analyze the Li ion diffusion kinetics of the initial charging process. As shown in Fig. S12 (ESI),† the Li ion diffusion coefficients (D_Li⁺) of materials are located in the range of 10⁻¹⁵–10⁻¹⁰ cm² s⁻¹. D_Li⁺ steadily decreases before 4.5 V for LRO-TB while sharply decreases around 4.0 V for LRO. The largely decreased D_Li⁺ of LRO is associated with structural instability with Li ions extracted from the Li layer.⁵¹ TBs, acting as fast diffusion channels, accelerate Li ion diffusion and provide more active sites, which can promote the electrochemical reaction kinetics of LRO-TB.^18,22 The inert Li₂MnO₃ activation control the reaction kinetics above 4.5 V, leading to severely decreased D_Li⁺, accompanied by Li ion extraction, oxygen loss and structural rearrangement.^7,41 It is noteworthy to mention that the Li⁺ diffusion coefficient of LRO-TB is an order of magnitude higher than that of LRO, which is associated with the presence of TBs that provide a preferable pathway and a lower migration energy for the migration of Li⁺. Electrochemical impedance spectroscopy (EIS) was adopted to analyze the electrochemical reaction kinetics of cathodes before and after 100 cycles at 2C, respectively. The results were fitted with the equivalent circuit model, as shown in Fig. 4d and e. All curves reveal similar characteristics with a mid-frequency semicircle and a low-frequency sloping line. The intercept of EIS in the high-frequency region refers to the ohmic resistance (R_o) of the electrolyte. The semicircle in the mid-frequency region is related to charge transfer resistance (R_ct) at the electrode/electrolyte interface. In addition, the sloping line in the low-frequency region is related to the Weber resistance associated with Li ion diffusion in solids.³⁷ LRO-TB delivers the smaller R_ct and Weber resistance before cycling, which are favourable for electron transfer and Li ion (de-)intercalation across the interface.⁵² The side reaction at the electrode–electrolyte interface leads to the accumulation of by-products, which increases the impedance of cathodes and decreases the charge transport rate through the CEI film during cycling. LRO-TB gets a smaller increase of impedance after cycling. These results suggest that TBs play a positive role in improving the comprehensive electrochemical performances of LRO-TB.

The excellent electrochemical performances of LRO-TB endow it with practical applications. The electrochemical performance of LRO-TB was evaluated by assembling full cells with graphite as the anode. The initial charge–discharge curves at 0.1C are shown in Fig. 4f. The voltage distribution of full cells is similar to that of half-cell. Specifically, the full cell delivers a charge capacity of 346 mA h g⁻¹, a discharge capacity of 276 mA h g⁻¹ and an initial CE of 79.8%. The cycling stability of LRO-TB‖graphite was further evaluated at 0.5C, as shown in Fig. 4g, with an initial reversible capacity of 247 mA h g⁻¹. The capacity of LRO-TB‖graphite rapidly decays in first 10 cycles and then is relatively stable in subsequent cycles, which can be attributed to the formation of the CEI film. The capacity retention is 92.4% after 50 cycles when the capacity is stable.

