Low-Cost Mn-Based Cathode Materials for Lithium-Ion Batteries
Low-Cost Mn-Based Cathode Materials for Lithium-Ion Batteries
Yi, Hongming;Liang, Ying;Qian, Yunlong;Feng, Yuchuan;Li, Zheng;Zhang, Xue
2023-04-26 00:00:00
batteries Review 1 , † 2 , † 1 1 , 2 1 1 , Hongming Yi , Ying Liang , Yunlong Qian , Yuchuan Feng , Zheng Li and Xue Zhang * QingTao (KunShan) Energy Development Co., Ltd., Shengxi Road No. 3, KET, Kunshan 215300, China State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China * Correspondence: zhangx_qingtao@163.com † These authors contributed equally in this work. Abstract: Due to a high energy density and satisfactory longevity, lithium-ion batteries (LIBs) have been widely applied in the fields of consumer electronics and electric vehicles. Cathodes, an essential part of LIBs, greatly determine the energy density and total cost of LIBs. In order to make LIBs more competitive, it is urgent to develop low-cost commercial cathode materials. Among all cathode materials, Mn-based cathode materials, such as layered LiNi Mn O and Li-rich materials, spinel 0.5 0.5 2 LiMn O and LiNi Mn O , olivine-type LiMnPO and LiMn Fe PO , stand out owing to their 2 4 0.5 1.5 4 4 0.5 0.5 4 low cost and high energy density. Herein, from the perspective of industrial application, we calculate the product cost of Mn-based cathode materials, select promising candidates with low cost per Wh, and summarize the structural and electrochemical properties and improvement strategies of these low-cost Mn-based cathode materials. Apart from some common issues for Mn-based cathode materials, such as Jahn–Teller distortions and Mn dissolution, we point out the specific problems of each material and provide corresponding improvement strategies to overcome these drawbacks. Keywords: lithium-ion batteries; Mn-based cathodes; low cost per Wh; high energy density 1. Introduction The widespread application of consumer electronics and the urgent requirements of electrical vehicles have greatly promoted the development of lithium-ion batteries Citation: Yi, H.; Liang, Y.; Qian, Y.; (LIBs) [1–3]. Among the main components of LIBs, such as the cathode, anode, electrolyte, Feng, Y.; Li, Z.; Zhang, X. Low-Cost and separator, cathode materials, one of the main components of LIBs along with an- Mn-Based Cathode Materials for ode materials, liquid electrolytes, and separators, are crucial to the whole energy density Lithium-Ion Batteries. Batteries 2023, and costing of LIBs. This is because of the relatively lower capacity of cathode materi- 9, 246. https://doi.org/10.3390/ als compared with that of anode materials (for example, commercial LiCoO cathode: batteries9050246 1 1 ~180 mAh g ; commercial graphite anode: ~350 mAh g ) and their non-negligibly higher Academic Editor: Seung-Wan Song cost (about 40–60% of total LIBs). The high cost of cathode materials has caused the high price of LIBs, which drives up the cost of LIB-based products, especially those that need Received: 7 December 2022 a large quantity of LIBs, such as electrical vehicles, energy storage devices and electrical Revised: 17 March 2023 tools. As lithium resources are limited and unevenly distributed, LIBs are becoming more Accepted: 23 April 2023 Published: 26 April 2023 and more expensive, which has resulted in the emergence of other competitive batteries, including sodium-ion batteries (SIBs), potassium-ion batteries (KIBs) and magnesium-ion batteries (MIBs). Meanwhile, low-cost lead-acid batteries (LAB) still occupy a large part of the battery market. Under the pressure of traditionally commercial batteries (LABs) and Copyright: © 2023 by the authors. emerging batteries (SIBs, KIBs, MIBs etc.), it is urgent to develop new inexpensive commer- Licensee MDPI, Basel, Switzerland. cial cathode materials for LIBs, so that LIBs can be more competitive in the battery market. This article is an open access article To screen low-cost cathode materials, it is necessary to take the cost into account from distributed under the terms and the elemental standpoint. At present, cobalt (Co), which is included in commercial LiCoO conditions of the Creative Commons and ternary materials (LiNi Co Mn O , NCM; LiNi Co Al O , NCA etc.), is the x y x y 1-x-y 2 1-x-y 2 Attribution (CC BY) license (https:// primary element of cathode materials. However, the price of Co is extremely high and keeps creativecommons.org/licenses/by/ rising due to its rare reserve and uneven distribution [4]. Another widely used element for 4.0/). Batteries 2023, 9, 246. https://doi.org/10.3390/batteries9050246 https://www.mdpi.com/journal/batteries Batteries 2023, 9, 246 2 of 23 cathode materials is Ferrum (Fe), which is included in commercial LiFePO for power bat- teries. However, Fe ions not only have a +2 and +3 valance state, which limits the diversity of Fe-based cathode materials, but also exhibit a low discharge voltage (3.5 V vs. Li/Li ), which greatly limits the discharge capacity of LiFePO [5]. By comparison, manganese (Mn) is plentiful in the Earth’s crust and has been utilized extensively in Ferrum and the steel industry, non-ferrous metallurgy, the chemical industry, electronics, batteries, agriculture, medicine and other fields. Additionally, Mn ions have manifold valance states ranging from +2 to +7, which enables them to produce various cathode materials with different types of crystal structures, such as spinel LiMn O , layered LiMnO and olivine LiMnPO . 2 4 2 4 Therefore, Mn-based materials should be the main emphasis when seeking next-generation low-cost cathode materials. Nowadays, LiMn O has become a commercial material by 2 4 virtue of its price advantage and is mainly applied in the two-wheel electric-vehicle indus- tries. Other Mn-based materials, including LiNi Mn O , LiNi Mn O , LiMnPO and 0.5 0.5 2 0.5 1.5 4 4 Li-rich Mn-based materials, will potentially be used as commercial cathode materials for LIBs because of their high energy density and inexpensive price. There exist several excellent reviews on Mn-based materials which normally focus on a specific type of Mn-based material, such as spinel LiMn O [6,7], layered LiMnO [8] 2 4 2 and Li-rich Mn-based materials [9,10], and collect recent advances in the academic area; however, they ignore practical industrial applications. Hence, the purpose of this review is to offer an essential understanding of the performance and preparation cost of various types of Mn-based materials. By calculating the preparation cost (raw materials and preparation technology), we selected potential low-cost cathode materials whose cost per Wh is lower than that of present NCM materials. The basic structures, electrochemical performance, and improvement strategies are also elucidated systematically, and their industrial-application prospects are also outlined. 2. Estimation of Energy Density and Cost for Mn-Based Materials For the industrial application of cathode materials in LIBs, the energy density (gravi- metric energy density, GED; volumetric energy density, VED) and cost (production cost, cost per Wh) are vital parameters. The theoretical specific capacity of a cathode material is computed according to the formula Q = nF/3.6 M , where n is the electron-transfer number, F is the Faraday constant, and M is the molecular weight of the cathode material [11]. In order to assess the realistic economic worth of the cathode materials, the data of practical specific capacities, average discharge voltages, and compaction densities of most materials (except LiNi Mn O Li Ni Mn O and Li Mn Nb O F ) are obtained 0.5 0.5 2, 1.2 0.2 0.6 2 1.2 0.625 0.175 1.95 0.05 from various suppliers of LIB’s materials. The material cost is computed according to the average price of raw materials on the CBCIE website (www.cbcie.com) in 2021 and the production cost. The raw materials include the transition-metal salts or transition-metal precursor, lithium salt, phosphate salts for phosphate materials, and additives. The pro- duction cost not only includes the cost of electricity, packaging and transportation, but also considers the cost of labor and depreciation of equipment. The GED is generated by multiplying practical specific capacity and average discharge voltage. The VED is generated by multiplying GED and compaction density. The final cost per Wh is the result of material cost/GED. Table 1 summarizes the results from the above. According to Table 1, LiCoO has the highest VED, implying its widespread appli- cation in consumer electrics that provide limited space for batteries. For electric vehicles, the GED and material cost are more important than VED due to its larger space and use of many more batteries to produce much more energy. Although NCM materials, such as LiNi Co Mn O (NCM523) and LiNi Co Mn O (NCM811) have a substantially 0.5 0.2 0.3 2 0.8 0.1 0.1 2 lower VED than LiCoO , their GED is comparable to that of LiCoO . NCM is widely used 2 2 in batteries of electrical vehicles due to its significantly cheaper material cost compared to LiCoO [12]. Although LiFePO has a relatively lower GED than NCM and LiCoO , 2 4 2 it is still a viable cathode material for electrical vehicles due to its low cost per Wh and –1 extraordinarily extended working life. The cost per Wh of LiFePO is 0.15 CNY Wh , 4 Batteries 2023, 9, 246 3 of 23 which is relatively low in present commercial materials, but its energy density is limited to a low level. This paper reviews the low-cost materials with lower cost per Wh than –1 –1 NCM (NCM523: 0.34 CNY Wh ; NCM811: 0.34 CNY Wh ). Then, we find that almost all Mn-based materials can meet the aforementioned requirement, and most of them have –1 prices of less than 0.20 CNY Wh , with the exception of Li Mn Nb O F . 