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• The continuous expanding application field of lithium-ion batteries (LIBs) brings about an urgent demand for high-energy-density and long-life LIBs [1-3]. Recently, lithium rich layered oxide materials (LRLOs) have received significant attention due to their large capacity and high redox potential. LRLOs can be formularized as xLi₂MnO₃·(1 − x)LiMO₂ (0 < x < 1), where Li₂MnO₃ is a special LiMO₂ transition metal (TM) layered oxide with one-third of the TM replaced by Li [4-6]. The ultra-high capacity of LRLOs results from the extra Li⁺ sited in Li₂MnO₃ ingredient, which can be activated with lattice oxygen oxidation when charging above 4.5 V [7-9]. However, the removal of Li⁺ in TM layer of Li₂MnO₃ is laborious in dynamics [10], and excessive Li⁺ deficiency easily suffers structural destruction, such as TM migration [11] and oxygen dimerization (O²⁻ → Oⁿ⁻) [12], leading to poor rate performance and serious LRLOs degeneration both in capacity and voltage [13]. As mentioned, Li₂MnO₃ is a pivotal component and plays as the main "contributor" in LRLOs, but Li₂MnO₃ itself has bad electrochemical activity and faces severe electrochemical deterioration [14].

  Although many strategies, such as atomic doping [15-17], atomic mixing [18, 19] and surface coating [20, 21], are effective for LRLOs and other cathode materials to keep their high energy densities during the cycling, it is still hard to greatly improve the electrochemical performance of monoclinic Li₂MnO₃. Many efforts have been reported to deeply understand the decay mechanism of Li₂MnO₃, introducing that the primary reasons are oxygen release and subsequent structural destruction caused by the excessive oxygen oxidation at high voltages [22, 23]. So far, unmodified Li₂MnO₃ cathode has bad capacity retention even within 30 cycles, and it can only operate at a very low current density below 20 mA/g [24-27]. Xia's group used Li₂RuO₃ as a model and designed a Li/Ru disorganized TM layer, which restrained oxygen release and alleviated the cathode deterioration, providing a new strategy to improve the stability of Li₂MnO₃ [19]. Yabuuchi's group prepared a xLi₂TiO₃(1 − x)LiMnO₂ binary system, which could deliver a reversible capacity of ~250 mAh/g even at a high x value of 0.75, indicating that titanium anchored in Li₂MO₃ structure could improve its electrochemical properties [28]. Ti⁴⁺ can replace Mn⁴⁺ and Li⁺ in the TM layer and help improve the capacity retention of Li₂MnO₃ cathode because of the same octahedral configuration and electrochemical inertness of Ti⁴⁺ at high voltages [29-32]. In addition to Ti⁴⁺-substitution, another effective strategy to improve electrochemical performance is to introduce vacancies in cathodes, which can optimize the electronic properties of host materials, facilitating charge transfer and redox kinetics for Li-ion storage [33, 34]. In the case of cation vacancies, it can also serve as extra host sites for Li-ions insertion, promoting Li-ion diffusion kinetics during charge/discharge processes [35].

  In this work, we designed Li₂MnO₃ cathodes with Ti⁴⁺ substitution for Mn⁴⁺/Li⁺ in TM layer accompanied by vacancies in lithium layer to improve the electrochemical performance. According to the density functional theory (DFT) calculation, Ti⁴⁺ substitution for Li⁺ in the TM layer results in partial interlayer Li loss due to the unequal chemical valence of Ti⁴⁺ and Li⁺, simultaneously improving the kinetic activity of Li-ion diffusion. The designed Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ cathode shows high capacity, stable cycling performance and enhanced rate capability, which delivers a capacity of 290 mAh/g (229 mAh/g for Li₂MnO₃) at 10 mA/g and a capacity of 242 mAh/g (167 mAh/g for Li₂MnO₃) with capacity retention of 84% (39% for Li₂MnO₃) after 60 cycles at 50 mA/g. The improved electrochemical performance of Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ cathode benefits from the unequal-valent substitution of Ti⁴⁺ for Mn⁴⁺/Li⁺ and the associated interlayer Li vacancies, which inhibits the excessive oxidation of O²⁻ above 4.5 V, promotes Li-ion diffusion kinetics and activates the reversible redox reaction of Mn^3+/4+. The unequal-valent substitution of Ti⁴⁺ is an effective strategy to improve the electrochemical performance of Li₂MnO₃.

