Open access peer-reviewed chapter

Advanced Chalcogen Cathode Materials for Lithium-Ion Batteries

Written By

Varishetty Madhu Mohan, Madhavi Jonnalagadda and VishnuBhotla Prasad

Submitted: 18 December 2021 Reviewed: 03 February 2022 Published: 09 May 2022

DOI: 10.5772/intechopen.103042

From the Edited Volume

Chalcogenides - Preparation and Applications

Edited by Dhanasekaran Vikraman

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As on today the main power sources of lithium-ion batteries (LIBs) research developments gradually approach their theoretical limits in terms of energy density. Therefore, an alternative next-generation of power sources is required with high-energy densities, low cost, and environmental safety. Alternatively, the chalcogen materials such as sulfur, selenium, and tellurium (SSTs) are used due to their excellent theoretical capacities, low cost, and no toxicity. However, there will be some challenges to overcome such as sluggish reaction of kinetics, inferior cycling stability, poor conductivity of S, and “shuttle effect” of lithium polysulfides in the Li-S batteries. Hence, several strategies have been discussed in this chapter. First, the Al-SSTs systems with more advanced techniques are systematically investigated. An advanced separators or electrolytes are prepared with the nano-metal sulfide materials to reduce the resistance in interfaces. Layered structured cathodes made with chalcogen ligand (sulfur), polysulfide species, selenium- and tellurium-substituted polysulfides, Se1-xSx uniformly dispersed in 3D porous carbon matrix were discussed. The construction of nanoreactors for high-energy density batteries are discussed. Finally, the detailed classification of flexible sulfur, selenium, and tellurium cathodes based on carbonaceous (e.g., carbon nanotubes, graphene, and carbonized polymers) and their composite (polymers and inorganics) materials are explained.


  • sulfur
  • selenium
  • tellurium (SSTs) electrodes
  • polysulfides
  • nano-metal chalcogen and flexible batteries
  • nanoreactors
  • electrochemical properties

1. Introduction

Recently, much attention has been focused on the development of more safe, high-energy density, long-life, and low-cost batteries to satisfy our energy demand as our daily life includes electric vehicles, portable electronics, and large-scale grids [1, 2, 3, 4]. However, lithium-ion batteries (LIBs) successfully prepared and available in the commercial market since the 1990s, even though their theoretical specific capacity, energy density of the electrode material is relatively low. Hence, still it is challenging to development of next generation of lithium-ion batteries to fulfill the demand by overcoming their hindrance by the intrinsic limitation [5].

Various metal sulfides, metal oxides, and metal poly-anions have been developed, and they exhibited preferable capacities, but their output voltages are not sufficient. Hence, chalcogen materials (sulfur, selenium, and tellurium (SSTs) are emerging conversion-type cathode materials for aluminum-ion batteries (AIBs) [6]. They exhibit the multi-electron transfer process; therefore, AlSST batteries can deliver very high capacities.

The low-cost and -toxicity, most-abundant sulfur (S) has theoretical gravimetric capacity of 1672 mA h g−1 due to the reaction between S and sulfide (S2−) exhibiting the highest capacity, over all cathode materials in AIBs [7]. Notably, lithium-sulfur (Li-S) batteries exhibited the highest theoretical energy density of 2600 Wh kg−1. Moreover, the sulfur-based cathodes require a large number of conductive additives due to their electrical insulating property to avoid leakage/loss issues, which greatly decreases the actual capacity [8]. Therefore, it is usually difficult to achieve the theoretical specific capacity. Based on these merits, metal chalcogens batteries (MCBs) are an attractive interest for alternative of new-generation secondary batteries.

The main research has been focused on lithium dendrite and “shuttle effect” of high-order lithium phosphate sulfur batteries. There is three-phase boundary between cathodes, electrolytes, and modified interface layers. The interface boundaries occupied by the sulfur species restrict the redox kinetics. Therefore, to improve the redox reaction, it is necessary to make strong bond of S to the host, thereby continuing to conduct mixed charge carriers (Li+ and e). There is another difficulty that the degradation of active materials between the electrodes due to shuttle process, in which cathode electrolyte interfaces bring rapid decay of the capacity, thereby reducing the coulomb efficiency of Li-S batteries [2]. The formation of lithium dendrites on the surface of lithium anode and also unstable solid electrolyte interface (SEI) leads to low columbic efficiency (CE) and poor cyclic performance [9]. To overcome these problems, most suitable nanomaterials are considered for energy applications due their unique crystal structure providing high surface-area-to-volume ratio and shortening lithium ions transport [9].

