Stabilization of lattice oxygen in lithium rich cathode materials via manipulating Ni content

Abstract

Li-rich layered oxides (LLOs) have been vigorously explored as promising cathodes for next generation Li-ion batteries due to their high capacity. But their adoption is hampered by poor stability and serious voltage decay during cycling due to the oxygen vacancies (V_o) existing within the active particles and the V_o sites can result in structural reconstruction of the material. Herein, we focus our research on controlling the oxygen vacancies by adjusting the Ni content in Li-rich cathodes to improve their cycling stability and suppress the voltage decay. Based on comprehensive physical characterization, electrochemical analysis, and theoretical calculations, the key factors that contribute to the improved electrochemical performance of the lithium-rich materials are: (1) a reduced number of V_o results in a more stable structure, thereby stabilizing the lattice oxygen; and (2) modulating the nickel content causes an increase of the potential for redox reactions and remarkably enhances the oxygen stability under high voltage. The designed cathode exhibits a high initial coulombic efficiency of 83%. And a specific capability of 125 mA h g⁻¹ can be reached at 5C with a capacity retention of 97% even after 200 cycles. This strategy of modulating the Ni content of LLOs has great potential for designing high-energy-density Li-ion batteries.

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1. Introduction

Lithium-ion batteries (LIBs) are one of the most commonly used energy storage systems that have been widely used in portable electronic devices such as mobile phones, notebook computers, and electric or hybrid vehicles.^1,2 As the core component of LIBs, cathode materials control the capacity and operation voltage of the batteries and thus determine their energy density.³ To satisfy the high standards of next-generation LIBs, the development of high specific capacity cathode materials has become the top priority.^4,5 Lithium-rich layered oxides (LLOs) have been reported to deliver a specific charge capacity of about 250 mA h g⁻¹ and a higher operating voltage of about 3.7 V when they are charged to over 4.5 V at room temperature, which are promising cathode materials for next-generation LIBs with high energy density.^6–8 However, the LLO cathodes always suffer from pernicious voltage and capacity decay, and low initial coulombic efficiency, which come from unstable lattice oxygen and the migration or dissolution of transition metal (TM) ions during charge and discharge processes.^9,10

To overcome these problems, many strategies, such as doping^11–15 and coating,^16–22 have been attempted. And in recent years, composition adjustment, especially increasing the Ni content, has also been found to be useful. A series of studies have been conducted to investigate the impact of Ni content on the operating voltage and cycling stability properties of the LLO cathodes.^2,23,24 Sun et al. adjusted the chemical composition by increasing the Ni content to regulate the local electronic structure, which leads to a higher operating voltage and improved electrochemical cycling performance of the LLO cathode.²³ Wu et al. demonstrated that high Ni content contributes to lowering the amount of the Li₂MnO₃ phase, which affects both the reaction mechanism and the kinetics of the LLO cathodes. And as a result, the LLO cathode shows improved electrochemical stability and suppressed voltage decay during cycling.² These studies pay more attention to the influence of the Ni content on the transition metal, chemical composition, local electronic structure, and anionic redox properties of LLOs. However, the effect of a high Ni content on the defect structure, especially oxygen vacancies of LLOs, has not been clarified up to now.

Recently, some studies have shown that fewer oxygen vacancies (V_o) are beneficial in improving the structural stability of LLO cathodes during cycling.²⁵ V_o are products of lattice oxygen loss during cycling, which have been found to have a strong correlation with cation mixing,²⁶ transition metal ion migration,²⁷ and phase degradation from layered to spinel and finally to rock-salt phase,²⁸ and thus have adverse effects on structural stability and electrochemical performance. Cui et al. showed that a swing-like non-isothermal sintering technique is effective for reducing the number of intrinsic oxygen vacancies of the Li-rich Mn-based oxides and as a result, their electrochemical performance is improved obviously.²⁹ Xu et al. successfully synthesized a Co-free spinel LiNi_0.5−xMn_1.5+xO₄ (LNMO) using a high-temperature-shock (HTS) method. And the cathode material delivers excellent electrochemical properties because of the fewer surface V_o.³⁰ In short, fewer intrinsic V_o in the LLO cathodes is beneficial for enhancing their structural stability during cycling because of the stable lattice oxygen. And it is commonly believed that the sintering conditions, such as temperature, time, and atmosphere, have an important effect on V_o concentration.