  The structure of Li₂MnO₃ originates from α-NaFeO₂ and presents monoclinic structure with space group C2/m [36]. As shown in Fig. 1a (left), the fully ordered Li₂MnO₃ can be regarded as a TM layered oxide with the chemical formula of Li[Li_1/3Mn_2/3]O₂, and the ordered TM slab is composed of one-third [Li_2bO₆] and two-thirds [Mn_4gO₆] octahedrons. We follow the design rules described in Fig. 1a. Ti⁴⁺ can efficiently replace Mn⁴⁺ on the 4g site because of the similar radius and same valence state of Ti⁴⁺ and Mn⁴⁺ ions. [Ti_4gO₆] can inhibit the excessive oxidation of O²⁻ above 4.5 V and improve the cyclic stability of Li₂MnO₃ structure due to the electrochemical inertness of Ti⁴⁺ at high voltage. Based on equivalent state substitution Li₂Mn_1-xTi_xO₃, we can adjust the ratio of lithium and titanium sources to make Ti⁴⁺ substitute Li⁺ on the 2b site and form Li vacancies in lithium layer, represented by a general stoichiometry of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂. The formation of Li vacancies is caused by the unequal-valent substitution between Ti⁴⁺ and Li⁺. And the amount of Li vacancies grows with the increase of the unequal-valent Ti⁴⁺-substitution. Introduction of Li vacancy will promote the Li-ion diffusion and activate the reversible redox reaction of Mn^3+/4+ during charging/discharging processes, leading to high Li⁺ storage ability.

  Figure 1

  

  Figure 1.  Design of unequal-valent Ti⁴⁺-substitution Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ with interlayer Li vacancies. (a) The atomic structure evolution from fully ordered Li₂MnO₃ (left) to equal-valent Ti⁴⁺ substituted Li₂Mn_1-xTi_xO₃ (middle), and then to unequal-valent Ti⁴⁺ substituted Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ with interlayer vacancies (right). (b) The theoretical calculation results of vacancy (Li) formation energy in Li₂MnO₃.

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  However, Li vacancy probably emerges at three sites of 2b, 2c and 4h. The theoretical calculation for each lithium-absence-state indicates that Li vacancy between TM layer is determined by the energetically favorable configuration. According to the simulation results, Li-ions would be preferentially extracted out from lithium layer. For the case of the first Li-ion extraction, the energetically favorable sites are 4h and 2c positions for both pure Li₂MnO₃ and Ti⁴⁺ substituted Li₂MnO₃. The formation of a Li vacancy in the LiMn₂ layer (TM layer) requires 0.18 eV higher energy than that in the lithium layer (Fig. 1b). It suggests that Li⁺ preferentially anchored at 2b site in TM layer in crystallization process, emerging Li vacancies mainly at 2c and 4h sites in lithium layer [37].

  The Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ can be easily synthesized by solid-phase reaction. Specifically, the prepared materials are Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, Li_0.5[Li_1/3 Mn_2/3]_0.5Ti_0.5O₂, which can be rewritten as Li_0.9[(Li_0.9Ti_0.1)_1/3 (Mn_0.9Ti_0.1)_2/3]O₂, Li_0.7[(Li_0.7Ti_0.3)_1/3(Mn_0.7Ti_0.3)_2/3]O₂, and Li_0.5[(Li_0.5Ti_0.5)_1/3(Mn_0.5Ti_0.5)_2/3]O₂, respectively. The inductively coupled plasma (ICP) results of these prepared materials are shown in Fig. 2a, which are in good agreement with their chemical formula. The particle morphology was characterized by scanning electron microscopy (SEM). As shown in Figs. 2b-e, all samples have the similar primary particle size of 2–3 µm via solid-phase synthesis. One difference observed from SEM images is that the primary particles of Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ locally agglomerate and form secondary particles (Fig. S1 in Supporting information).