3.3 Structural evolution of cathodes

The electrochemical performance of cathodes is closely related to evolution of structure and morphology during cycling, for which reason the mechanism of the excellent performance of LRO-TB was analyzed by ex situ XRD, SEM and TEM. Fig. 5a shows the crystal structures of LRO-TB and LRO after 100 cycles at 0.5C. The (018)/(110) diffraction peaks of LRO almost disappear, suggesting the structural deterioration of LRO. However, the (018)/(110) diffraction peaks of LRO-TB still exist with splitting. As shown in Fig. 5b, the shift and decreased intensity of the (003) peak for LRO-TB are significantly smaller than those of LRO. These results indicate that LRO-TB well maintains the layered structure after cycling.⁵³ From the SEM images shown in Fig. 5c, S13a and b,† it can be observed that LRO-TB particles maintain the integrity and smoothness after 100 cycles at 0.5C, which is the same for the non-cycled material. For comparison, as shown in Fig. 5d, S12c and d,† LRO particles suffer from severe cracking due to surface and intergranular corrosion, which originates from the following three aspects: (1) continuous generation of cracks and expansion from bulk to surface, (2) micro-strain induced by volume change of primary particles, and (3) corrosion of particles by acidic compounds generated from electrolyte decomposition.⁵⁴ The cracks create more electrolyte-exposing fresh surface as well as disrupt bonding of primary particles, which lead to extensive CEI formation and separation of cathode particles, and seriously degrade the rate and cycling performance.^2,53 The decreased cracks of LRO-TB compared with LRO are because TBs can refrain the changes in anisotropic lattice parameters through mitigating intragranular and intergranular cracks, as well as alleviate the accumulation of micro-strains during cycling.⁵⁵ These are beneficial to maintain the spherical morphology and improve the cycling stability of LRO-TB. As shown in Fig. 5e–h and S14 (ESI),† the TEM and FFT pattern images indicate that LRO-TB maintains a better layered structure while LRO has experienced irreversible phase transformation from layered to spinel phase on the surface after cycling. The phase transformation results in structural degradation as well as attenuation of capacity and voltage. The more stable structure of LRO-TB can be attributed to the presence of the robust rock-salt phase without expansion and contraction, which can afford enhanced structural stability and inhibit further formation of harmful phase during cycling.⁵⁶ The formation of rock-salt phase can be further attributed to the decreased migration energy of TM ions by TBs. The elemental mappings (Fig. S15, ESI†) disclose uniform distribution of Ni, Co, Mn and O of materials after 200 cycles at 2C.

4 Conclusion

In summary, metastable Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ with abundant TBs as a high-performance cathode material for LIBs has been successfully synthesized by an ultrafast HTS method. The results indicated that the non-equilibrium conditions of rapid heating and cooling can introduce various crystal defects including TBs and oxygen vacancies into the synthesized Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ electrode. These defects exert significant effects on the improvement of the electrochemical performances of LRO-TB. Compared with LRO prepared by tube furnace calcination, LRO-TB with special TBs delivers less cracks and maintains a better layered structure, which yield remarkable electrochemical performance in terms of reversible capacity and cycling stability. The LRO-TB electrode expresses a reversible discharge capacity of 210 mA h g⁻¹ with a capacity retention of 89.4% after 100 cycles at 0.5C, which is much better than that of the LRO electrode (187 mA h g⁻¹, 81.0%). The structural evolution results indicate that TBs are beneficial to decrease the cracks of particles due to the alleviated anisotropic changes in lattice parameters and promoted formation of rock-salt phase to mitigate the notorious phase transformation from layered to spinel. This work provides a new strategy for the synthesis of electrode materials with beneficial defects through a non-equilibrium HTS technique, which can guide the development of high-performance LIBs in the future and achieve long cycling stability.

Data availability

The data supporting this article have been included as part of the ESI.†

Author contributions

C. Y. N. planned and supervised the project. L. Z. D. conducted material synthesis, characterization. L. Z. D. wrote the manuscript, and all authors contributed to the discussion and provided feedback on the manuscript.

Conflicts of interest

The authors declare no competing interests.

Acknowledgements

The authors acknowledge the financial support from the National Natural Science Foundation of China (92372107, 52171219).