1.2 0.625 0.175 1.95 0.05 Among these low-price Mn-based materials, spinel LiMn O and LiNi Mn O demon- 2 4 0.5 1.5 4 –1 strated the lowest cost per Wh at 0.13 CNY Wh . LiMn O has been applied in two-wheel 2 4 electric vehicles and electric tools that need higher cost performance and tolerance for relative lower energy density and cycle performance, and it has the potential to replace LABs. The layered LiNi Mn O has a comparable GED to NCM523 and its cost of Wh 0.5 0.5 2 (0.19 CNY Wh ) is much lower than that of NCM523, which has the potential to replace –1 NCM materials by virtue of its low price. The cost per Wh for LiMnPO (0.16 CNY Wh ) –1 and LiMn Fe PO (0.15 CNY Wh ) is similar to that of LiFePO , which has shown 0.5 0.5 4 4 a commendable advantage in price. Meanwhile, the higher GED of these two materials make them more attractive, and some material companies are developing them. Li-rich –1 –1 materials, such as layered Li Ni Mn O (0.19 CNY Wh , 825.0 Wh kg ) and rock-salt 1.2 0.2 0.6 2 –1 –1 Li Mn Nb O F (0.30 CNY Wh , 1056 Wh kg ) are becoming research-and- 1.2 0.625 0.175 1.95 0.05 development hotspots for many material companies. Although they do not have a low cost per Wh, they have the highest GEDs among all materials, which can greatly improve the range of electric vehicles. Batteries 2023, 9, 246 4 of 23 Table 1. The comparison of various Mn-based materials. Most data are obtained from the materials’ suppliers in China except LiNi Mn O Li Ni Mn O 0.5 0.5 2, 1.2 0.2 0.6 2 and Li Mn Nb O F . The currency type “CNY” is “Chinese Yuan”. 1.2 0.625 0.175 1.95 0.05 Practical Theoretical Average Specific Initial Compaction Material Specific Discharge Gravimetricenergydensity Volumetricenergydensity Cost per Wh Main Cycling Air Materials Capacity Coulombic Density Cost –1 –1 –1 Capacity Voltage (Wh kg ) (Wh L ) (CNY Wh ) Drawbacks Performance Stability –1 –3 –1 (mAh g ) Efficiency (g cm ) (CNY kg ) –1 (mAh g ) (V) @ 0.1C LiCoO 274 210 94% 3.80 3.80 798.0 3032 308.6 0.39 High cost FFFII FFFFF Relatively NCM523 276 180 88% 3.71 3.45 667.8 2304 226.4 0.34 low energy FFFFI FFFII density Poor stability, NCM811 275 210 91% 3.66 3.40 768.6 2613 235.2 0.31 FFFII FFIII bad safety Low energy LiFePO 170 160 98% 3.35 2.20 536.0 1179 81.1 0.15 FFFFF FFFFI density Low energy LiMn O 148 120 95% 3.80 3.00 456.0 1368 57.5 0.13 FFIII FFFFI 2 4 density Spinel High voltage LiNi Mn O 147 133 94% 4.70 3.00 625.1 1875 82.7 0.13 FFFII FFFFI 0.5 1.5 4 plateau LiNi Mn O Poor rate 0.5 0.5 2 280 199 96% 3.40 / 676.6 / 128.3 0.19 FFIII FFFII [13] capability Layered Low initial Li Ni Mn O 1.2 0.2 0.6 2 378 240 82% 3.45 2.80 845.0 2367 158.3 0.19 coulombic FFIII FFFFI [14] efficiency High Li Mn Nb 1.2 0.625 0.175 Rock salt 353 330 94% 3.20 2.70 1056 2851 319.2 0.30 production FIIII FFFFI O F [15] 1.95 0.05 cost Low LiMnPO 171 154 83% 3.90 / 600.6 / 96.5 0.16 electronic IIIII FFFFI conductivity Olivine Low energy LiMn Fe PO 170 160 93% 3.72 2.40 595.2 1428 89.1 0.15 FFFFF FFFFI 0.5 0.5 4 density Batteries 2023, 9, x FOR PEER REVIEW 5 of 23 Batteries 2023, 9, 246 5 of 23 3. Structure, Performance and Improvement Strategies of Mn-Based Materials 3. Structure, Performance and Improvement Strategies of Mn-Based Materials The typical low-cost Mn-based materials listed in Table 1 include layered, spinel, The typical low-cost Mn-based materials listed in Table 1 include layered, spinel, rock-salt and olivine structures. Most layered transition-metal oxides often undergo a rock-salt and olivine structures. Most layered transition-metal oxides often undergo a transformation during cycling, first becoming layered, then spinel, and, finally, forming a transformation during cycling, first becoming layered, then spinel, and, finally, forming a rock-salt structure [16,17] (Figure 1). For the layered structure, lithium ions and transition- rock-salt structure [16,17] (Figure 1). For the layered structure, lithium ions and transition- metal (TM) ions are situated in two independent layers. When some Li sites are replaced metal (TM) ions are situated in two independent layers. When some Li sites are replaced with TM, the layered structure can change into a defect spinel/spinel-like structure. After with TM, the layered structure can change into a defect spinel/spinel-like structure. After a long cycling process, all metal ions, including Li and TM ions, take up the metal sites a long cycling process, all metal ions, including Li and TM ions, take up the metal sites randomly, leading to the formation of the rock-salt structure. randomly, leading to the formation of the rock-salt structure. Figure 1. Figure 1. Failu Failur re model of lay e model of layer ee red transition d transition me metal oxides (t tal oxides (takin aking Li g Li ric−h rich materials as an materials as an exam example) ple) [16]. [16]. 3.1. Layered Mn-Based Oxides LiMnO is the primary Mn-based material, from which many other cathode materials are 3.1. Layered Mn-Based Oxides derived, including LiNi Mn O , LiMn O (Li MnO ), and Li MnO (Li(Li Mn )O ). 0.5 0.5 2 2 4 0.5 2 2 3 1/3 2/3 2 LiMnO2 is the primary Mn-based material, from which many other cathode materials LiMnO has three different types of structures, with space group Pmnm, C2/m and R3m, are derived, including LiNi0.5Mn0.5O2, LiMn2O4 (Li0.5MnO2), and Li2MnO3 (Li(Li1/3Mn2/3)O2). respectively: orthorhombic, monoclinic and rhombohedral structures [18]. In these, the LiMnO2 has three different types of structures, with space group Pmnm, C2/m and R3m, monoclinic LiMnO with a layered structure can be utilized as a cathode material for LIBs respectively: orthorhombic, monoclinic and rhombohedral structures [18]. In these, the –1 owing to its high theoretical capacity of 285 mAh g [19]. However, orthorhombic LiMnO monoclinic LiMnO2 with a layered structure can be utilized as a cathode material for LIBs is usually produced during the high-temperature solid-state methods [20], indicating that –1 owing to its high theoretical capacity of 285 mAh g [19]. However, orthorhombic LiMnO2 monoclinic LiMnO (Figure 2a) has a metastable phase. The monoclinic LiMnO is usually 2 2 is usually produced during the high-temperature solid-state methods [20], indicating that synthesized by low-temperature ion-exchange methods [21], which sharply raises the monoclinic LiMnO2 (Figure 2a) has a metastable phase. The monoclinic LiMnO2 is usually production cost of LiMnO . At present, LiMnO has not been commercialized not only 2 2 synthesized by low-temperature ion-exchange methods [21], which sharply raises the pro- due to the high production cost, but also because of some performance disadvantages, duction cost of LiMnO2. At present, LiMnO2 has not been commercialized not only due to such as Jahn–Teller distortions, Mn dissolution, low electronic conductivity, and structure the high production cost, but also because of some performance disadvantages, such as transformation to spinel phase, which results in low stability [22]. In Mn-based materials, Jahn–Teller distortions, Mn dissolution, low electronic conductivity, and structure trans- the Jahn–Teller distortion is a common problem which leads to the phase transformation formation to spinel phase, which results in low stability [22]. In Mn-based materials, the 3+ 2+ and the disproportionation reaction of Mn which generates Mn ions and results in Jahn–Teller distortion is a common problem which leads to the phase transformation and Mn dissolution. According to the Jahn–Teller theory, nonlinear molecules with a spatially 3+ 2+ the disproportionation reaction of Mn which generates Mn ions and results in Mn dis- degenerate electronic ground state tend to experience a geometrical distortion to lower solution. According to the Jahn–Teller theory, nonlinear molecules with a spatially degen- 3+ the energy of total molecule. For Mn in MO octahedrons, four d orbital electrons will erate electronic ground state tend to experience a geometrical distortion to lower the en- 3 1 4 be arranged in two e and three t orbitals, forming high-spin t e or low-spin t , g 2g 2g g 2g 3+ ergy of total molecule. For Mn in MO6 octahedrons, four d orbital electrons will be ar- which results in an odd number of electrons in d orbitals and Jahn–Teller distortion [18]. To 3 1 4 ranged in two eg and three t2g orbitals, forming high-spin t2g eg or low-spin t2g , which alleviate the Jahn–Teller distortion, various strategies, such as doping, surface modification, results in an odd number of electrons in d orbitals and Jahn–Teller distortion [18]. To alle- and constructing nanostructures with distinctive morphologies have already been used. viate the Jahn–Teller distortion, various strategies, such as doping, surface modification, and constructing nanostructures with distinctive morphologies have already been used. 3.1.1. LiNi Mn O 0.5 0.5 2 When the Mn cations in LiMnO are partially substituted by Ni cations, a LiNi Mn O 2 x 1-x 2 3.1.1. LiNi0.5Mn0.5O2 solid solution is formed. The structure of LiNi Mn O is determined by the ratio of x 1-x 2 Ni/Mn. When the Mn When the cation ratios in of L Ni/Mn iMnO2< are p 1, the artispinel ally sub str sti uctur tuted by Ni e is formed; cations, otherwise, a LiNixMn the 1- x O -NaFeO 2 solid sostr lut uctur ion is e forme with R d3. The m symmetry structure o (Figur f LiNi e 2 xb) Mn will 1-xObe 2 is generated determined by [23]. The the r hexago- atio of Ni nal/Mn. When the ra -NaFeO -structur tio of Ni ed LiNi/Mn Mn <1, t O he spine delivers l str a u dischar cture is gefor specific med; otherwise, the capacity of about α- 2 0.5 0.5 2 –1 NaF 200 mAh eO2 struc g ,ture with R as well as a3m symmet plateau potential ry (Figur ofe 2b 3.8 ) wi V [24 ll be generat ]. As the valence ed [23]. The hex state of Mn agon and al Batteries 2023, 9, 246 6 of 23 Ni ions in LiNi Mn O is +4 and +2, respectively, only the Ni ions undergo the redox 0.5 0.5 2 reaction from +2 to +4 during the charging and discharging process. The Mn ion keeps 3+ the valance state of +4 without the appearance of Mn , which can effectively avoid Mn 2+ dissolution and the Jahn–Teller distortion [25]. Nevertheless, with the utilization of Ni , the Li/Ni cation mixing problem will appear in LiNi Mn O , which reduces the Li 0.5 0.5 2 diffusion rate and the reversible capacity [26]. In order to enhance the electrochemical performance of LiNi Mn O , various strate- 0.5 0.5 2 gies were employed, including nanostructure construction, element doping and surface coating. The nano-structured LiNi Mn O is mainly synthesized using template meth- 0.5 0.5 2 ods. Yuan et al. [27] carried out an in-situ conversion from