  Figure 2

  

  Figure 2.  (a) The ICP results of the prepared Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5). The SEM images of (b) Li₂MnO₃, (c) Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, (d) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, and (e) Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂. The HAABF-STEM images along the [100] zone axis for (f) Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, (g) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ and [110] zone axis for (h) Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂. Yellow dashed boxes in (f, g) indicate the same nine layers in Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂. The ball represents atom in (f-h). Pink: Mn/Ti; Yellow: Li; Red: O.

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  To confirm the atomic arrangement of the above designed materials, we conducted high-angle annular bright field-scanning transmission electron microscopy (HAABF-STEM). Fig. 2f shows the HAABF-STEM image of Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ along the [100] zone axis. The atomic arrangement of Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ is similar to that of C2/m Li₂MnO₃ with a layer spacing of 4.8 Å and a repeating sequence of two-atom TM–TM dumbbells in TM layer [36]. A few (002) plane slip is caused by the partial absence of interlayer Li. In the HAABF-STEM image of Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ along the [100] zone axis (Fig. 2g), most regular domains with consecutive superlattice planes and the enclosed domains (red dotted lines) with continuous TM atoms coexist, indicating that the partial Li atoms at 2b sites are replaced by Mn or Ti atoms. In addition, the spacing distance of Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ expands to 4.9 Å. Fig. 2h displays a HAABF-STEM image of Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ along the [110] orientation. Interestingly, the atomic arrangement of Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ likes the structure of a spinel phase [38], which is quite different from the atomic arrangement of Li₂MnO₃. By the way, the Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ with less Li content can be written as Li[Li_1/3Mn_2/3Ti]O₄ resembling spinel-phase LiMn₂O₄.

  The structures of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5) were detected by X-ray diffractometer (XRD) and Raman spectroscopy measurements. Fig. 3a shows the XRD patterns of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂, and all strong diffraction peaks are similar to the layered α-NaFeO₂ structure. The 2θ around 20°−22° is attributed to the superlattice plane in monoclinic Li₂MnO₃, corresponding to (020) and (110) planes [39]. The superlattice peaks become weak with the augment of the x values, which reflects a growing tendency to shape a chaotic arrangement of atoms (Mn/Ti/Li) in TM layer. Finally, the superlattice diffraction peaks disappear in Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ with spinel structure. As shown in the enlarged XRD patterns at 18°−23° (Fig. 3b) and 35°−40° (Fig. 3c), the low-2θ-shift of (001) and (130) diffraction peaks corresponds to the increase of the interplanar spacing, which fits well with the HAABF-STEM images (Figs. 2f and g) and the XRD-rietveld results (Fig. S2 and Table S1 in Supporting information), indicating increased Li vacancies in Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂. Moreover, a new diffraction peak (indexed to (311) plane in spinel LiMn₂O₄ [40]) is observed in Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ (Fig. 3c), ascribed to the interlayer Li vacancies.

  Figure 3

  

  Figure 3.  (a-c) XRD patterns at different ranges, (d) Raman shifts and (e, f) high-resolution XPS spectra for Mn and Ti of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5).

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  Raman shift (Fig. 3d) was used to analyze the molecular structures of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂, which were difficult to be detected by XRD measurements. In monoclinic Li₂MnO₃, there are two strong stretching-vibration (A_1g and A_g) peaks and a distinct bending-vibration (E_g) peak [41]. Specifically, the A_g peak located at 425 cm⁻¹ stands for a characteristic vibration of monoclinic Li₂MnO₃ structure [42], while E_g and A_1g peaks at 487 cm⁻¹ and 605 cm⁻¹ are associated with the Raman-sensitive layered structure [43]. For spinel-phase Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂, the A_g vibration disappears and the A_1g peak shifts to 630 cm⁻¹, which is obviously different from the monoclinic Li₂MnO₃ structure [44]. For Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, a slight red shift of A_1g peak indicates that they tend to transform to spinel phase [45]. However, the stretching A_g vibrations still exist in their molecular structures, illustrating that they still contain the monoclinic structure (like Li₂MnO₃). The E_g peak gradually weakens because of the variations in the local vibration symmetry of [TMO₆] octahedron [46], which mainly generates by the Li vacancies and Ti⁴⁺ substitutions.