References

1. X. Li, Y. Qiao, S. Guo, Z. Xu, H. Zhu, X. Zhang, Y. Yuan, P. He, M. Ishida and H. Zhou, Adv. Mater., 2018, 30, e1705197 CrossRef PubMed .
2. Y. Liu, J. Wang, J. Wu, Z. Ding, P. Yao, S. Zhang and Y. Chen, Adv. Energy Mater., 2019, 10, 1903139 CrossRef .
3. Y. Liu, H. Zhu, H. Zhu, Y. Ren, Y. Zhu, Y. Huang, L. Dai, S. Dou, J. Xu, C. J. Sun, X. L. Wang, Y. Deng, Q. Yuan, X. Liu, J. Wu, Y. Chen and Q. Liu, Adv. Energy Mater., 2021, 10, 2003479 CrossRef .
4. Z. Sun, L. Xu, C. Dong, H. Zhang, M. Zhang, Y. Ma, Y. Liu, Z. Li, Y. Zhou, Y. Han and Y. Chen, Nano Energy, 2019, 63, 103887 CrossRef CAS .
5. X. Ji, Q. Xia, Y. Xu, H. Feng, P. Wang and Q. Tan, J. Power Sources, 2021, 487, 229362 CrossRef CAS .
6. T. Nakamura, K. Ohta, X. Hou, Y. Kimura, K. Tsuruta, Y. Tamenori, R. Aso, H. Yoshida and K. Amezawa, J. Mater. Chem. A, 2021, 9, 3657–3667 RSC .
7. C. Cui, X. Fan, X. Zhou, J. Chen, Q. Wang, L. Ma, C. Yang, E. Hu, X. Q. Yang and C. Wang, J. Am. Chem. Soc., 2020, 142, 8918–8927 CrossRef CAS PubMed .
8. W. He, F. Ye, J. Lin, Q. Wang, Q. Xie, F. Pei, C. Zhang, P. Liu, X. Li, L. Wang, B. Qu and D. L. Peng, Nano-Micro Lett., 2021, 13, 205 CrossRef CAS PubMed .
9. Q. Li, Z. Yao, E. Lee, Y. Xu, M. M. Thackeray, C. Wolverton, V. P. Dravid and J. Wu, Nat. Commun., 2019, 10, 1692 CrossRef PubMed .
10. X. Fan, G. Hu, B. Zhang, X. Ou, J. Zhang, W. Zhao, H. Jia, L. Zou, P. Li and Y. Yang, Nano Energy, 2020, 70, 104450 CrossRef CAS .
11. L. Nie, Z. Wang, X. Zhao, S. Chen, Y. He, H. Zhao, T. Gao, Y. Zhang, L. Dong, F. Kim, Y. Yu and W. Liu, Nano Lett., 2021, 21, 8370–8377 CrossRef CAS .
12. E. Wang, D. Xiao, T. Wu, X. Liu, Y. Zhou, B. Wang, T. Lin, X. Zhang and H. Yu, Adv. Funct. Mater., 2022, 32, 2201744 CrossRef CAS .
13. H. Liu, C. Du, G. Yin, B. Song, P. Zuo, X. Cheng, Y. Ma and Y. Gao, J. Mater. Chem. A, 2014, 2, 15640–15646 RSC .
14. Q. Wang, W. He, L. Wang, S. Li, H. Zheng, Q. Liu, Y. Cai, J. Lin, Q. Xie and D.-L. Peng, ACS Sustain. Chem. Eng., 2020, 9, 197–206 CrossRef .
15. J. Wan, W. Chen, C. Jia, L. Zheng, J. Dong, X. Zheng, Y. Wang, W. Yan, C. Chen, Q. Peng, D. Wang and Y. Li, Adv. Mater., 2018, 30, 1705369 CrossRef .
16. Y. Zhang, L. Tao, C. Xie, D. Wang, Y. Zou, R. Chen, Y. Wang, C. Jia and S. Wang, Adv. Mater., 2020, 32, e1905923 CrossRef .
17. M. M. Rahman, W. Y. Chen, L. Mu, Z. Xu, Z. Xiao, M. Li, X. M. Bai and F. Lin, Nat. Commun., 2020, 11, 4548 CrossRef CAS .
18. R. Wang, X. Chen, Z. Huang, J. Yang, F. Liu, M. Chu, T. Liu, C. Wang, W. Zhu, S. Li, S. Li, J. Zheng, J. Chen, L. He, L. Jin, F. Pan and Y. Xiao, Nat. Commun., 2021, 12, 3085 CrossRef CAS PubMed .