-MnO hollow nanospheres to LiNi Mn O nanoarchitecture spheres (Figure 2c). In comparison to LiNi Mn O 0.5 0.5 2 0.5 0.5 2 particles synthesized via a conventional solid-state reaction process, the nanostructured –1 LiNi Mn O displays a significant improvement in rate performance (121.9 mAh g 0.5 0.5 2 at 3.2 C, 70% of the capacity at 0.2 C), because much faster interfacial kinetics and higher Li insertion/removal rates are realized by reducing the size of LiNi Mn O particles. 0.5 0.5 2 Later, Yuan et al. [28] utilized cryptomelane-type octahedral molecular sieve manganese dioxide (OMS-2) in the form of dendritic nanostructures as templates to prepare the three- dimensional LiNi Mn O nanostructures, which also showed improved rate performance. 0.5 0.5 2 Element doping is an efficacious strategy to enhance the electrochemical properties of cathode materials. Numerous elements have been attempted to dope into LiNi Mn O , 0.5 0.5 2 such as Al [29,30], Mg [31], Ba [32], Cu [33], Sb [34], Ti [35], Zr [36], Si [37], Mo [38] and F [39]. All these doping ions can suppress the Li/Ni cation mixing in LiNi Mn O , and further 0.5 0.5 2 improve the rate capability, specific capacity, and cycling performance. In addition, the special element has particular effects. For instance, LiNi Mn O doped with Al element 0.5 0.5 2 –1 –1 displayed the highest initial specific capacity of 206 mAh g [30] and 215 mAh g [29]. This is because Al doping could narrow the size distribution and decrease Li migration resistance with extended lattice parameter of c axle. Ba doping can effectively promote the cycling performance of LiNi Mn O because the Ba-O bond has a greater dissociation 0.5 0.5 2 –1 –1 energy than the Ni-O bond (563 kJ mol vs. 391.6 kJ mol ), which helps to stabilize the crystal structure [32]. As a result, after 100 cycles, Ba-doped LiNi Mn O can still 0.5 0.5 2 operate at 97% of its initial specific capacity. Cu doping could extend the migration channels of lithium ions in LiNi Mn O , which, in turn, improves the rate performance 0.5 0.5 2 effectively [32]. In addition to single-element doping, the double-element co-doping is applied in LiNi Mn O , which has shown better improvement than single-element 0.5 0.5 2 doping. Jia et al. [40] realized a Na-Al dual-doped LiNi Mn O material, improving 0.5 0.5 2 the reversible capacity, cycling stability, and the stability of discharge midpoint potential (Figure 2d). Na successfully entered into the lithium layer and played a “pillar” role to promote the structural stability, as opposed to hindering the diffusion of Li . The Al-O bond with a higher bond dissociation energy further reinforced the crystal structure. As a result, the Na-Al dual-doped LiNi Mn O displayed a high initial specific capacity of 0.5 0.5 2 –1 216 mAh g and a capacity retention of 90.56% after 180 cycles (Figure 2e). Surface coating can function as a barrier between cathode materials and liquid elec- trolytes, preventing direct contact and further enhancing the interface stability. Meanwhile, the conductivity of cathode materials can also be promoted by coating a conductive layer. Hashem et al. [41] prepared carbon-coated LiNi Mn O using an oxalate coprecipitation 0.5 0.5 2 method, with table sugar as a carbon source. In comparison to pristine LiNi Mn O , the 0.5 0.5 2 as-prepared sample exhibited superior capacity retention (92% after 50 cycles vs. 75% after 30 cycles). Jia et al. [42] coated a stable, 10 nm thick, and lithium super conductive Li TiO 2 3 layer on the surface of LiNi Mn O , which effectively enhanced the Li diffusion rate, 0.5 0.5 2 protected the particle morphologies, and helped to maintain a better structure stability, by reducing side reactions between the cathode materials and liquid electrolytes. In order to further improve the electrochemical properties of LiNi Mn O , element doping and 0.5 0.5 2 surface coating were combined, such as Sb-doped and Sb O -coated LiNi Mn O [34], 2 3 0.5 0.5 2 and Zr-doped and Li ZrO -coated LiNi Mn O [36]. In these, element doping could 2 3 0.5 0.5 2 Batteries 2023, 9, 246 7 of 23 enhance structural stability by reducing the degree of Li/Ni cation mixing, while surface coating could efficaciously shield the active material from direct contact with liquid elec- trolytes, and increase the Li diffusion rate at the interface between electrode and electrolyte. However, present methods for these materials usually follow two steps: preparation of Batteries 2023, 9, x FOR PEER REVIEW 7 of 23 precursors (without doping agents) and preparation of final materials. Since the doping ion is diffused from surface to bulk, resulting in a low doping amount [36], it is important to investigate how to add a doping agent to the precursor. Figure 2. (a) Crystal models of monoclinic LiMnO2 (blue: MnO6 octahedron, green: Li) [23]. (b) Crys- Figure 2. (a) Crystal models of monoclinic LiMnO (blue: MnO octahedron, green: Li) [23]. 2 6 tal models of LiNi0.5Mn0.5O2 [35]. (c) Schematic diagram of the synthesizing nanostructured (b) Crystal models of LiNi Mn O [35]. (c) Schematic diagram of the synthesizing nanostructured 0.5 0.5 2 LiNi0.5Mn0.5O2 derived from the in-situ conversion of γ−MnO2 hollow nanospheres [27]. Schematic LiNi Mn O derived from the in-situ conversion of
MnO hollow nanospheres [27]. Schematic 0.5 0.5 2 2 diagrams (d) and cycling performance (e) of crystal structure of Li1−xNaxNi0.5-yAlyMn0.5O2 [40]. diagrams (d) and cycling performance (e) of crystal structure of Li Na Ni Al Mn O [40]. 1 x x 0.5-y y 0.5 2 Surface coating can function as a barrier between cathode materials and liquid elec- 3.1.2. Li-Rich Mn-Based Oxides trolytes, preventing direct contact and further enhancing the interface stability. Mean- –1 Owing to their high specific capacity (>250 mAh g @0.1C), low cost, and outstanding while, the conductivity of cathode materials can also be promoted by coating a conductive safety, Li-rich Mn-based (LRM) layered oxides are regarded as the most promising can- layer. Hashem et al. [41] prepared carbon-coated LiNi0.5Mn0.5O2 using an oxalate coprecip- didate of cathode materials for the next-generation high-specific-energy LIBs. The main itation method, with table sugar as a carbon source. In comparison to pristine component of LRM cathode materials is the inexpensive Mn element; low Co content LiNi0.5Mn0.5O2, the as-prepared sample exhibited superior capacity retention (92% after 50 (< 10 mol%) or no Co content is the main way to reduce the material cost of LRM. The cycles vs. 75% after 30 cycles). Jia et al. [42] coated a stable, 10 nm thick, and lithium super low cost of LRMs combined with the high energy density can significantly reduce the cost conductive Li2TiO3 layer on the surface of LiNi0.5Mn0.5O2, which effectively enhanced the per Wh. Since Numata et al. [43] reported LiCoO -Li MnO for the first time in 1997 and 2 2 3 Li diffusion rate, protected the particle morphologies, and helped to maintain a bett er Gopukumar et al. [44] originally used Li MnO as a cathode material for LIBs in 1999, 2 3 structure stability, by reducing side reactions between the cathode materials and liquid numerous studies have been conducted on LRM cathode materials. However, the structure electrolytes. In order to further improve the electrochemical properties of LiNi0.5Mn0.5O2, of LRMs has not been figured out yet; there are still two arguments about their structure, element doping and surface coating were combined, such as Sb-doped and Sb2O3-coated one is a two-phase nanodomain (xLi MnO (1-x) LiNi Co Mn O ), and the other one is a x y z 2 3 2 LiNi0.5Mn0.5O2 [34], and Zr-doped and Li2ZrO3-coated LiNi0.5Mn0.5O2 [36]. In these, element single-phase solid solution (Li Ni Co Mn O ) [10]. Nevertheless, both nano-domain 1+x y z 1-x-y-z 2 doping could enhance structural stability by reducing the degree of Li/Ni cation mixing, and solid-solution structures are composed of monoclinic Li MnO (C2/m) and trigonal 2 3 while surface coating could efficaciously shield the active material from direct contact LiMO (R3m) (Figure 3a). with liquid electrolytes, and increase the Li diffusion rate at the interface between elec- Although LRM show great potential as cathode materials for LIBs due to their high trode and electrolyte. However, present methods for these materials usually follow two specific capacity, their low initial coulombic efficiency, serious voltage fading, bad cycling steps: preparation of precursors (without doping agents) and preparation of final materi- and rate properties have limited their practical utilization. Several studies have focused als. Since the doping ion is diffused from surface to bulk, resulting in a low doping amount on the causes of the LRM’s poor initial coulombic efficiency. Oxygen activation brought [36], it is important to investigate how to add a doping agent to the precursor. about by the Li-O-Li configuration (Figure 3b) is not only the cause of LRMs’ high specific capacity [45], but also the main origin of the problems of oxygen loss and transition 3.1.2. Li-rich Mn-Based Oxides –1 Owing to their high specific capacity (>250 mAh g @0.1C), low cost, and outstanding safety, Li-rich Mn-based (LRM) layered oxides are regarded as the most promising candi- date of cathode materials for the next-generation high-specific-energy LIBs. The main component of LRM cathode materials is the inexpensive Mn element; low Co content (< 10 mol%) or no Co content is the main way to reduce the material cost of LRM. The low cost of LRMs combined with the high energy density can significantly reduce the cost per Wh. Since Numata et al. [43] reported LiCoO2-Li2MnO3 for the first time in 1997 and Batteries 2023, 9, 246 8 of 23 metal migration. During the initial charging/discharging process, O gas is generated and released due to the oxidation of some active O, causing an amount of irreversible capacity. Meanwhile, some TM ions will transfer from TM sites to Li sites and form an irreversible spinel phase in the first cycle because the irreversible oxygen release reduces the binding energy of TM ions and oxygen, which further increases irreversible capacity