In this work, two kinds of LLO materials with different Ni contents, Li_1.2Ni_0.2Mn_0.6O₂ (LR) and Li_1.13Ni_0.305Mn_0.565O₂ (HLR), were synthesized. And the effect of higher Ni content on the intrinsic V_o and electrochemical properties of the LLO materials was investigated. It was found that under the same sintering conditions, there exists a relatively lower number of intrinsic V_o in the LLO with higher Ni content. And as a result, the HLR material with a more stable lattice oxygen shows reversible redox reactions and structural stability during cycling. This work reveals the effect of higher Ni content on the properties of LLO materials from another aspect and also provides a method to regulate the defect structure of the LLO material.

2. Results and discussion

The crystal structures of LR and HLR were examined using powder X-ray diffraction (XRD) spectra, as shown in Fig. 1a and the corresponding Rietveld refinements were conducted using the GSAS program as shown in Fig. S1 and 2.† The main characteristic peaks of the two samples can be indexed as the standard α-NaFeO₂ hexagonal-layered structure with the R3̄m space group. The peaks in the 2θ range of 20–25° are distinct, demonstrating the formation of the super-lattice of monoclinic Li₂MnO₃ with C2/m symmetry.^31,32 With a higher Ni content, the intensity of super-lattice reflections is weaker in HLR, which indicates a decreased Li₂MnO₃ character in HLR. The fractions of the Li₂MnO₃ component in LR and HLR are about 34.5% and 23.2%, respectively, as obtained from Rietveld refinement of the XRD patterns (Fig. S1 and S2 in the ESI†). The evident splitting peaks of the (006)/(102) and (018)/(110) planes are observed, indicating a layered structure with high crystallization.³³ The values of I(003)/I(104) for LR and HLR are 1.508, and 1.582, respectively, which are larger than 1.2, evidencing their good layered structure.³⁴ Meanwhile, the refined results show that the degree of Li⁺/Ni²⁺ mixing is 3.99% and 5.12% for LR and HLR, respectively. Recent research has shown that appropriate cation disordering can support Li slabs and reduce the neighboring oxygen layer repulsion, which may maintain the layer structure during charging/discharging.³⁵ At the same time, cation disordering can inhibit the continuous migration of TM ions, which mitigates the phase transition from the layered to spinel phase during cycling. As the Ni content increases, the lattice parameter c and the unit cell volume V of the α-NaFeO₂ structure increase as shown in Table S1 in the ESI.† Oxygen atom occupation increases from 0.9238 (LR) to 0.9578 (HLR), indicating the decreased V_o in HLR.³⁶

Fig. 1 Structural characterization of LR and HLR: (a) XRD patterns; (b) Raman spectra; (c) survey XPS spectra; high resolution XPS spectra of (d) Ni 2p and (e) Mn 2p; (f) electron paramagnetic resonance (EPR) spectra; high resolution XPS results of O 1s at surface and depths in 50 nm of (g) LR sample and (h) HLR sample; (i) the percentage composition of lattice oxygen/oxygen vacancies at 50 nm etched depths and surfaces of LR and HLR samples.

To further investigate the local structures of the two samples, we measured the Raman spectra of the LR and HLR powders, as shown in Fig. 1b. For LR and HLR, the Raman spectrum contains two broad peaks at around 485 and 600 cm⁻¹, which can be assigned to the bending E_g and stretching A_1g modes in LiTMO₂.³⁷ The weak peak at ∼415 cm⁻¹ coming from the monoclinic Li₂MnO₃ component becomes weak for HLR, suggesting that the higher Ni content can suppress the formation of the Li₂MnO₃ component, which is consistent with XRD refinement results. Compared with the HLR sample, the Raman peaks of the LR sample show a slight red shift, indicating the weakening of M–O bonds which may be caused by the formation of oxygen vacancies, inferring that there are more oxygen vacancies in the LR sample.²⁹