In order to create the metal-ligand covalency is the one of strategies by replacing the oxide ligand with the chalcogen (S, Se) to achieve an anion redox stabilization, where the less electronegative nature of the chalcogen improves the ligand p band penetration into the metal d band. Tarascon et al. investigated layered chalcogen structures as well as their electrochemical performance for the next generation of cathodes [10]. The results exhibited the superior performance in voltage and capacity fade with voltage hysteresis. Hence, chalcogen anion redox plays a critical role in a Li-rich cathode batteries. Some of research was carried out on chalcogen cathodes, both Li-rich and conventional chalcogen cathodes for the evolution of chalcogen anion redox cathode [11, 12]. Mespoulie et al. [13] introduced fast Li-ion conductors of mixed anionic and cationic redox activity of Li2SnS3, by introducing the Fe redox couple in the host cation Sn site.

Selenium (Se) is another class of material present in Al-Se battery that showed higher voltage plateau resulting a desirable energy density [14]. Selenium metal has high theoretical volumetric capacity (3253 mA h cm−3, ρ = 4.81 g cm−3), which is more suitable especially in hybrid electric vehicles and in the mobile smart phones due to restrictions of the battery volume [15]. Selenium showed higher electronic conductivity (1 × 10−3 S m−1) and excellent kinetic behavior than sulfur [16]. The chemical compound Se1-xSx with different Se-S ratios shows higher theoretical capacity as well as better electronic conductivity due to fast reaction of kinetics than pristine S [17]. Even though, Se1-x Sx cathode materials also suffer from poor cycle performance, lower coulombic efficiency due to the dissolution and shuttling of intermediates [18]. The electrochemical performance of Li-Se1-xSx batteries improved by carbon coating, which provide a strong chemical affinity of polarized surface, which can effectively trap the soluble intermediates to minimize the shuttle effect and side reactions in the electrolyte [19]. A series of Se1-xSx cathodes were prepared by Se/S ratio and the presence of supercritical CO2. NC@SWCNTs@Se1-xSx cathodes exhibited higher conductivity and strong adsorption leading to superior cyclic efficiency.

Tellurium (Te) material has the highest atomic weight among sulfur and selenium, high electrical conductivity and a 6-electron transfer reaction process made it to be promising cathode material in AIBs [20]. However, still there are several challenges remaining to overcome for the development of batteries such as low electronic conductivity of S, shuttle effect, slow kinetics of ionic liquids as well as undesirable reaction mechanism [21]. Tellurium exhibits a higher theoretical volumetric capacity of 2619 mAh cm−3 due to its intrinsic electrical conductivity of Te (2 × 102 S m−1), much better than that of S (5 × 10−16 S m−1) and Se (1 × 10−4 S m−1). Therefore, the high utilization ratio of active material of Te leads to good performance at the large current density. The fabricated batteries based on Te/porous carbon (Te-G-CNT) electrode materials deliver a high volumetric capacity up to 2493.13 mAh cm−3 [22].

The polysulfide (Li2Sn) species have strong tendency to catenate and form reactive polysulfide dianions as well as radical anions (Sn2− and Sn/n/2, 2 < n < 8). These conversion reactions of sulfur ↔ Li2S kinetically favored in the mediated solution and their deposition degrading the lithium surface and the cyclic stability [23, 24, 25]. The dissolved species shows shuttle effect by insulating deposition of Li2S/Li2S2. Polysulfide molecules modified by substituting chalcogen atoms minimized the intrinsic shuttle effect [26]. By substituting S, Se and Te can be facilely formed as the polyselenosulfides (Li2SexSy) and polytellurosulfides (Li2TexSy). However, selenium and tellurium lead to significant differences in the electrochemical performance compared with Li-S batteries. The substitution of selenium and tellurium has significant impact on the metal-chalcogen batteries and solid-state batteries by employing chalcogenide solid electrolytes. Therefore, in this chapter, the strategies to improve electrochemical performance are elaborated, and the development of new trends for next-generation lithium-ion batteries is provided.