loss. The continuous structural transition from layered to spinel phase has been considered as the cause of the voltage decay and poor cycling performance during cycling [46,47]. In addition, some researchers found that the capacity and voltage decay are caused by the decrease in the redox couple of TM ions [48]. The diffusion kinetics of Li in the crystal structure often sets a limit on the rate performance of LRMs. The pristine structure of cathode materials is destroyed during the phase transition from the layered to spinel structure, leading to a poor rate performance [49,50]. Interestingly, some scholars have explored how the Li content affects the structure of LRM cathode materials. LRMs with lower lithium content had more chemical ordering defects, while the spinel-structured surface had no obvious structural change with the change in lithium content [51]. Some advanced characterization methods are used to reveal the failure mechanism of LRM positive electrode during the cycle [52,53]. Li/Ni and Li/Mn anti-site defects were also discussed in the case of LiNi Co Mn O and Li MnO cathode materials [54]. In order x y (1-x-y) 2 2 3 to avoid the anti-site defects, LiNi Co Mn O and Li MnO were prepared under x y (1-x-y) 2 2 3 Ni-rich condition and under O-rich and Mn-poor conditions, respectively. In order to solve the above problems, the main strategies are doping, coating and surface modification. The elements that are usually utilized to dope cations at TM sites have a stronger bond with oxygen than Ni, Co, and Mn, which can efficaciously inhibit the structural shift from layered to spinel [55]. Surface coating is utilized in LRMs to keep the structural stability of the electrode/electrolyte interface and reduce voltage fading [55]. However, doping and coating are unable to address the issue of oxygen loss, which is the essential problem of LRMs. One present effective strategy to reduce the oxygen loss is creating defects/vacancies which is mostly realized by surface modification. For exam- ple, Qiu et al. [56] modified the Li[Li Ni Co Mn ]O with CO by gas–solid 0.144 0.136 0.136 0.544 2 2 interface reaction (GSIR) to form a layer of oxygen vacancy with a thickness of 10–20 nm on the surface of material (Figure 3c). Oxygen vacancy can effectively inhibit the release of lattice oxygen; hence, the obtained sample has shown an improved initial coulomb efficiency of 93.2% and cycling performance; Peng et al. [57] treated the LRM with oleic acid to manufacture cation and anion double defects and an in-situ surface reconstruction layer to reduce oxygen release and improve structural stability, which can enable precise control of the ICE from 84.1% to 100.7%. At the same time, this sample presents a high 1 1 specific capacities of 330 mAh g at 0.1 C with a large energy density of 1143 Wh kg , 1 + and 276 and 250 mAh g at 1 and 5 C due to the fast kinetics of Li and its electron (Figure 3d). In contrast to the surface modification, Zhu et al. [58] developed the method of molten molybdate-assisted LiO extraction to create gradient Li-rich single crystals, which can inhibit oxygen loss. Lithium is rich in the core of the particle, poor on the surface, and continuous in between (Figure 3e). The prepared material shows a good voltage and cycling stability with high discharge specific capacity of 250.4 mAh g and high average voltage of 3.368 V after 200 cycles at 0.2 C. Since the price of Cobalt salts has increased sharply in recent years, LRMs with low Co content or without Co are more attractive in industrial application due to their rela- tively low cost. However, LRMs with less Co (<10%) show poor structure stability and ionic/electron conductivity, such as Li Ni Mn O . To enhance the electrochemical prop- 1.2 0.2 0.6 2 erties, Chen et al. [59] prepared a hierarchical Li Ni Mn O quasi-sphere with a plane- 1.2 0.2 0.6 2 based surface, in which the electrochemical active planes allow for fast Li transport kinetics due to efficient ion and electron transport (Figure 3f). The as-prepared Li Ni Mn O 1.2 0.2 0.6 2 displayed improved cycling and rate performance. On the surface of the Li Ni Mn O 1.2 0.2 0.6 2 microsphere, Ding et al. [14] pyrolyzed urea to form a multifunctional surface modification, which could simultaneously construct oxygen vacancy, integrated spinel phase and N- Batteries 2023, 9, 246 9 of 23 doped carbon nanolayer (Figure 3g). The integration of oxygen vacancy and spinel phase Batteries 2023, 9, x FOR PEER REVIEW 9 of 23 could not only inhibit the irreversible release of O but also promote the diffusion of Li . Meanwhile, the N-doped carbon nanolayer with high electrical conductivity could promote electron transport and reduce electrolyte corrosion. The electrochemical results show that in between (Figure 3e). The prepared material shows a good voltage and cycling stability the surface-modified Li Ni Mn O can retain its initial specific capacity up to 89.9% 1.2 0.2 0.6 2 −1 with high discharge specific capacity of 250.4 mAh g and high average voltage of 3.368 after 500 cycles at 1 C, and its voltage decay rate per cycle is merely 1.09 mV (Figure 3h), V after 200 cycles at 0.2 C. which significantly inhibits the capacity and voltage decay. Figure 3. (a) Schematic diagram of Li−rich materials [10]. (b) Structural origin of the preferred oxy- Figure 3. (a) Schematic diagram of Li rich materials [10]. (b) Structural origin of the preferred gen oxidation along the Li−O−Li configuration [45]. (c) Schematic diagram of gas−solid interface oxygen oxidation along the Li O Li configuration [45]. (c) Schematic diagram of gas solid interface reaction (GSIR) between Li−rich layered oxides and carbon dioxide [56]. (d) Schematic diagram of reaction (GSIR) between Li rich layered oxides and carbon dioxide [56]. (d) Schematic diagram of oleic acid−assisted interface engineering [57]. (e) Schematic of the Li gradient Li−rich material [58]. oleic acid assisted interface engineering [57]. (e) Schematic of the Li gradient Li rich material [58]. (f) Schematic illustration of hierarchical structured Li1.2Mn0.6Ni0.2O2 material [59]. (g) Schematic il- (f) Schematic illustration of hierarchical structured Li Mn Ni O material [59]. (g) Schematic 1.2 0.6 0.2 2 lustration and cycling performance (h) (blue lines) of Li1.2Mn0.6Ni0.2O2 by three−in−one surface mod- illustration and cycling performance (h) (blue lines) of Li Mn Ni O by three in one surface 1.2 0.6 0.2 2 ification [14]. modification [14]. Since the price of Cobalt salts has increased sharply in recent years, LRMs with low 3.2. Spinel Mn-Based Oxides Co content or without Co are more att ractive in industrial application due to their rela- Spinel Mn-based oxides show great structural stability and have been widely applied tively low cost. However, LRMs with less Co (<10%) show poor structure stability and in LIBs, MIBs [60,61], and so on. LiMn O and LiNi Mn O are, at present, the main 2 4 0.5 1.5 4 ionic/electron conductivity, such as Li1.2Ni0.2Mn0.6O2. To enhance the electrochemical prop- spinel Mn-based oxides for lithium batteries. Although their theoretical specific capacity is erties, Chen et al. [59] prepared a hierarchical Li1.2Ni0.2Mn0.6O2 quasi-sphere with a plane- only about half that of layered Mn-based oxides, their more stable structures and quicker based surface, in which the electrochemical active planes allow for fast Li transport ki- Li diffusion compared with layered forms have attracted much attention in scientific netics due to efficient ion and electron transport (Figure 3f). The as-prepared research and industrial application. LiMn O has become the most mature commercial 2 4 Li1.2Ni0.2Mn0.6O2 displayed improved cycling and rate performance. On the surface of the material in Mn-based cathode materials, while LiNi Mn O was originally prepared and 0.5 1.5 4 Li1.2Ni0.2Mn0.6O2 microsphere, Ding et al. [14] pyrolyzed urea to form a multifunctional investigated as one of the metal-substituted LiMn O derivatives. Since the low valance 2 4 surface modification, which could simultaneously construct oxygen vacancy, integrated 2+ state of Ni can increase the valence state of Mn and further prevent the Jahn–Teller spinel phase and N-doped carbon nanolayer (Figure 3g). The integration of oxygen va- 2+ effect and Mn dissolution, Ni doping can stabilize the crystal structure [62]. Similar to cancy and spinel phase could not only inhibit the irreversible release of O2 but also pro- the aforementioned LiNi Mn O , the valance state of Mn in LiNi Mn O is +4, so 0.5 0.5 2 0.5 1.5 4 mote the diffusion of Li . Meanwhile, the N-doped carbon nanolayer with high electrical only Ni ions undergo oxidation and reduction during cycling. Compared with LiMn O , 2 4 conductivity could promote electron transport and reduce electrolyte corrosion. The elec- –1 LiNi Mn O has a comparable cost per Wh (0.13 CNY Wh ), but a higher potential 0.5 1.5 4 trochemical results show that the surface-modified Li1.2Ni0.2Mn0.6O2 can retain its initial specific capacity up to 89.9% after 500 cycles at 1 C, and its voltage decay rate per cycle is merely 1.09 mV (Figure 3h), which significantly inhibits the capacity and voltage decay. Batteries 2023, 9, 246 10 of 23 –1 plateau of ~4.7 V as well as a higher energy density of 650 Wh kg . When matched with a suitable high-voltage electrolyte, LiNi Mn O has great potential to be widely applied. 