We further obtained the X-ray photoelectron spectra (XPS) to investigate the valence of ions and the surface chemistry of HLR and LR, as shown in Fig. 1c–e. For both samples, the main Ni 2p3/2 peak located at about 855.2 eV can be resolved into two regions, 854.7 eV and 855.8 eV, corresponding to Ni²⁺ and Ni³⁺, respectively. The area of Ni³⁺ for HLR is larger than that for LR, indicating that the average valence state of Ni increases in HLR. In the meantime, Fig. 1e illustrates that the Mn 2p3/2 peak of both samples can be assigned to Mn³⁺ at 641.8 eV and Mn⁴⁺ at 642.7 eV.^38,39 However, the area ratio of the Mn³⁺ peak in LR is slightly larger than that of HLR. This result is most likely related to the significant removal of oxygen from the lattice.^29,30 Reducing the amount of Mn³⁺ is beneficial for mitigating the Jahn–Teller effect of high-spin Mn³⁺. In addition, for the O 1s spectrum, one peak located at 529.5 eV is attributed to the lattice oxygen (O–TM–O) and that at 531.4 eV belongs to V_o. Based on the peak areas, we can semi-quantitatively analyze the contents of V_o and lattice oxygen at the surface for the two samples.⁴⁰ Additionally, the XPS information after etching 50 nm is collected to identify the V_o in bulk. As shown in Fig. 1g–i, the relative content of V_o for LR is larger than that for HLR, both on the surface and after etching 50 nm.

The difference in the V_o concentration between the LR and HLR samples was also confirmed using EPR, as shown in Fig. 1f. It is worth noting that the responses in EPR include all possible V_o in the bulk and surface regions. The EPR response at a g factor of approximately 2.003 indicates the presence of V_o.⁴¹ Notably, the EPR response of HLR is lower than that of LR, indicating the relatively lower V_o concentration of HLR. Based on all the above analyses, it is inferred that the higher Ni content contributes to the lower V_o concentration of the LLO cathode material.

To further understand the structure of HLR and LR, we performed aberration-corrected high-angle annular-dark-field scanning transmission electron microscopy (HAADF-STEM), TEM and SEM for the two samples. The morphologies of LR and HLR and the enlarged images of the two materials are shown in Fig. 2a, d and Fig. S3, S4 in the ESI.† These images clearly indicate that the secondary particles of HLR and LR consist of close-packed primary particles and that the morphology of LLO did not change with the increase of nickel content. TEM images (Fig. 2b and e) reveal that the bulk and surface of the two samples have the same layered structure, matching well with the (003) plane of the R3̄m space. Energy dispersive spectroscopy (EDS) mappings (Fig. S5 in the ESI†) indicate that the O and TM elements are evenly distributed in the two materials. HAADF images of the two materials were used to further observe the cation mixing phenomenon. The HAADF imaging intensity of each atomic column reflects the average atomic number of each atomic column (∼Z 1.5 to Z 1.8) which is therefore termed the Z-contrast image and is sensitive to heavy elements. Hence the bright atomic columns correspond to the TM atoms, and the lighter elements Li and O are not visible in the image due to their low atomic weights. In Fig. 2c, little Ni/Li ion exchange is evident in LR and no TM ions are observed in the Li slab. In comparison, bright spots are observed in the Li layers of HLR, indicating more Li/Ni mixing in the sample with a higher Ni content, analogous to the XRD refined results.

Fig. 2 (a) SEM image of LR; (b) TEM images with fast Fourier transform (FFT) of LR; (c) HAADF image of LR; (d) SEM image of HLR; (e) TEM images with fast Fourier transform (FFT) of HLR; (f) HAADF image of HLR.

The electrochemical performances of LR and HLR are compared in Fig. 3. All cells are electrochemically investigated in the voltage range of 2.0–4.8 V at 20 °C. The initial charge and discharge profiles of the two samples at 0.1C are displayed in Fig. 3a. It is obvious that all the charging profiles consist of two parts: a ramp below 4.5 V and a long platform at 4.5 V. The charging reaction below 4.5 V attributed to the oxidation of TM (mainly Ni) in LiTMO₂ ⁴² is prolonged from LR to HLR due to the increase of the LiTMO₂ component. The platform region at 4.5 V related to the activation of the monoclinic Li₂MnO₃ component and O₂ release shortens along with the Ni content increase owing to the decreased Li₂MnO₃ component. However, HLR exhibits a high initial coulombic efficiency of 83%, about 1.1 times higher than that of the LR electrode. We conducted in situ differential electrochemical mass spectrometry (DEMS) tests on the two materials respectively during the first cycle to monitor lattice oxygen loss behavior and interfacial side reactions. As shown in Fig. 3b and c, when LR is charged above 4.5 V, an obvious signal peak of O₂ can be detected, as well as that of CO₂. In comparison, HLR releases less CO₂. The release of CO₂ is related to the decomposition reaction of the carbonate electrolyte. In general, high Ni content can increase the side reaction between Ni⁴⁺ and the electrolyte, thus leading to the decomposition of the electrolyte. The relatively lower CO₂ release in this work may be related to the reduction of oxygen free radicals generated by lattice oxygen loss,⁴³ because oxygen free radicals can also cause electrolyte decomposition. One can find that the O₂ release in HLR is undetectable (Fig. 3b and c). The release of O₂ is related to the loss of lattice oxygen. Therefore, it can be concluded that the lattice oxygen of HLR is more stable, which is related to the decreased Li₂MnO₃ phase and fewer oxygen vacancies caused by the increase of Ni content.