  X-ray photoelectron spectroscopy (XPS) was employed to identify valence states of Mn and Ti in these materials. As shown in Figs. 3e and f, both Mn and Ti present as tetravalence in Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, which indirectly proves the existence of Li vacancies. By the way, a right-shoulder peak visualizes at the Mn 2p_3/2 spectra of Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂, indicating the co-existence of Mn⁴⁺ and Mn³⁺ in spinel-phase Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂.

  The charge/discharge measurements of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5) were conducted to explore the influence of the unequal-valent Ti⁴⁺-substitution with interlayer Li vacancies on the electrochemical performance. Figs. 4a-d display the charge/discharge curves of the above cathode materials for the first two cycles at a current density of 10 mA/g. To facilitate comparison, their initial charge/discharge curves are placed together in Fig. S3 (Supporting information). The initial charge capacities of Li₂MnO₃ (247 mAh/g), Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ (231 mAh/g) and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ (205 mAh/g) are mainly contributed by the oxidation of O²⁻ above 4.5 V. With the increase of unequal-valent Ti⁴⁺-substitution, the decrease of the charging capacity is ascribed to the effort of Ti⁴⁺ substitution, which inhibits the lattice oxygen oxidation at high voltages. Meanwhile, the slightly increasing capacity between 3.5–4.5 V is caused by the oxidation of Mn³⁺. By contrast, the spinel-phase Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ cathode exhibits much lower first charge capacity (117 mAh/g) than other samples. The initial discharge capacities of Li₂MnO₃, Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, and Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ cathodes are 229, 241, 290 and 252 mAh/g, respectively. Compared with Li₂MnO₃, the Ti⁴⁺-substituted samples present a new plateau, located at around 2.8 V in Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, but 2.6 V in spinel-phase Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂. This unique lithiation process is related to the reduction of Mn⁴⁺ accompanying with the intercalation of Li⁺ [18], which benefits from interlayer Li vacancies. The plateau related to Mn^3+/4+ redox extends with the unequal-valent Ti⁴⁺ substitution, increasing until the phase transferred to spinel after Ti⁴⁺ substituted Li₂MnO₃. The long plateau means a deep reduction reaction of Mn^3+/4+, indicating that unequal-valent Ti⁴⁺-substitution and interlayer Li vacancies can increase the capacity of the sample. And the voltages associated with the increased capacities of Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ are completely retained in the second cycle, indicating good reversibility of Mn^3+/4+ redox. The corresponding differential capacity versus voltage (dQ/dV) plots of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5) exhibit similar electrochemical changes to their charge/discharge curves. Fig. 4e shows initial dQ/dV plots of charging process, the oxygen dimerization peak weakens with the increase of x value and finally disappears in spinel-phase Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂. The paired dQ/dV plots during discharge are shown in Fig. 4f, a sharp peak emerges around 2.8 V after unequal-valent Ti⁴⁺ substituting, which corresponds to the Mn^3+/4+ redox activated by Ti⁴⁺ substitution and interlayer Li vacancies. The sharpest reduction peak appears in Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, and the reduction peak shifts to 2.6 V in spinel-phase Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂.

  Figure 4

  

  Figure 4.  First-two cycles charge/discharge curves of (a) Li₂MnO₃, (b) Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, (c) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, and (d) Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ at 10 mA/g. The corresponding dQ/dV plots of (e) charging process and (f) discharging process related to the first cycle.

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  To further understand the contribution of interlayer Li vacancies caused by unequal-valent Ti⁴⁺-substitution to Li-ion storage, we compare the electrochemical characteristics of equal-valent Ti⁴⁺ substituted Li₂Mn_1-xTi_xO₃, including Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃, where Ti⁴⁺ only replaces Mn⁴⁺ without Li vacancy formation. The XRD patterns of Li₂Mn_1-xTi_xO₃ (x = 0, 0.1 and 0.3) (Fig. S4 in Supporting information) manifest high similar structure. Fig. S5 (Supporting information) shows their charge/discharge curves at 10 mA/g. The simply equal-valent Ti⁴⁺-substitution (Ti substituted Mn) can also activate Mn^3+/4+ redox, but reversibility is worse than that of unequal-valent Ti⁴⁺-substituted Li₂MnO₃, which is our designed strategy to achieve Ti⁴⁺-substitution and build interlayer Li vacancies. The Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃ show the lower Coulombic efficiency (CE) than that of Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ in the second cycle, indicating that interlayer Li vacancies can also improve the stability of electrochemical process. Compared with the designed unequal-valent Ti⁴⁺-substitution cathodes, there are no sharp peaks in the dQ/dV plots of equal-valent Ti⁴⁺ substituted Li₂Mn_1-xTi_xO₃ during the discharge process (Fig. S6 in Supporting infromation), demonstrating that the activation of Mn^3+/4+ redox is mainly contributed by the unequal substitution with interlayer Li vacancies.