19. Y. Pei, Q. Chen, M. Wang, B. Li, P. Wang, G. Henkelman, L. Zhen, G. Cao and C.-Y. Xu, Nano Energy, 2020, 71, 104644 CrossRef CAS .
20. J. F. Nie, Y. M. Zhu, J. Z. Liu and X. Y. Fang, Science, 2013, 340, 957–960 CrossRef CAS PubMed .
21. A. Nie, L. Y. Gan, Y. Cheng, Q. Li, Y. Yuan, F. Mashayek, H. Wang, R. Klie, U. Schwingenschlogl and R. Shahbazian-Yassar, Nano Lett., 2015, 15, 610–615 CrossRef CAS .
22. F. Zhang, J. Dong, D. Yi, J. Xia, Z. Lu, Y. Yang and X. Wang, Chem. Eng. J., 2021, 423, 130122 CrossRef CAS .
23. C. Liu, Y. Shen, J. Zhang, G. Li, X. Zheng, X. Han, L. Xu, S. Zhu, Y. Chen, Y. Deng and W. Hu, Adv. Energy Mater., 2022, 12, 2103505 CrossRef CAS .
24. H. Guo, Z. Wei, K. Jia, B. Qiu, C. Yin, F. Meng, Q. Zhang, L. Gu, S. Han, Y. Liu, H. Zhao, W. Jiang, H. Cui, Y. Xia and Z. Liu, Energy Storage Mater., 2019, 16, 220–227 CrossRef .
25. S. Hy, J.-H. Cheng, J.-Y. Liu, C.-J. Pan, J. Rick, J.-F. Lee, J.-M. Chen and B. J. Hwang, Chem. Mater., 2014, 26, 6919–6927 CrossRef CAS .
26. W. Huang, C. Lin, M. Zhang, S. Li, Z. Chen, W. Zhao, C. Zhu, Q. Zhao, H. Chen and F. Pan, Adv. Energy Mater., 2021, 11, 2102646 CrossRef CAS .
27. J. Han, H. Zheng, Z. Hu, X. Luo, Y. Ma, Q. Xie, D.-L. Peng and G. Yue, Electrochim. Acta, 2019, 299, 844–852 CrossRef CAS .
28. F.-D. Yu, L.-F. Que, Z.-B. Wang, Y. Zhang, Y. Xue, B.-S. Liu and D.-M. Gu, J. Mater. Chem. A, 2016, 4, 18416–18425 RSC .
29. A. O. Kondrakov, A. Schmidt, J. Xu, H. Geßwein, R. Mönig, P. Hartmann, H. Sommer, T. Brezesinski and J. Janek, J. Phys. Chem. C, 2017, 121, 3286–3294 CrossRef CAS .
30. W. He, P. Liu, B. Qu, Z. Zheng, H. Zheng, P. Deng, P. Li, S. Li, H. Huang, L. Wang, Q. Xie and D. L. Peng, Adv. Sci., 2019, 6, 1802114 CrossRef PubMed .
31. Y. Liu, Y. Chen, J. Wang, W. Wang, Z. Ding, L. Li, Y. Zhang, Y. Deng, J. Wu and Y. Chen, Nano Res., 2021, 15, 3178–3186 CrossRef .
32. J. Li, M. Li, L. Zhang and J. Wang, J. Mater. Chem. A, 2016, 4, 12442–12450 RSC .
33. H. Li, X. Wei, P. Yang, Y. Ren, S. Wang, Y. Xing and S. Zhang, Electrochim. Acta, 2018, 261, 86–95 CrossRef CAS .
34. L. Li, L. Wang, X. Zhang, M. Xie, F. Wu and R. Chen, ACS Appl. Mater. Interfaces, 2015, 7, 21939–21947 CrossRef CAS .
35. L. Li, L. Wang, X. Zhang, Q. Xue, L. Wei, F. Wu and R. Chen, ACS Appl. Mater. Interfaces, 2017, 9, 1516–1523 CrossRef CAS .
36. H. Li, S. Zhang, X. Wei, P. Yang, Z. Jian and J. Meng, RSC Adv., 2016, 6, 79050–79057 RSC .
37. J. Li, S. Xiong, Y. Liu, Z. Ju and Y. Qian, Nano Energy, 2013, 2, 1249–1260 CrossRef CAS .
38. T. Lin, T. U. Schulli, Y. Hu, X. Zhu, Q. Gu, B. Luo, B. Cowie and L. Wang, Adv. Funct. Mater., 2020, 30, 1909192 CrossRef CAS .