Fig. 3 (a) Initial charge–discharge profiles of the two samples in a voltage range of 2.0–4.8 V at 0.1C; in situ DEMS results of (b) LR, and (c) HLR during the first cycle; (d) rate capabilities of the materials from 0.1 C to 5 C (e) GITT curves during the first charging–discharging process (f) relationship between Li⁺ ion diffusion coefficient and potential. (g) Cycling performances of the two samples at 0.5C; dQ/dV of (h) LR and (i) HLR obtained from the discharge curves every ten cycles in 200 cycles.

The galvanostatic intermittent titration technique (GITT) test was used to assess the changes in the Li⁺ diffusion coefficient (D_Li⁺) during the discharge/charge process. The calculation process of D_Li⁺ is depicted in Supplementary Note 1 in the ESI† and the results are shown in Fig. 3e and f. We can clearly see that the D_Li⁺ value of HLR is evidently higher than that of LR whether during charging or discharging, indicating the faster Li⁺ diffusion in HLR. As a result, the improved Li⁺ ion transport property will contribute to the excellent rate performance of HLR. The rate performance further highlights the advantage of HLR, especially under high current densities such as 1, 2, and 5C, as shown in Fig. 3d.

Fig. 3g shows the cycling performance of the two samples at 0.5C. Compared with the LR sample (78.4% remaining after 200 cycles), the retained discharge capacity of the HLR sample was still over 198.3 mA h g⁻¹ after 200 cycles, with a high capacity retention rate of 97%. As illustrated in Fig. S6 in the ESI,† the average discharge voltages of the HLR and LR cathodes decreased by 191 and 268 mV, respectively, during 200 cycles. The corresponding dQ/dV curves of the two samples taken at various cycles were analyzed to investigate the evolution of the discharge voltage (Fig. 3h and i). These dQ/dV curves of the discharge process present three broad peaks, 1, 2 and 3, which are related to the reduction of Oⁿ⁻/O²⁻, Ni⁴⁺/Ni³⁺/Ni²⁺ and Mn⁴⁺/Mn³⁺ respectively.^44,45 In Fig. 3h, as the number of cycles increases, the reduction peaks of O, Mn and Ni all move to a lower voltage for LR, indicating that the LR sample has experienced pronounced capacity and voltage decay which is caused by the formation of the spinel phase.⁴⁶ The voltage position of the 2 peaks in LR decreases by 259.5 mV while that in HLR decreases by only 72.9 mV. This implies that both the capacity and voltage attenuation are effectively suppressed in HLR and the transition of the layered-to-spinel phase has been suppressed to some extent. It is concluded that, as the Ni content increases, electrochemical activation of the Li₂MnO₃ component is restrained, resulting in mitigated voltage decay.⁴⁷ Therefore, the excellent cycling stability of HLR is importantly attributed to a more stable lattice oxygen which can effectively suppress cation migration and the formation of microstructural defects during the electrochemical process.