  The cycling performance of Li₂MnO₃ based cathodes was measured at a current density of 50 mA/g. As shown in Fig. 5a, the initial capacities of Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂, Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, and Li₂MnO₃ cathodes are 220, 242, 143 and 167 mAh/g, respectively, and their capacity retention after 60 cycles are 59%, 83%, 84%, and 39%, respectively. Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃ cathodes have 139 and 143 mAh/g capacities with 47% and 61% retentions at 50 mA/g after 60 cycles (Fig. S7 in Supporting information). Among these cathodes, Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ cathode shows the highest CE in 60 cycles, especially in the first five cycles (Fig. S8 in Supporting information). In view of these results, the unequal-valent Ti⁴⁺-substitution with interlayer Li vacancies building and Ti⁴⁺-substitution can increase capacities and improve cyclic stability of Li₂MnO₃ based cathodes. In addition, Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ cathode delivers a high energy of 661 Wh/kg (457 Wh/kg for Li₂MnO₃) and a residual energy of 447 Wh/kg (144 Wh/kg for Li₂MnO₃) after 60 cycles (Fig. 5b and Fig. S9 in Supporting information). The rate ability of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5) cathodes also was investigated. The rate performance of Li₂MnO₃ cathode is enhanced with the appropriate unequal-valent Ti⁴⁺-substitution (Fig. 5c), implying that Ti⁴⁺ substitution and interlayer Li vacancies induce a thermodynamically fast Li-ion intercalation process in the Li₂MnO₃ structural phase.

  Figure 5

  

  Figure 5.  (a) Capacity, (b) energy density, and (c) rate performance of the Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5) at a current density of 50 mA/g.

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  DFT calculations (Fig. 6) in terms of structural relaxation and total energy were carried out to further understand the improved LIBs performance of the modified Li₂MnO₃ cathodes. The simplified model of Ti replacing Mn is adopted to show the effect of Ti⁴⁺ substitution in Li₂MnO₃. Calculation results show that the average length of the TM—O bond increased compared to the length of the Mn—O bond in Li₂MnO₃ after Ti⁴⁺ substitution (Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃) (Fig. 6a). O1-O6 stand for six TM—O bonds in [TMO₆] octahedron as shown in Fig. S10 (Supporting information). In addition, the substitution of Ti⁴⁺ pushes the surrounding atoms away from the [TiO₆] octahedron, which is consistent with the Mo doped Li₂MnO₃ configuration [47], indicating that the surrounding structure is loose. For the Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃ structures, the calculated supercell volume will increase from 413.5 ų to 418.0 ų and 417.7 ų, respectively, as listed in Table S2 (Supporting information). According to Gao et al. [47], the expanded lattice due to Ti doping can facilitate the Li-ion diffusion. The correlations between the delithiation energy and the lithium content under different substitution conditions are shown in Fig. 6b, where n represents the number of Li-ions extraction. With respect to the pristine Li₂MnO₃, the delithiation energy is reduced by 0.58–0.75 eV after the Ti⁴⁺ substitution, indicating that the electrochemical activity of Li₂MnO₃ is enhanced by the incorporation of Ti⁴⁺.

  Figure 6

  

  Figure 6.  (a) Calculated bond length of TM-O between oxygen and transition metals in Li₂MnO₃, Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃, respectively. (b) Calculated formation energies of Li vacancies in Li₂MnO₃, Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃. (c) Calculated reaction Gibbs free energy of O₂ release from the lattice of Li_2-xMn_1-yTi_yO₃, with x = 0, 0.125, 0.25, 0.375, 0.5, 0.625, 0.75, and y = 0, 0.1, 0.3.