To fabricate flexible lithium-ion batteries using sulfur-based cathodes, there are two main synthetic approaches: (1) Post-sulfur loading: The formation of a flexible skeleton then loaded with sulfur by using vapor infusion, melt diffusion, or reprecipitation of sulfur from a solution (generally carbon disulfide (CS2) or toluene). (2) Pre-sulfur loading: pre-synthesized sulfur composites into a flexible cathode. By keeping the flexible cathode required features in mind such as (1) high content of active materials with respect to total mass of the electrode, (2) mechanically robust skeleton, (3) long-range interpenetrated conductive network, (4) porous structure, and (5) three-dimensional (3D) scaffold to improve areal sulfur loading [27]. The flexible energy Li-S batteries, flexible alkali metal-chalcogen batteries, and two special flexible batteries such as prototypes of foldable and cable-type Li-S batteries are discussed.


2. Results and discussion of various topics

2.1 Sulfur, selenium and tellurium batteries

Many efforts have focused on the cathode material design, electrolyte optimizations, separator modification, still some of the challenges remain due to slow kinetics, electrolyte compatibility, and inferior cycling stability. Hence, there are many possibilities for the development of more reliable sulfur, selenium tellurium (SSTs) batteries. He et al. [5] showed schematic representation of various components for the development of lithium-ion batteries based on SST as shown in Figure 1. The carbon-based materials, conductive materials, and their nanostructure with a porous matrix refereed as a host due to low electrical conductivity of SSTs (S and Se) and the soluble properties of the chalcogenide. Therefore, the host materials can provide necessary contact with SSTs to reduce the formation of inactive regions and satisfy the adsorbing as well as accommodation of soluble active materials. Therefore, the mass loading of SSTs cathodes can efficiently be increased. The various approaches such as melting diffusion, chalcogen vaporization are used to increase the mass loading of active SSTs into the conductive host materials. An introduced of metal atoms to form a bonding with SSTs is another possible method for reducing the reaction barriers in Al-SSTs batteries.

Figure 1.

Perspective for Al-SSTs batteries. Reprinted from ref. [5].

2.1.1 Al-sulfur batteries

The most abundant high-surface-area carbon porous materials are possible to absorb the SSTs materials, which is more impartment for limiting chalcogenide dissolution [28]. Thus, well-designed porous structure carbon materials composites can enhance not only the charge transport but also improve the retention of SSTs cathode during electrochemical reaction [29]. Therefore, carbon materials are attractive to be host for the insulating S with a regular matrix. A melt-diffusion method was conducted to prepare the S/activated carbon cloth (ACC) composite cathode material in Al-S battery as shown in Figure 1a. As prepared ACC material exhibited type I adsorption, corresponding to microporous structure with a pore size below 2 nm. The Brunnauer-Emmett-Teller (BET) measurement showed that after compositing of S into the ACC, specific surface area decreased from 2376.6 to 1532.8 m2 g−1 and also decreased to its micropore volume from 0.93 to 0.61 cm3 g−1, indicating that the S material was uniformly impregnate into the microporous structure. The Al-S battery based on S/ACC cathode exhibited a high specific capacity of 1320 mA h g−1 and discharged voltage about 0.65 V as shown in Figure 2b.

Figure 2.

The galvanostatic charge/discharge curves of a, b) Al-S battery, c, d) Al-Se battery (n ≥ 1), and e, f) Al-Te battery. Reprinted from ref. [5].