0.5 1.5 4 LiMn O is a cubic spinel with an Fd3m space group, and can be synthesized via 2 4 a simple solid-state reaction in an air atmosphere at a high temperature. In the spinel structure of LiMn O , the Li , Mn ions and O ions are situated at the 8a tetrahedral sites, 2 4 16d octahedral sites and 32e positions, respectively. O ions are cubic-close-packed [63] (Figure 4a). The edge-shared MnO octahedrons can construct a continuous three-dimensional cubic array, and, further, cause the formation of the robust Mn O spinel framework, in 2 4 which fast diffusion of Li is realized. Since the theoretical specific capacity and average –1 discharge voltage of LiMn O is relatively low (146 mAh g , ~3.8 V), it exhibits a rela- 2 4 tively low energy density. When the discharge voltage is lower than 3 V, the lithiation of LiMn O occurs, which leads to irreversible phase transformation from spinel to rock-salt 2 4 (Li Mn O ) [63]. Besides the low energy density, another problem of LiMn O is its poor 2 2 4 2 4 cycling performance due to its Mn dissolution, especially at high temperatures. In LiMn O , 2 4 3+ 4+ the average valance state of Mn ions is +3.5, mostly composed of Mn and Mn . Mn 2+ dissolution has been mainly ascribed to the dissolution of Mn , which originates from 3+ 3+ 2+ 4+ the disproportionation of Mn (2Mn !Mn +Mn ) [64,65]. Meanwhile, the currently used LiPF -based carbonate electrolyte will produce large amounts of hydrofluoric acid 4+ (HF) with the presence of trace water, which accelerates the dissolution of Mn . Surface coating can also effectively promote the electrochemical performance of LiMn O [66–70]. 2 4 Coating with solid electrolytes can avoid the Mn reactions of LiMn O in the case of liquid 2 4 electrolytes. Li La Zr Nb O (LLZNO) electrolytes were used by Bi et al. [71] to 6.375 3 1.375 0.625 12 coat LiMn O , which not only suppressed Mn reactivity but also enhanced the interface 2 4 between Li La Zr Ta O (LLZTO) and LiMn O . After 100 cycles at 0.2C and 55 C, 6.4 3 1.4 0.6 12 2 4 the solid battery with LMO@LLZNO cathode, LLZTO electrolyte, and Li metal showed a high discharge capacity retention of 81.3%. LiNi Mn O , one of the most significant LiMn O derivatives, was originally syn- 0.5 1.5 4 2 4 thesized in 1996 via substituting part of Mn sites with Ni. LiNi Mn O contains two 0.5 1.5 4 types of spinel crystal structures, one is face-centered cubic (Fd3m) with disordered TM ions, and the other is primitive simple cubic (P4 32) with ordered TM ions [72] (Figure 4b,c). In the Fd3m space group, Li occupies tetrahedral 8a sites, Mn/Ni ions occupies octahedral 2- + 16d sites randomly, and O occupies 32e sites. In P4 32 space group, Li occupies tetra- 4+ 2+ 2- hedral 8a sites, Mn occupies 12d sites, Ni occupies 4b sites, and O occupies 24e and 2+ 4+ 8c sites. It is worth noting that Ni replaces part of Mn sites orderly in the P4 32 space group, different from that in the Fd3m space group. Compared with P4 32, LiNi Mn O 3 0.5 1.5 4 with the Fd3m space group exhibits outstanding electrochemical performance and struc- tural reversibility. LiNi Mn O has been deemed as one of the most promising cath- 0.5 1.5 4 ode materials for LIBs, owing to its high working voltage (4.7 V), high specific capacity –1 (146.7 mAh g ), high energy density, low cost, and good cycle stability [73,74]. However, in addition to some common obstacles for Mn-based cathode such as TM ions’ dissolution 3+ and the Jahn–Teller effect of Mn , LiNi Mn O also has great side reactions with liquid 0.5 1.5 4 electrolytes due to its high voltage plateau [75]. To solve the above problems of LiMn O and LiNi Mn O , various strategies, 2 4 0.5 1.5 4 including element doping, surface coating, an appropriate synthesis method, and elec- trolyte modification, etc., are proposed to improve their electrochemical performance. 3+ Element doping can restrain the Mn disproportionation reaction and Jahn–Teller distor- tion for spinel Mn-based oxides. Present doping ions can be divided into two categories, cations and anions. Present doping cations include Al, Cr, Co, Ga, Pr, Gd, La, Ce, Nd, Sm, Sc, Y, Tb, Er, B, Fe, Mg, Ti, Ru, Si, Ni, Zn, Cu, Sn, Li, and Na ions, etc., and anions include F, Cl, Br S, and PO ions, etc., which have been summarized by Cui et al. [7]. Among the various single cation doping options, Al doping has shown the best improve- ment effect on the cycling performance of LiMn O . Sun et al. [75], Xia et al. [76] and 2 4 Xiao et al. [77] synthesized the Al-doped samples, which show excellent cycling perfor- mances with capacity retentions of 98.5%, 99.34% and 99.3% after 50 cycles at ~50 C, Batteries 2023, 9, 246 11 of 23 respectively. The improvement effects can be summarized as: (1) an Al–O bond with a 1 1 higher bonding energy than Mn–O (512 kJ mol vs. 402 kJ mol ) could enhance the stability of the spinel structure during insertion/de-insertion of lithium; (2) the smaller lattice parameter of Al-doped LiMn O alleviates the dissolution of the active material 2 4 and maintains its structural integrity. Meanwhile, a multi-doped strategy with three ele- ments also shows potential to greatly enhance the electrochemical properties of LiMn O . 2 4 Sun et al. prepared the Li, Co, Gd multi-doped LiMn O , which shows an outstanding ca- 2 4 pacity retention of 98.3% after 100 cycles at 25 C due to its improved structure stability [78]. Manthiram et al. synthesized several Li-M-F (M = Ti, Ni, Cu, Fe, Co, Zn) multi-doped + 2+ LiMn O with excellent electrochemical performance. In these samples, Li and Ni are 2 4 used to prevent Mn dissolution which is caused by a much smaller lattice-parameter differ- ence between the two cubic phases formed during the charge–discharge process. However, the low-valance-state ion doping will increase the valance state of Mn, causing low specific capacity (<100 mAh g ). Then, F doping was applied to decrease the valence state of Mn in the Li-Ni co-doped LiMn O . Among these materials, LiMn Li Zn O F 2 4 1.85 0.075 0.075 3.85 0.15 displayed a high specific capacity of 113 mAh g , good cycling performance with a capac- ity retention of 94.6% after 50 cycles at 60 C, and excellent rate capability with a retention of 96% at 4 C of its initial specific capacity at 0.1 C [79]. Anion doping can also inhibit the structural changes during the charging/discharging process, thereby improving the elec- trochemical properties of LiNi Mn O . Previous research has shown that doping with F 0.5 1.5 4 can suppress the generation of NiO impurity, reducing TMs dissolution and improving the rate performance of LiNi Mn O . It is a pity that doping is unable to prevent undesired 0.5 1.5 4 side reactions between LiNi Mn O and liquid electrolytes [80]. 0.5 1.5 4 Coating can effectively inhibit side reactions by impeding the direct contact between cathodes and liquid electrolytes. Ideal coating materials for spinel Mn-based oxides should have a good match with the spinel lattice and good diffusion ability of Li and elec- trons. Chong et al. [81] synthesized Li PO -coated LiNi Mn O by solid-state reaction. 3 4 0.5 1.5 4 With the Li PO layer (<6 nm), coated LiNi Mn O had a disordered crystal structure, 3 4 0.5 1.5 4 protected cathode-electrolyte interface, and dramatically enhanced cycling performance. Jang et al. [82] synthesized LiFePO -modified spinel LiNi Mn O through a single-step 4 0.5 1.5 4 coating process. LiFePO coating greatly improved the thermal stability and high tempera- ture performance, with negligible discharge-capacity reduction. Fang et al. [83] employed atomic layer deposition (ALD) to coat an ultrathin Al O layer onto LiNi Mn O parti- 2 3 0.5 1.5 4 cles (Figure 4d). Al O coating protected LiNi Mn O from direct exposure to liquid 2 3 0.5 1.5 4 electrolytes, which improved the cycling performance of LiNi Mn O . The Al O -coated 0.5 1.5 4 2 3 LiNi Mn O showed 63% capacity retention after 900 cycles (Figure 4e), whereas the 0.5 1.5 4 bare LiNi Mn O maintained 75% of the original capacity after 200 cycles. Apart from 0.5 1.5 4 inorganic coating, organic materials such as polyimide (PI) [84] and polypyrrole (PPy) [85] can also improve the electrochemical performance of spinel Mn-based oxides. Kim et al. [86] utilized thermal polymerization to produce PI coating from polyamic acid and found that 0.3 wt % PI coated LiNi Mn O delivered excellent cycle ability with capacity retention 0.5 1.5 4 of >90% at 55 C. Choosing an appropriate synthesis method can also improve the electrochemical prop- erties of spinel Mn-based oxides, especially LiNi Mn O . The solid-phase method, sol- 0.5 1.5 4 gel method, and molten-salt method are the three main methods to synthesize LiNi Mn O , 0.5 1.5 4 while each technique has its own advantages and problems. The solid-phase method is a simple, economical and time-saving method [87]. However, the produced LiNi Mn O 0.5 1.5 4 not only has a lot of oxygen defects and impurities such as NiO, but also uneven particle- size distribution and severe agglomeration of particles. Oxygen defects usually lead to 3+ the generation of more Mn , which will have Jahn–Teller distortion, a disproportionation 2+ reaction to producing soluble Mn , and, thereby, impair the electrochemical properties of LiNi Mn O [88]. Rosedhi et al. [89] synthesized LiNi Mn O via ball-milling 0.5 1.5 4 0.5 1.5 4 and following calcination at 750 C. The as-prepared LiNi Mn O displayed an initial 0.5 1.5 4 –1 –1 discharge capacity of 81 mAh g and 86 mAh g after 100 cycles at 1C. The widely used Batteries 2023, 9, 246 12 of 23 sol-gel method can synthesize LiNi Mn O with good crystallization, small particle size, 0.5 1.5 4 homogenous dispersion, and outstanding electrochemical performance [90]. In comparison to other synthetic methods, the sol-gel method is much more complicated, time-consuming, and relatively expensive. Cui et al. [91] synthesized nanosized LiNi Mn O with an 0.5 1.5 4 average particle size of 80–100 nm, via a high-oxidation-state manganese sol-gel method. The produced LiNi Mn O materials not only have an impurity-free cubic spinel struc- 0.5 1.5 4 ture, but also exhibit excellent dispersion. Both the solid-phase method and sol-gel method need a high-temperature annealing process, which will easily cause impurities and oxygen Batteries 2023, 9, x FOR PEER REVIEW 12 of 23 defects. The molten-salt method is a simple method for preparing complex oxides with pure phase in a low-melting point flux [92]. Wen et al. [93] synthesized spherical LiNi Mn O 0.5 1.5 4 materials using the molten-salt method. The as-prepared material displayed a high dis- –1 –1 charge capacity of 129 mAh g in the first cycle and 127 mAh g after 50 cycles. The good polyamic acid and found that 0.3 wt % PI coated LiNi0.5Mn1.5O4 delivered excellent cycle retention of capacity is credited to significantly fewer impurity phases than that prepared ability with capacity retention of >90% at 55 °C. via a solid-state reaction. Figure 4. Figure 4.Cry Crystal stal stru structur cture of e of LiMn LiMn2O O 4 (( aa ) [6 ) [62 2], LiNi ], LiNi0.5Mn Mn 1.5OO 4 with with FdFd3m 3 sp m space ace gr gr ou oup p (b (b ) a ) and nd P4 P4 332 32 2 4 0.5 1.5 4 3 space group (c) [72]. Diagram showing (d) and long−term cycling performance (e) of ALD−based space group (c) [72]. Diagram showing (d) and long term cycling performance (e) of ALD based Al2O3 coating LiNi0.5Mn1.5O4 [83]. Al O coating LiNi Mn O [83]. 