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  Furthermore, we evaluate the O₂ release reaction to reflect the stability of the oxygen sublattice in Ti⁴⁺ substituted Li₂MnO₃ using DFT calculations. Considering the entropy of gas phase O₂ under standard states (−TΔS = −0.63 eV) [48], the Gibbs free energy ΔG can be obtained. Fig. 6c shows the reaction Gibbs free energy of O₂ release from the lattice of Li_2-xMn_1-yTi_yO₃. Upon delithiation, the Gibbs free energies decrease monotonically, indicating that the lattice stability of Li_2-xMn_1-yTi_yO₃ will decrease with increasing amount of delithiation. In the case of unsubstituted Li₂MnO₃, the formation Gibbs energy is below zero when x > 0.4, indicating that the oxygen release reaction would occur spontaneously. After Ti doping, the spontaneous oxygen release will delay to x > 0.5 and x > 0.6 for Li₂Mn_0.7Ti_0.3O₃ and Li₂Mn_0.9Ti_0.1O₃, respectively, suggesting that the O₂ release reaction is suppressed and consequently the structural stability is improved.

  We carried out ex-situ XRD measurement to detect the structural evolution of Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ and Li₂MnO₃ at open circuit (OC), full charge (FC) and full discharge (FD) states under a current density of 10 mA/g. The ex-situ XRD patterns are shown in Figs. 7a and b. At the first FC state, the diffraction peak of (001) slightly shifts to high degree (clearly shown in Fig. S11a in Supporting information) in Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ due to the lattice oxygen oxidation with Li⁺ removing from TM layer. On contrast, the (001) diffraction peak of Li₂MnO₃ shifts to low degree at initial FC state (Fig. S11b in Supporting information), indicating the enlarging of the interlayer spacing caused by the excessive lattice oxygen dimerization and the oxygen release, which represents an unstable structure during the charge process. At the initial FD state, the interlayer spacing of both Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ and Li₂MnO₃ becomes large due to Li⁺ intercalation, which is a typical lithiation phenomenon in layered oxide cathodes. The (001) diffraction peak of Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ moves back, which means that Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ with enlarged interlayer spacing can alleviate the strain perpendicular to z axis during Li⁺ intercalation. It avoids the lattice collapse in layered structure. Differently, the (130) diffraction peak of FD state in Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ disappears, which means that the Mn⁴⁺ reduction induced extra embedded Li⁺ are not arranged as regularly as Li⁺ in Li₂MnO₃. The extra Li⁺ intercalation exhibits a good reversible behavior as the (130) diffraction peak reappears in Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ electrode at the second FC state.

  Figure 7

  

  Figure 7.  Ex-situ XRD patterns of (a) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ and (b) Li₂MnO₃ cathodes. And ex-situ XPS spectra of the Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ electrode for (c) Mn and (d) Ti elements.

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  Furthermore, the ex-situ XPS tests were taken out to analyze the charge compensation mechanisms and identify the source of the increased capacity. Fig. 7c shows the ex-situ XPS spectra of Mn element in Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ at different states. Mn⁴⁺ is not oxidized to a high valence state during the charge process, while it is partially reduced to Mn³⁺ as expected with a Mn⁴⁺/Mn³⁺ ions ratio of 2.2 (closing to the Li/V_Li ratio of 0.7/0.3 in lithium layer) in the first FD state. It can further confirm that the increased capacity is originate from the Mn^3+/4+ redox reaction. Subsequently, almost all of the Mn³⁺ are oxidized to Mn⁴⁺ at the second FC state, representing a good reversibility of the Mn^3+/4+ redox. The ex-situ XPS spectra of Ti element in Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ are shown in Fig. 7d. Ti⁴⁺ always exists without charge transfer, further verifying the stable and inert of titanium in cathodes at high voltages, which is beneficial to improve the electrochemical stability of unequal-valent Ti⁴⁺-substituted Li₂MnO₃ cathodes.