39. X. Ding, D. Luo, J. Cui, H. Xie, Q. Ren and Z. Lin, Angew. Chem., Int. Ed., 2020, 59, 7778–7782 CrossRef CAS PubMed .
40. L. Xu, P. Hou, Y. Zhang, H. Zhang, D. Song, X. Shi, X. Wang and L. Zhang, J. Mater. Chem. A, 2015, 3, 21219–21226 RSC .
41. Q. Li, D. Ning, D. Wong, K. An, Y. Tang, D. Zhou, G. Schuck, Z. Chen, N. Zhang and X. Liu, Nat. Commun., 2022, 13, 1123 CrossRef CAS .
42. C. Zhang, B. Wei, W. Jiang, M. Wang, W. Hu, C. Liang, T. Wang, L. Chen, R. Zhang, P. Wang and W. Wei, ACS Appl. Mater. Interfaces, 2021, 13, 45619–45629 CrossRef CAS PubMed .
43. T. Erichsen, B. Pfeiffer, V. Roddatis and C. A. Volkert, ACS Appl. Energy Mater., 2020, 3, 5405–5414 CrossRef CAS .
44. S. Li, Z. Yao, J. Zheng, M. Fu, J. Cen, S. Hwang, H. Jin, A. Orlov, L. Gu, S. Wang, Z. Chen and D. Su, Angew. Chem., Int. Ed., 2020, 59, 22092–22099 CrossRef CAS PubMed .
45. S. Hwang, S. M. Kim, S.-M. Bak, K. Y. Chung and W. Chang, Chem. Mater., 2015, 27, 6044–6052 CrossRef CAS .
46. Y. Cho, P. Oh and J. Cho, Nano Lett., 2013, 13, 1145–1152 CrossRef CAS PubMed .
47. C. Zhang, B. Wei, M. Wang, D. Zhang, T. Uchiyama, C. Liang, L. Chen, Y. Uchimoto, R. Zhang, P. Wang and W. Wei, Energy Storage Mater., 2022, 46, 512–522 CrossRef .
48. Y. Zhou, H. Cui, B. Qiu, Y. Xia, C. Yin, L. Wan, Z. Shi and Z. Liu, ACS Mater. Lett., 2021, 3, 433–441 CrossRef CAS .
49. D. Luo, G. Li, C. Fu, J. Zheng, J. Fan, Q. Li and L. Li, Adv. Energy Mater., 2014, 4, 1400062 CrossRef .
50. S. Y. Kim, C. S. Park, S. Hosseini, J. Lampert, Y. J. Kim and L. F. Nazar, Adv. Energy Mater., 2021, 11, 2100552 CrossRef CAS .
51. C. Han, J. Wu, S. Li, T. Li, J. Li and H. Liu, Solid State Ionics, 2021, 364, 115629 CrossRef CAS .
52. H. Li, Y. Wang, J. Zhang, L. Zhu, J. Shi, B. Chen, Y. Hu, H. Zhang, X. Deng and Y. Peng, Nanoscale, 2022, 14, 7692–7701 RSC .
53. E. Wang, Y. Zhao, D. Xiao, X. Zhang, T. Wu, B. Wang, M. Zubair, Y. Li, X. Sun and H. Yu, Adv. Mater., 2020, 32, e1906070 CrossRef PubMed .
54. H.-H. Ryu, K.-J. Park, C. S. Yoon and Y.-K. Sun, Chem. Mater., 2018, 30, 1155–1163 CrossRef CAS .
55. H. Chung, Y. Li, M. Zhang, A. Grenier, C. Mejia, D. Cheng, B. Sayahpour, C. Song, M. H. Shen, R. Huang, E. A. Wu, K. W. Chapman, S. J. Kim and Y. S. Meng, Chem. Mater., 2022, 34, 7302–7312 CrossRef CAS .
56. Y. Su, Y. Yang, L. Chen, Y. Lu, L. Bao, G. Chen, Z. Yang, Q. Zhang, J. Wang, R. Chen, S. Chen and F. Wu, Electrochim. Acta, 2018, 292, 217–226 CrossRef CAS .

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4ta03098b

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