2 3 0.5 1.5 4 For LiNi Mn O , the high working voltage at around 4.7 V makes it promising Choosing an 0.5 appropriate 1.5 4 synthesis method can also improve the electrochemical –1 –1 due to its high theoretical energy density (690 Wh kg = 147 mAh g 4.7 V), but it properties of spinel Mn-based oxides, especially LiNi0.5Mn1.5O4. The solid-phase method, greatly limits the compatibility of LiNi Mn O with conventional liquid electrolytes. sol-gel method, and molten-salt method a 0.5 re the three ma 1.5 4 in methods to synthesize The mainstream liquid electrolyte is prepared by mixing organic carbonate esters and LiNi0.5Mn1.5O4, while each technique has its own advantages and problems. The solid- LiPF and will undergo a continuous decomposition above 4.5 V versus Li /Li, resulting phase method is a simple, economical and time-saving method [87]. However, the pro- in increased thickness of solid electrolyte interphase (SEI), aggravated dissolution of Mn duced LiNi0.5Mn1.5O4 not only has a lot of oxygen defects and impurities such as NiO, but ions, and destroyed electrode structure [94]. Hence, electrolyte additives, fluorinated also uneven particle-size distribution and severe agglomeration of particles. Oxygen de- electrolytes and solid electrolytes are utilized to improve the compatibility between high- 3+ fects usually lead to the generation of more Mn , which will have Jahn–Teller distortion, voltage LiNi Mn O and electrolytes. Some additives can be polymerized to form a 2+ 0.5 1.5 4 a disproportionation reaction to producing soluble Mn , and, thereby, impair the electro- protective layer and thereby suppress the side reaction between material and electrolytes, chemical properties of LiNi0.5Mn1.5O4 [88]. Rosedhi et al. [89] synthesized LiNi0.5Mn1.5O4 prior to the decomposition of electrolyte. For example, Li et al. [95] utilized lithium bis(2- via ball-milling and following calcination at 750 °C. The as-prepared LiNi0.5Mn1.5O4 dis- methyl-2-fluoromalonato)borate (LiBMFMB) as an electrolyte additive for LiNi Mn O . –1 –1 0.5 1.5 4 played an initial discharge capacity of 81 mAh g and 86 mAh g after 100 cycles at 1C. With the addition of LiBMFMB, a thinner and stabler SEI was formed to prevent further The widely used sol-gel method can synthesize LiNi0.5Mn1.5O4 with good crystallization, decomposition of the carbonate solvent. Johannes et al. [96] synthesized a novel electrolyte small particle size, homogenous dispersion, and outstanding electrochemical perfor- with
-Butyrolactone (GBL) as solvent, fluoroethylene carbonate (FEC) as solid electrolyte mance [90]. In comparison to other synthetic methods, the sol-gel method is much more interphase (SEI) additive, and lithium tetrafluro borate (LiBF4) as electrolyte salt, and its complicated, time-consuming, and relatively expensive. Cui et al. [91] synthesized na- cutoff potential was increased to 4.6 V [97,98]. Apart from the continuous side reactions nosized LiNi0.5Mn1.5O4 with an average particle size of 80–100 nm, via a high-oxidation- between electrolytes and cathode materials, the reactions between LiPF and trace water state manganese sol-gel method. The produced LiNi0.5Mn1.5O4 materials not only have an will produce HF, which will corrode the cathodes and then cause the TMs dissolution [99]. impurity-free cubic spinel structure, but also exhibit excellent dispersion. Both the solid- phase method and sol-gel method need a high-temperature annealing process, which will easily cause impurities and oxygen defects. The molten-salt method is a simple method for preparing complex oxides with pure phase in a low-melting point flux [92]. Wen et al. [93] synthesized spherical LiNi0.5Mn1.5O4 materials using the molten-salt method. The as- –1 prepared material displayed a high discharge capacity of 129 mAh g in the first cycle –1 and 127 mAh g after 50 cycles. The good retention of capacity is credited to significantly fewer impurity phases than that prepared via a solid-state reaction. For LiNi0.5Mn1.5O4, the high working voltage at around 4.7 V makes it promising due –1 –1 to its high theoretical energy density (≈690 Wh kg = 147 mAh g × 4.7 V), but it greatly limits the compatibility of LiNi0.5Mn1.5O4 with conventional liquid electrolytes. The Batteries 2023, 9, 246 13 of 23 To solve the above-mentioned problems, additives that have strong bonding with HF, - + F , and H have been utilized to scavenge the detrimental HF. Li et al. [100] utilized pentafluorophenyltriethoxysilane (TPS) as an additive to improve the cycling stability of LiNi Mn O . TPS not only constructed an ionically conductive cathode electrolyte 0.5 1.5 4 interphase (CEI) film, but also captured detrimental species in electrolytes. As a result, after 400 cycles at 1 C, LiNi Mn O presented an improved discharge capacity retention 0.5 1.5 4 from 28% to 85%, with the addition of 1 wt % TPS. In addition, fluorinated electrolytes can promote the capacity of LiNi Mn O because of their high oxidation potential. Zhang 0.5 1.5 4 et al. [101] substituted ethyl methyl carbonate (EMC) and (ethylene carbonate) EC with fluorinated EMC and fluorinated cyclic carbonate (F-AEC) and found that fluorinated electrolytes greatly improve the voltage limits of the electrolyte, and thereby enhance the electrochemical properties of full batteries at elevated temperatures. Developing solid electrolytes can completely solve the decomposition and side reactions of liquid electrolytes, because some solid electrolytes show a wide voltage window beyond 5 V [102]. Li et al. [103] realized high-voltage cycling in solid-state systems using LiNi Mn O cathode, Lipon 0.5 1.5 4 solid electrolyte, and Li anode. The solid-state high-voltage battery delivered a remarkable capacity retention of 90% over 1000 cycles, a high coulombic efficiency of 97% and a round-trip energy efficiency of 97%. 3.3. Cation-Disordered Rock-Salt Mn-Based Li-Rich Materials The cation-disordered rock-salt (DRX) structure is mostly considered to be the product of material that has lost its electrochemical activity after cycling. As opposed to the layered structure’s ordered arrangement of metal ions, Li and TM ions of DRXs are randomly mixed in each other ’s positions [104]. Therefore, the DRX structure was thought to be harmful to Li transport and incapable of providing reversible capacity, until Ceder et al. discovered the electrochemical activity of Li Mo Cr O in 2014 [105]. In the DRX 1.211 0.467 0.3 2 structure, the Li diffusion between two octahedral (o) sites must undergo an intermediate tetrahedral (T ) site, known as o-t-o diffusion. The intermediate tetrahedral (Td) site has four face-sharing octahedrons that can be filled with Li or M, which forms a “tetrahedral cluster”, generating five possible situations: Li (0-TM), Li M (1-TM), Li M (2-TM), LiM 4 3 2 2 3 (3-TM) and M (4-TM) [104] (Figure 5a). The o-t-o Li diffusion needs a minimum of two + + Li , making a Li , Li M, Li M environment possible through Li diffusion pathways. 4 3 2 2 However, the Li mobility in structures that only contain 1-TM channel or 2-TM channels would be negligible in DRX compounds. Then, in order to guarantee the Li diffusion in DRX compounds, sufficient 0-TM channels are required to form the 0-TM percolating network [105], which can be constructed by excess Li in a DRX structure (Figure 5b,c). Therefore, the investigations of DRX compounds mainly focus on the Li-rich materials. Since Ni, Co and Mn will not migrate to Li sites when a significant amount of Li is removed, causing further structural change, they are the main elements in the majority of layered cathode materials [106]. In contrast, the DRX Li-rich materials show an important advantage of using a large range of TM species, such as V, Mn, Nb, Mo, Ti, Cr and so on. –1 Another advantage of DRX Li-rich materials is the higher specific capacity (>250 mAh g ), compared with present NCM materials [104]. Meanwhile, the intrinsic cation disorder can cause a small volume change in the DRX structure which, in principle, is advanta- geous for the cycling of cathode materials. Even the nearly zero-strain Li-ion cathodes of Li V Nb O and Li V Nb O F have been designed and synthesized by 1.3 0.4 0.3 2 1.25 0.55 0.2 1.9 0.1 Ceder et al. in order to enhance the cycling of cathode materials [107] (Figure 5d). However, the rate capability and cycling performance of DRX Li-rich materials are extremely poor and far from that of today’s commercial cathode materials. What is worse, only few of the improvement strategies are effective at promoting the electrochemical properties of DRX materials, such as fluorination. Among various DRX Li-rich cathodes, the Mn-based materials have shown many advantages, including low cost, resource friendliness, and high energy density (about 1000 Wh kg ) [108,109], demonstrating great potential as low-cost high-energy-density Batteries 2023, 9, x FOR PEER REVIEW 14 of 23 Since Ni, Co and Mn will not migrate to Li sites when a significant amount of Li is removed, causing further structural change, they are the main elements in the majority of layered cathode materials [106]. In contrast, the DRX Li-rich materials show an important advantage of using a large range of TM species, such as V, Mn, Nb, Mo, Ti, Cr and so on. Another advantage of DRX Li-rich materials is the higher specific capacity (>250 mAh g ), compared with present NCM materials [104]. Meanwhile, the intrinsic cation disorder can cause a small volume change in the DRX structure which, in principle, is advanta- geous for the cycling of cathode materials. Even the nearly zero-strain Li-ion cathodes of Li1.3V0.4Nb0.3O2 and Li1.25V0.55Nb0.2O1.9F0.1 have been designed and synthesized by Ceder et al. in order to enhance the cycling of cathode materials [107] (Figure 5d). However, the rate capability and cycling performance of DRX Li-rich materials are extremely poor and far from that of today’s commercial cathode materials. What is worse, only few of the improvement strategies are effective at promoting the electrochemical properties of DRX materials, such as fluorination. Among various DRX Li-rich cathodes, the Mn-based materials have shown many ad- Batteries 2023, 9, 246 14 of 23 vantages, including low cost, resource friendliness, and high energy density (about 1000 −1 Wh kg ) [108,109], demonstrating great potential as low-cost high-energy-density mate- rials. The Mn-based Li-rich Li2Mn2/3Nb1/3O2F and Li2Mn1/2Ti1/2O2F materials were synthe- materials. The Mn-based Li-rich Li Mn Nb O F and Li Mn Ti O F materials were 2 2/3 1/3 2 2 1/2 1/2 2 sized via high-energy ball-milled methods and exhibited a high specific capacity of >300 synthesized via high-energy ball-milled methods and exhibited a high specific capacity of −1 –1 mAh g and high energy density of around 1000 Wh kg without the use of O redox [108]. 