Density functional theory (DFT) calculations were performed to gain a deeper understanding of how Ni content affects V_o formation and the electrochemical properties of LR and HLR cathodes. To simplify the structure, we excluded the scenario where TM ions occupy Li sites in the Li layer. The V_o formation energy of HLR was found to be 3.024 eV, higher than that of LR (2.129 eV). This shows that LR materials are more likely to form V_o under the same conditions, andthat HLR has the advantage of possessing more stable lattice oxygen, which effectively suppresses cationic migration, layered-to-spinel conversion, and the formation of microstructural defects during the electrochemical process. To further investigate the diffusion kinetics of Li⁺, nudged elastic band (NEB) calculations were performed to determine the Li⁺ diffusion energy barrier. As shown in Fig. 4b, the Li⁺ diffusion energy barrier in HLR is 0.683 eV lower than that in LR. A higher energy barrier indicates poorer Li⁺ diffusion kinetics. Therefore, HLR demonstrates superior kinetic performance during cycling, which aligns with the measured rate performance in Fig. 3d.

Fig. 4 (a) Oxygen vacancy formation energy of LR and HLR. (b) The diffusion energy barrier of Li⁺ and the illustration shows the Li⁺ migration path, LR on the left and HLR on the right. (c) Density of states of O 2p states and TM 3d states in LR and HLR sample. (d) Schematic diagrams of the electronic structure in LR (gray) and HLR (orange) samples according to the above density of states.

Fig. 4c depicts the density of states for the O 2p and TM 3d states in both samples. The O 2p states consist of two possible oxygen coordination environments: non-bonding (NB) states that contribute to anionic redox reactions, and higher-energy O 2p orbitals involved in TM–O bond formation. As shown in Fig. 4c, compared to LR, the non-bonding O 2p band in HLR shifts to a lower energy (approximately 0.22 eV) which can enhance remarkably the oxygen stability under high voltage. Furthermore, the TM 3d–O 2p state in HLR is lower than that in LR. The redox voltage is determined by the energy state of TM 3d and O 2p electrons involved in the redox reaction, and those with a lower energy state will show a higher redox reaction potential. As a result, HLR demonstrates a higher operational voltage compared to LR. Additionally, the energy-reduced TM 3d–O 2p bands can also mitigate the variation of the O 2p band throughout cycling, thereby contributing to improved cycling stability.⁴⁸

To investigate the effects of manipulating the Ni content on the crystal structural evolution, we performed in situ XRD during the initial three charging–discharging cycle processes, aiming to obtain more insight into the origin of the superior electrochemical performance of HLR. During the first cycle, the (003) peak of LR and HLR slowly shifts to a low angle before the voltage plateau at 4.5 V upon charging. And then it remains almost unchanged from 4.5 V to 4.8 V. At the same time, the (101) peak shifts to a high angle before 4.5 V, implying the decrease of lattice parameter a coming from the oxidation of TM ions followed by the decrease of ionic radii.⁴⁹ And then it basically remains unchanged at 4.5 V. In the process of discharge, the (003) peak initially shifts towards lower angles and then slightly moves to higher angles, indicating an increase in the lattice parameter c due to the intercalation of Li⁺ ions. Subsequently, it decreases slightly due to reduced repulsion between adjacent oxygen layers caused by continuous intercalation of Li⁺ ions. For the second cycle, both peaks show almost the same trend as the first cycle, except that the peaks change continuously rather than remaining constant at a voltage higher than 4.5 V. Comparing Fig. 5a and b, the difference is that the unit cell volume change of HLR is 1.88% which is much smaller than that of LR (2.87%) after three cycles, proving the better structural stability of the former during Li⁺ (de)intercalation processes.

Fig. 5 The in situ XRD results of (a) LR and (b) HLR electrodes during the initial three cycles. The potential profiles and corresponding color-coded images of the (003) and (101) peaks and (003) peak's 3D pictures are shown to clearly present the structural evolution. Changes of lattice parameters a and unit cell volume during the initial three cycles of LR and HLR.

During cycling, the (003) peaks of the two samples exhibit fluctuations in peak intensity and position, as revealed by the 3D color-coded contour maps. In the charging stage, the extraction of lithium leads to a decrease in crystallinity, resulting in a gradual weakening of peak strength. Conversely, during the discharging stage, the insertion of lithium ions causes a gradual increase in peak strength. It is evident that the (003) peaks of HLR demonstrate superior reversibility, indicating more reversible electrochemical reactions and enhanced structural stability. These can be attributed to the fact that the stable lattice oxygen in HLR inhibits the release of lattice oxygen and cationic migration, and layered-to-spinel conversion, resulting in an improvement of crystal structure stability.