The microporous structure ACC host consists of pore size less than 2 nm and can effectively provide the fast solid-state reaction kinetics favoring to its ready electron access, large reaction area, and decreasing the ionic diffusion length. Similarly, a free-standing CNF host was also introduced in Al-S battery as shown in 2b. The carbon nanofiber with a diameter of 100–200 nm occupies between the interspaces in micrometer scale level. This CNF structure provides a spacious, robust, conductive matrix to accommodate the active S and their products. The S and EMIC/AlCl3 slurry dispersed into the freestanding CNF host, the Al-S battery exhibited a good capacity of 1250 mAhg−1. These above free-standing carbon materials not only provide a conductive matrix for S materials, but also reduce the side reaction from the binder, thereby enhanced stability of Al-S battery. The porous carbonized Cu-based metal organic frameworks (MOFs) called as HKUST-1-C also introduced as a host to the S in Al-S batteries [30]. The HKUST-1-C carbon materials exhibited high hierarchical porous structure with surface area of 179 m2 g−1 and a pore sizes in <5 nm range. These are more suitable for being host in S cathode batteries. The metallic Cu can react with polysulfides to form S-Cu ionic clusters, thereby reducing the kinetic barrier of the electrochemical conversion reaction and facilitating the reversibility of S during charge/discharge processes. Therefore, the S@HKUST-1-C cathode battery exhibited a stable performance with a reversible capacity of 600 mA h g−1 at the 75th cycle and retained 460 mA h g−1 even after 500 cycles at 1 A g−1. These results indicate that the metallic material provides a valuable strategy to develop stable Al-S batteries. A nitrogen-doped hierarchical porous carbon called as HPCK used as a host for S in Al-S batteries [31]. Hierarchical micro-, meso-, and macro-pores of HPCK was prepared by carbonizing a N-rich polymer precursor combined with zinc nitrate and followed by a KOH etching process. Interestingly, KOH etching process greatly improved the surface area (2513 m2 g−1) and created more micro- and meso-pores. The Al-S batteries based on S/HPCK cathode delivered a capacity of 1027 mA h g−1 at 0.2 A g−1 for 50 cycles and exhibited excellent cyclic ability of 405 mA h g−1 at 1 A g−1 for 700 cycles. The hierarchical porous structure of HPCK with high surface area and large pores confined S materials. The huge macropores provide fast ion transport in the electrolyte. Therefore, the porous carbon powders, metal content along with structure optimization for host are impartment factors for Al-S batteries.

2.1.2 Al-selenium batteries

Selenium composite prepared with carbon meso-porous material (2 nm < pore size <50 nm) can enhance the encapsulation of Se in the Al-Se batteries as shown in Figure 2C [32]. The CMK-3 carbon demonstrated that it exhibited a hexagonal meso-porous structure favorable for being host material in Al-Se batteries as shown in Figure 2c. The CMK-3 carbon nanorods showed a large pore volume of 1.78 cm3 g−1, a high surface area about 1632 m2 g−1, and a uniform pore size of 3.4 nm. The Al-Se battery based on Se@CMK-3 cathode delivers an initial discharge capacity of 218 mA h g−1 at 100 mA g−1 and a relatively high discharged potential above 1.5 V as shown in Figure 2d. Hollow selenium carbon nanotube (Se@CT) with a specific surface area of 61.49 m2 g−1 and pore diameter of 3.36 nm was also demonstrated as cathode material in Al-Se batteries [6]. The Se@CT cathode exhibited an initial discharge capacity of 447.2 mA h g−1 at 200 mA g−1 with a voltage of about 1.6 V. The capacity retains 83.5% even after the 200 cycles at 500 mA g−1. The carbon materials in Se@ CT cathode reside their structural stability of Se, reduce the dissolution of selenide products, and also avoid the volume change of Se during the electrochemical process. Therefore, mesoporous structure and pore size are important to stabilize the Se cathode as well as enhance the electrochemical performance. Mesoporous carbon fibers (MCFs) size from 2.7 to 8.9 nm prepared and demonstrated the effect of pore size on the electrochemical performance of Al-Se batteries [32]. The Al-Se battery based on MCFs material with pore size of 7.1 nm exhibited a good capacity of 366 mA h g−1. The chloroaluminate ion diffusivity greatly affects in the mesopore size of MCFs composite Se cathode. The carbon well-designed structures may provide the chloroaluminate ion transportation as well as charge transportation during charge/discharge processes, which enhances the electrochemical behavior of Al-Se batteries.