1 –1 >300 mAh g and high energy density of around 1000 Wh kg without the use of O re- Another series of Mn-based LixMn2-xO2-yFy compounds with a high capacity of 350 mAh dox [108]. Another series of Mn-based Li Mn O F compounds with a high capacity of –1 x 2-x 2-y y g was also synthesized using ball-milling methods [109] (Figure 5e). Although the high- –1 350 mAh g was also synthesized using ball-milling methods [109] (Figure 5e). Although energy ball-milling process is a good method to produce DRX compounds even for mate- the high-energy ball-milling process is a good method to produce DRX compounds even rials whose disorder cannot be accessed thermally, its high energy consumption and low for materials whose disorder cannot be accessed thermally, its high energy consumption productivity make it unsuitable for industrial application. To realize the industrial prod- and low productivity make it unsuitable for industrial application. To realize the industrial uct of Mn-based DRX compounds, the traditional solid-state method has been employed. product of Mn-based DRX compounds, the traditional solid-state method has been em- However, these Mn-based DRX compounds must contain d0 element to stabilize the dis- ployed. However, these Mn-based DRX compounds must contain d0 element to stabilize 5+ ordered structure, and Nb is often used. For instance, Ceder et al. [110] and Tong et al. 5+ the disordered structure, and Nb is often used. For instance, Ceder et al. [110] and [15] prepared Li-Mn-Nb-O-F compounds by sintering at 1000 °C under an argon atmos- Tong et al. [15] prepared Li-Mn-Nb-O-F compounds by sintering at 1000 C under an argon phere for 7h using Li2CO3, MnO2, Nb2O5 and LiF as raw materials. The as-prepared atmosphere for 7h using Li CO , MnO , Nb O and LiF as raw materials. The as-prepared 2 3 2 2 5 –1 Li1.2Mn0.625Nb0.175O1.95F0.05 showed a high specific capacity of 330 mAh g and high average –1 Li Mn Nb O F showed a high specific capacity of 330 mAh g and high 1.2 0.625 0.175 1.95 0.05 discharge voltage of 3.2 V. However, to improve its electrochemical performance, fluori- average discharge voltage of 3.2 V. However, to improve its electrochemical performance, nated DRX needs a larger F content since the F solubility in DRX using LiF as the F source fluorinated DRX needs a larger F content since the F solubility in DRX using LiF as the F is limited at 7.5 at%. Chen et al. [111] synthesized the Li-Mn-Nb-O-F compounds using a source is limited at 7.5 at%. Chen et al. [111] synthesized the Li-Mn-Nb-O-F compounds us- solid-state calcination method using poly(tetrafluoroethylene) (PTFE) as an F source; the ing a solid-state calcination method using poly(tetrafluoroethylene) (PTFE) as an F source; incorporation of F content was up to 12.5 at%. In this, the Li-Mn-Nb-O-F compound with the incorporation of F content was up to 12.5 at%. In this, the Li-Mn-Nb-O-F compound −1 10 at% of F substitution displays a reversible discharge capacity of ≈255 mAh g and good with 10 at% of F substitution displays a reversible discharge capacity of 255 mAh g cycling performance (123% capacity retention after 30 cycles) (Figure 5f). and good cycling performance (123% capacity retention after 30 cycles) (Figure 5f). Figure 5. (a) Cation disorder causes the formation of all varieties of tetrahedral clus- ters (0 TM, 1 TM, 2 TM, 3 TM and 4 TM channels) [104]. (b) Computed probability of discovering a percolating network of 0-TM channels versus Li content (x in Li TM O ) 2-x 2 and cation mixing (TM /TM 100%) [105]. (c) Accessible Li content by a per- Li layers TM layers colating 0 TM network versus Li content and cation mixing [105]. (d) Structure change of Li V Nb O F during charging and discharging [107]. (e) Electrochemical perfor- 1.25 0.55 0.2 1.9 0.1 mance of Li Mn(III) Mn(IV) O F (HLF33) [109]. (f) Capacity retention of 1.3333 0.3333 0.3333 1.6667 0.3333 Li Mn Nb O (F0), Li Mn Nb O F (F2.5), Li Mn Nb O F (F5), and 1.2 0.6 0.2 2 1.2 0.625 0.175 1.95 0.05 1.2 0.65 0.15 1.9 0.1 Li Mn Nb O F (F10) cathodes [111]. 1.2 0.7 0.1 1.8 0.2 3.4. Phospho-Olivine Mn-Based Compounds Since the groundbreaking investigation of Goodenough et al. in 1996 [112], the Phospho-olivine LiMPO compound (where M = Fe, Mn, Co or Ni) has been recognized as a viable cathode material for LIBs. Among the olivine phosphate family, LiMnPO and LiFe Mn Fe PO are excellent candidates for stable and high-energy-density cathode x 1-x x 4 materials. Phospho-olivine LiMnPO is made up of a hexagonal close-packed (hcp) frame- work of oxygen with Pnma space group, where Mn and Li occupy octahedral Figure 4a,c Batteries 2023, 9, 246 15 of 23 octahedral sites, and P atom being in Figure 4c tetrahedral site, respectively (Figure 6a). LiFe Mn Fe PO has a similar crystal structure, with Li and Mn/Fe atoms located in the x x 1-x 4 octahedral Figure 4a,c sites, respectively, while the tetrahedral Figure 4c position harbors the P atoms [113]. LiMnPO is attractive due to its high theoretical discharge capacity –1 + (170 mAh g ), high operating voltage (4.1 V vs. Li/Li ), high structural stability during –1 the charging/discharging process, superior theoretical energy density (701 Wh kg = –1 171 mAh g 4.1 V), and high safety due to its strong P–O covalent bond. Additionally, LiMnPO shows low toxicity, environmental friendliness and low cost. However, the low –10 –1 –16 2 –1 electronic conductivity (<10 S cm ) and low lithium–ion diffusion (<10 cm S ) of LiMnPO significantly affects its rate capability and cycling performance at high rates and limits its industrial application [114]. LiFe Mn PO is a solid-solution material between 1-x 4 2+ LiMnPO and LiFePO with uniform distribution of Mn and Fe elements [115–117]. Fe 4 4 doping can reduce electrochemical polarization, leading to an enhancement in aspects such as the reversibility of electrodes, the ability of de-lithiation/lithiation and the diffusion of + 3+ Li . Moreover, Jahn–Teller lattice distortion of Mn ions will hinder the lithium extrac- tion/insertion process in LiMnPO , which will further cause a large volume change and interface strain between the LiMnPO and MnPO phases [118]. In order to overcome these 4 4 problems, several methods have been applied to improve the electrochemical properties of LiMnPO and LiFe Mn PO , including particle-size reduction, carbon coating, ion 4 x 1-x 4 doping, and an optimized synthesis process. Reducing particle size to a nanometer scale can efficaciously enhance the electrochemical properties of phospho-olivine materials, such as nanoparticles [119] and nanosheets [120–122], because nanoparticles can provide shorter diffusion pathways and larger surface-area contact with electrolytes for electron and ion transfer [123]. Particle size and shape may be greatly controlled via synthesis methods. Common preparation methods include the solid-phase method, sol-gel method, hydrothermal/solvothermal method, precipitation method, spray pyrolysis method, polyol synthesis and so on. The solid-phase method is a conventional and economical synthesis method, which heats a mixture of lithium, man- ganese and phosphorus sources at a high temperature to form olivine-phase products [124] This method can be easily scaled up for commercial use. However, the solid-phase method not only needs a high-temperature environment, but also produces large or agglomerated particles with poor electrochemical performance [125]. In comparison to the solid-phase method, the liquid-phase method is beneficial for controlling particle morphology and synthesizing nanoparticles. For example, powders produced using the sol-gel method have high purity, uniformity, increased crystallinity, accurate stoichiometric control, and minimal size [126]. Sol-gel synthesis is a low-temperature wet chemical method which involves the formation of sols, the gelation of sols into gels, the drying of gels into “xerogels” with reduced volume, and the densification to obtain final powder products [127]. The sol-gel method can adjust the product’s structure and morphology within the nanometric range, by controlling reaction time, pH value, calcination temperature, concentration, viscosity and so on. Kwon et al. [128] studied the influence of calcination temperature on particle size and successfully synthesized LiMnPO nanoparticles with a diameter of 130 10 nm via a glycolic-acid-assisted sol-gel approach and this sample showed a reversible capacity of 134 mAh g at 0.1 C. Liu et al. [129] synthesized carbon-coated nano-sized LiMnPO and LiMn Fe PO /C of 100–150 nm in width and 200–400 nm in length using a high-energy 0.5 0.5 4 ball-milling-assisted sol-gel method (Figure 6b,c). The as-prepared LiMn Fe PO /C 0.5 0.5 4 attained considerable electrochemical performance with initial discharge capacities of 128.6 mAhg and capacity retentions of 93.5% and 90.3% after 100 cycles at 1 C and 2 C rates, respectively (Figure 6d). Although nanoparticles provide shorter diffusion lengths and better electrochemical performances, they