Microstructure changes of HLR and LR after 200 cycles at 0.5C were studied by XRD. In Fig. 6a, the peaks at the 2θ region 10–20° indexed to the monoclinic Li₂MnO₃ distortion disappear after cycling. All the peaks shift to lower 2θ angles after 200 cycles, which is due to the structural rearrangements occurring upon cycling. But the shift of HLR is smaller than that of LR (Fig. S9 in the ESI†), further implying the structural stability of HLR during cycling. In addition, the I(003)/I(104) of HLR is 1.50, which is larger than that of LR (1.279), suggesting a minor level of cation mixing disorder of the former. As a consequence, some of the Li⁺ cannot insert into the former location and cation mixing does not occur, which leads to irreversible capacity loss.

Fig. 6 (a) XRD patterns of LR and HLR after 200 cycles. High-resolution XPS spectra of (b) Ni 2p and (c) O 1s in the fully discharged state after 200 cycles; discharge-state ex situ TEM and HRTEM images with FFT of (d) LR and (e) HLR samples after 200 cycles.

To further study the changes after cycling, high resolution XPS curves of LR and HLR electrodes at the end of the 200th discharge were obtained. Before the test, all the cycled electrodes were washed with dimethyl carbonate (DMC) and dried. From the C 1s and O 1s spectra (Fig. S11 in the ESI† and Fig. 6c), the peak intensities of C–O/C=O (286.4/288.6 eV), O–Fx (533.6 eV) and O=C–O (532.3 eV) of HLR, corresponding to the by-products of electrolyte decomposition, are weaker compared with LR, revealing that electrolyte decomposition is significantly suppressed for the former.^50–52 In the Ni 2p spectrum (Fig. 6b), both samples show peaks at about 858 eV, corresponding to NiF₂ which comes from the reaction between harmful HF and the cathode material.⁵³ The NiF₂ peak observed in HLR is significantly decreased compared to that of LR, suggesting the successful suppression of electrolyte decomposition and HF corrosion reactions. By analyzing the Mn 2p_3/2 spectra in Fig. S10 in the ESI,† LR includes more Mn³⁺ than HLR, caused by the loss of excess oxygen.²⁹ It is well known that Mn³⁺ is prone to disproportionation, which results in manganese dissolution. Conversely, the relatively high Mn⁴⁺ content in HLR indicates a favorable interface and superior structural stability during cycling.

To further evaluate the structural stability of LR and HLR, HRTEM analysis (Fig. 6d and e) was conducted to examine the internal and surface structures of the discharge-state LR and HLR samples after 200 cycles at 0.5C. The HLR sample exhibits a thin rock salt phase and limited spinel phases in the surface area, while the bulk phase largely maintains an intact layered structure. In contrast, the LR sample undergoes significant phase conversion in both the surface and bulk regions. Specifically, the lattice fringe of 4.7 Å in region I (marked by a yellow dotted box) corresponds well to the (003) plane of the layered structure.²⁹ Conversely, region II (marked by a red dotted oval) reveals a rhombic structure with a lattice fringe of 2.4 Å, indicating the presence of the spinel phase.⁵⁴ Notably, upon Li ion removal, TM ions predominantly migrate to neighboring Li sites with oxygen vacancies. The continuous migration of TM ions and irreversible degradation of the structure are key factors contributing to the inferior cyclic performance observed in the LR sample.

3. Conclusion

In this study, the effect of Ni content on the intrinsic V_o and electrochemical properties of the LLO was investigated. Increasing the Ni content in LLO helps to reduce the number of intrinsic V_o, thereby stabilizing the lattice oxygen. Theoretical calculations prove that a higher Ni content in LLO can enhance the potential for redox reactions, decrease the activation energy for Li⁺ diffusion, and increase the energy required for oxygen vacancy formation. Combining the experimental results, it can be concluded that the HLR material with a higher Ni content shows reversible redox reactions and the suppressed loss of lattice oxygen. Consequently, this mitigates TM reduction and prevents the irreversible transition from the layered structure to the spinel-like phase. And as a result, the HLR material exhibits improved cycling stability and lower voltage decay. This work gives an understanding of how higher nickel content affects the properties of LLO materials from another perspective.

Data availability

The data supporting this study's findings are available from the corresponding author upon reasonable request.

Conflicts of interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4qi01720j

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