  In summary, a series of unequal-valent Ti⁴⁺-substituted Li₂MnO₃ (Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂, x = 0, 0.1, 0.3 and 0.5) samples are successfully designed. After unequal-valent Ti⁴⁺-substitution, the interlayer Li vacancies are built. Our strategy can stabilize the structure and improve the cycling performance of Li₂MnO₃-based cathodes. In particular, the substitution of Ti⁴⁺ for Li⁺ in TM layer leads to the formation of Li vacancies in lithium layer, which is in good agreement with the DFT calculation results that the Li vacancies preferentially existed at 2c and 4h sites (lithium layer) in Li₂MnO₃. The interlayer Li vacancies facilitate the Li-ion transport and greatly increase the capacity by activating the reversible Mn^3+/4+ redox during the charge/discharge processes. In addition, DFT calculations simultaneously confirm that the Ti⁴⁺ substituted Li₂MnO₃ can promote Li-ion diffusion and inhibit the oxidation of O²⁻ above 4.5 V. The as-designed Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ cathode displays a high capacity of 290 mAh/g, stable cycling and enhanced rate performance. Therefore, the unequal-valent Ti⁴⁺-substitution is an efficient strategy for improving the electrochemical performance of Li₂MnO₃.

  Declaration of competing interest

  There are no conflicts to declare.

  Acknowledgments

  This work was financially supported by the National Natural Science Foundation of China (Nos. 51972258 and 22109186) and Open Fund by Sanya Science and Education Innovation Park of Wuhan University of Technology (No. 2021KF0021). Y. Sun is supported by 21C Innovation Laboratory, Contemporary Amperex Technology Ltd. by Project No. 21C-OP-202002.

  Supplementary materials

  Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.cclet.2022.05.008.

  
  
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• 

  Figure 1  Design of unequal-valent Ti⁴⁺-substitution Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ with interlayer Li vacancies. (a) The atomic structure evolution from fully ordered Li₂MnO₃ (left) to equal-valent Ti⁴⁺ substituted Li₂Mn_1-xTi_xO₃ (middle), and then to unequal-valent Ti⁴⁺ substituted Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ with interlayer vacancies (right). (b) The theoretical calculation results of vacancy (Li) formation energy in Li₂MnO₃.

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  Figure 2  (a) The ICP results of the prepared Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5). The SEM images of (b) Li₂MnO₃, (c) Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, (d) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, and (e) Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂. The HAABF-STEM images along the [100] zone axis for (f) Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, (g) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ and [110] zone axis for (h) Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂. Yellow dashed boxes in (f, g) indicate the same nine layers in Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂ and Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂. The ball represents atom in (f-h). Pink: Mn/Ti; Yellow: Li; Red: O.

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  Figure 3  (a-c) XRD patterns at different ranges, (d) Raman shifts and (e, f) high-resolution XPS spectra for Mn and Ti of Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5).

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  Figure 4  First-two cycles charge/discharge curves of (a) Li₂MnO₃, (b) Li_0.9[Li_1/3Mn_2/3]_0.9Ti_0.1O₂, (c) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂, and (d) Li_0.5[Li_1/3Mn_2/3]_0.5Ti_0.5O₂ at 10 mA/g. The corresponding dQ/dV plots of (e) charging process and (f) discharging process related to the first cycle.

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  Figure 5  (a) Capacity, (b) energy density, and (c) rate performance of the Li_1-x[Li_1/3Mn_2/3]_1-xTi_xO₂ (x = 0, 0.1, 0.3 and 0.5) at a current density of 50 mA/g.

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  Figure 6  (a) Calculated bond length of TM-O between oxygen and transition metals in Li₂MnO₃, Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃, respectively. (b) Calculated formation energies of Li vacancies in Li₂MnO₃, Li₂Mn_0.9Ti_0.1O₃ and Li₂Mn_0.7Ti_0.3O₃. (c) Calculated reaction Gibbs free energy of O₂ release from the lattice of Li_2-xMn_1-yTi_yO₃, with x = 0, 0.125, 0.25, 0.375, 0.5, 0.625, 0.75, and y = 0, 0.1, 0.3.

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  Figure 7  Ex-situ XRD patterns of (a) Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ and (b) Li₂MnO₃ cathodes. And ex-situ XPS spectra of the Li_0.7[Li_1/3Mn_2/3]_0.7Ti_0.3O₂ electrode for (c) Mn and (d) Ti elements.

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