have low tap density because of their high surface area, which causes low tap density and loading of active material in electrodes [130]. Since spherical particles have the largest tap density, it is crucial to control the particle morphology. The hydrothermal/solvothermal method can produce homogenous and mor- phology controllable particles. Hydrothermal/solvothermal method involves wet chemical Batteries 2023, 9, x FOR PEER REVIEW 16 of 23 obtain final powder products [127]. The sol-gel method can adjust the product’s structure and morphology within the nanometric range, by controlling reaction time, pH value, cal- cination temperature, concentration, viscosity and so on. Kwon et al. [128] studied the influence of calcination temperature on particle size and successfully synthesized LiMnPO4 nanoparticles with a diameter of 130 ± 10 nm via a glycolic-acid-assisted sol-gel −1 approach and this sample showed a reversible capacity of 134 mAh g at 0.1 C. Liu et al. [129] synthesized carbon-coated nano-sized LiMnPO4 and LiMn0.5Fe0.5PO4/C of 100–150 nm in width and 200–400 nm in length using a high-energy ball-milling-assisted sol-gel method (Figure 6b,c). The as-prepared LiMn0.5Fe0.5PO4/C att ained considerable electro- −1 chemical performance with initial discharge capacities of 128.6 mAhg and capacity re- tentions of 93.5% and 90.3% after 100 cycles at 1 C and 2 C rates, respectively (Figure 6d). Although nanoparticles provide shorter diffusion lengths and bett er electrochemical per- formances, they have low tap density because of their high surface area, which causes low tap density and loading of active material in electrodes [130]. Since spherical particles have Batteries 2023, 9, 246 the largest tap density, it is crucial to control the particle morphology. The hydrot 16her- of 23 mal/solvothermal method can produce homogenous and morphology controllable parti- cles. Hydrothermal/solvothermal method involves wet chemical processes that take place in an aqueous solution of mixed precursors above the solvent’s boiling temperature [131]. processes that take place in an aqueous solution of mixed precursors above the solvent’s Cao et al. [132] synthesized LiMnPO4 particles with good sphericity and high tap density boiling temperature [131]. Cao et al. [132] synthesized LiMnPO particles with good spheric- using the hydrothermal method with Li3PO4 as the precursor (Figure 6e). This sample ex- ity and high tap density using the hydrothermal method with Li PO as the precursor 3 4 hibited excellent cycling performance with capacity retention of 95.55% after 500 cycles (Figure 6e). This sample exhibited excellent cycling performance with capacity retention (Figure 6f). Luo et al. [133] synthesized a LiMn0.8Fe0.2PO4/C nanocrystal using a facile sol- of 95.55% after 500 cycles (Figure 6f). Luo et al. [133] synthesized a LiMn Fe PO /C 0.8 0.2 4 vothermal reaction and studied the transformation law of morphology from nanosheet to nanocrystal using a facile solvothermal reaction and studied the transformation law of nanoellipsoid. By modifying pH value and precursor ions, nanoellipsoid S-2.6 delivers morphology from nanosheet to nanoellipsoid. By modifying pH value and precursor ions, excellent cycling performance and chemical stability. nanoellipsoid S-2.6 delivers excellent cycling performance and chemical stability. Figure 6. (a) Crystallographic structures of LiMnPO4. SEM images (b), TEM images (c) and cycling Figure 6. (a) Crystallographic structures of LiMnPO . SEM images (b), TEM images (c) and cy- performance (d) of nano−sized LiMn0.5Fe0.5PO4/C [129]. Process diagram (e) and cycling perfor- cling performance (d) of nano sized LiMn Fe PO /C [129]. Process diagram (e) and cycling 0.5 0.5 4 mance (f) of LiMnPO4/C synthesized by hydrothermal method [132]. performance (f) of LiMnPO /C synthesized by hydrothermal method [132]. Ion doping at Li, Mn, and O sites can also enhance the electrochemical performances Ion doping at Li, Mn, and O sites can also enhance the electrochemical performances of LiMnPO4. Li-site doping can reduce the charge transfer resistance and broaden the one- of LiMnPO . Li-site doping can reduce the charge transfer resistance and broaden the dimensional diffusion channels of Li [134], but the transition metal in the Li layer will one-dimensional diffusion channels of Li [134], but the transition metal in the Li layer hinder Li diffusion to a certain extent [135]. Thus, Mn-site doping has received a lot of will hinder Li diffusion to a certain extent [135]. Thus, Mn-site doping has received a lot att ention, such as doping with Fe [136], V [137], Mg [138], Ni [139], Cu [140], Cr [141], and of attention, such as doping with Fe [136], V [137], Mg [138], Ni [139], Cu [140], Cr [141], Zn [135]. Oukahou et al. [136] synthesized LiMn1-xMxPO4 (M = Ni, Fe) with improved elec- and Zn [135]. Oukahou et al. [136] synthesized LiMn M PO (M = Ni, Fe) with improved 1-x x 4 + + 2+ 2+ 2+ 2+ tronic conductivity and reduced Li diffusion energy barrier, by doping Ni and Fe cat- electronic conductivity and reduced Li diffusion energy barrier, by doping Ni and Fe io cations ns at M at n Mn sites sites. . Hu e Hu t alet . [142] synthesized Fe-doped LiMnPO al. [142] synthesized Fe-doped LiMnPO 4 nanoparticles th nanoparticles rough through the solvothermal method, and found that Fe doping can significantly increase the initial the solvothermal method, and found that Fe doping can significantly increase the initial reversible capacity, cycle performance and rate capacity. LiMn Fe PO exhibits a high 0.5 0.5 4 –1 discharge capacity of 147 mAh g and nearly 100% capacity retention after 100 cycles at 1 C. Anion doping at O sites can also enhance the electrochemical properties of LiMnPO by facilitating Li migration of lithium ions in the diffusion channels and enhancing electronic conductivity [134]. Zhang et al. [129] prepared Sulphur-doped LiMn Fe PO via a 0.5 0.5 4 one-step solvothermal process. Since doping Sulphur with a less electronegative atom is advantageous for increasing electronic conductivity, the as-prepared cathode material delivers a high specific discharge capacity of 166.83 mAh g at 0.1 C. Carbon coating is a common method to increase the conductivity of phospho-olivine materials by enhancing the electron conductivity between particles [143] and improving the contact between the active material and liquid electrolyte. The most important thing for carbon coating is to select high-quality and low-cost carbon sources. Mizuno et al. [144] studied the discharge capacities of LiMnPO coated with carbon prepared from differ- ent carbon sources and found that LiMnPO with carbon obtained from carboxymethyl cellulose exhibits the highest discharge capacity of 94 mAh g at 0.01 C. However, this method needs additional heat treatment to form a carbon coating. The polyol method can efficaciously prepare nanosized particles with in-situ carbon coating. The polyol method Batteries 2023, 9, 246 17 of 23 includes four main steps, in which the polyol acts as solvent, fuel, energy supplier and carbon source, respectively [145]. Long et al. [146] synthesized LiMnPO with a carbon coating using a microwave-assisted polyol method at 130 C for 30 min. The obtained LiMnPO /C sample contains a 2 nm thick carbon layer and delivers a discharge capacity of 126 mAh g with a capacity retention of ~99.9% after 50 cycles at 1 C. By using a two-step mechanochemically assisted solid-state synthesis, Podgornova et al. [147] syn- thesized carbon-coated LiFe Mn PO cathode materials. They found that the capacity 0.5 0.5 4 and rate capability of LiFe Mn PO improved owing to the good graphitization of the 0.5 0.5 4 carbon coating and the tight combination between the carbon coating and LiFe Mn PO . 0.5 0.5 4 Although these aforementioned methods can enhance the electrochemical performances of LiMnPO and LiFe Mn PO , obtaining all these desired properties is still challenging. 4 0.5 0.5 4 To prepare LiMnPO and LiFe Mn PO cathodes for high-performance lithium-ion bat- 4 0.5 0.5 4 teries, a simple, economic, easy to control, and reliable synthesis technique still needs to be developed. 4. Summary and Perspectives Mn-based materials have shown advantages in material cost and energy density among LIBs’ cathode materials, due to their richness in the Earth’s crust and wide range of valance states. This has led to the generation of series of potential cathode materials with low cost per Wh, such as spinel LiMn O and LiN Mn O , layered LiNi Mn O 2 4 i0.5 1.5 4 0.5 0.5 2 and Li-rich Li Ni Mn O , rock-salt Li Mn Nb O F , Olivine LiMnPO and 1.2 0.2 0.6 2 1.2 0.625 0.175 1.95 0.05 4 LiMnFePO , etc. However, the Mn dissolution and Jahn–Teller effect of Mn-based materials are not negligible, which is detrimental to electrochemical performance. Element doping and surface coating are the major strategies to solve the above problems, which can effica- ciously promote the stability and kinetics of Mn-based materials, then improve the cycling performance and rate capability. Some special problems need special improvement strate- gies; for example, the oxygen release of Li-rich materials can be eliminated by constructing oxygen vacancy and the stability of DRX compounds can be improved by F fluorination. Among the materials mentioned in this review, spinel LiMn O has been success- 2 4 fully commercialized, while layered LiNi Mn O , olivine-type LiMnPO and LFMP, as 0.5 0.5 2 4 well as LRM have the potential to be applied as cathode materials for LIBs in the near future. The applications of LiN Mn O and DRXs, however, are limited due to their i0.5 1.5 4 high discharge potential and poor cycle performance, respectively. For the Mn-based cathode materials, some prospects are suggested: (1) achieving a uniform coating by simple methods, (2) introducing doping elements into the precursors, (3) producing high-voltage electrolytes with low cost, and (4) analyzing the synergistic effect of multi-element doping and the influences of different components on the cathode’s electrochemical properties, via theoretical calculations. 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