2.1.3 Al-Te cathode batteries

The electrical conductivity of Te cathode is very high [12]. Hence, utilization ratio of active material is large and also exhibited good rate performance in Al ion batteries. Te cathode is easily prepared by coating the slurry of Te powder, acetylene black, and binder on current collector, without any host materials as shown in Figure 2e [20]. The Al-Te battery exhibited with raw material an initial capacity of 913 mA h g−1 at 20 mA g−1 with a potential of 1.4 V as shown in Figure 2f. The Al-Te battery delivers a good rate performance at different current densities due to the high electrical conductivity of Te. However, Te batteries exhibited capacity fading due to its leaching of soluble telluride from the cathode. Further, the rGO materials were introduced to encapsulate Te nanowires in Al-Te battery [6]. This Te/rGO cathode battery exhibited a capacity of 1026 mA h g−1 at 500 mA g−1 and also delivers a considerable capacity beyond 100 cycles at 1.0 and 2.0 A g−1. The rGO materials suppress the dissolution of telluride into electrolyte indicating better utilization of Te. Further, N-doped porous carbon materials coupled with rGO also introduced as a host materials for the improvement of the stability of Al-Te batteries [33]. The rGO materials are easily encapsulated soluble tellurium species under physical and chemical confinements. Therefore, the Al-Te batteries exhibited excellent cyclic ability. It exhibited initial specific capacity of 935.5 mA h g−1 and 467.5 mA h g−1 after 150 cycles with the Te loading of 70 wt%. Thus, a host material with well-designed structure, such as porous conductive matrix with specific components is necessary for cathode materials in Al-SSTs batteries.

2.2 Interfacial problems in sulfur batteries

In the case of sulfur lithium-ion batteries, during the discharge/charge reaction, the sulfur required to be tightly attached to a host with sustainable conduction of Li+ and e. Generally, the cathodic reaction occurs at the host/elemental sulfur/liquid electrolyte interface. The carbon nano materials such as graphene, carbon nanotube (CNT), or carbon nanofiber and metal sulfide are the indispensable 2D materials to sulfur host. The modified host materials with only nano-sized pores alone cannot accommodate the sulfur and completely reduce suppression of the shuttle effect in the LiPS [34]. Therefore, in recent years the development of interface components for Li-S batteries is most impartment [28, 35]. One of the main strategies is that the coating of materials exhibited several merits such as fast electrical and ionic transmission capability, uniform thickness, and stable distribution of composition on the surface of cathode materials without effect of volume.

The another most prominent strategy is that in situ growth of nano transition metal oxides (sulfide) [36] or the loading of nano transition metals [37, 38] on the carbon surface leads to overcome the poor contact between sulfur and carbon materials. In this process, (sulfides, oxides, nitrides, etc.) compounds are added or doped some of the elements (N, S and co-NS) and their derivatives. Therefore, polar bonds are generated between host and sulfur; those provided fast transmission of electrons as well as increase the ions redox reactions at the interface.

2.2.1 Inter facial problems in nano metal sulfides (oxide)

Various nano metal oxides have interacted with LiPS through strong chemical bonds, which are reducing the shuttle effect. In particular, oxygen-rich compounds such as V2O3 [39], TiO2 [40, 41], SnO2 [42], Co3O4 [43], MnO2 [44, 45]) successfully prepared as LiPS traps to enhance the cyclic stability. The high conductive Ni foam/graphene/carbon nanotubes/MnO2 nanoflakes (NGCM) were proposed in which interconnected Ni foam, graphene, and carbon nanotubes of the NGCM sponge facilitated efficient electron transfer. The NGCM sponge showed good wettability and interfacial contact with the Li-S electrolyte, and the MnO2 nano flakes exhibited electro-catalytic effects as well as strong chemisorption on LiPS [46]. The porous and double-shelled architecture decreases the ion transfer distance, Uniform sulfur distribution offers active interfaces as well as decreases volume changes. Luo et al. developed spinel Ni-Co oxide double-shelled microspheres (NCO-HS), which consisted of defective spinel NiCo2O4–x, as the multifunctional sulfur host material. The S@NCO-HS prepared under high sulfur loading exhibited minimum capacity fading rate of 0.045% per cycle over 800 cycles with high areal specific capacity of 6.3 mAh cm−2 at 5 C.

2.2.2 Capping layer of MTO-CNTs for sulfur cathode

The thin barrier layers designed with light weight and with high polysulfide-trapping capabilities showed high weight density (usually >0.3 mg cm−2), unexpectedly reducing the overall energy densities of Li-S batteries as shown in Figure 3 [47, 48, 49, 50, 51, 52, 53, 54, 55]. The development of lightweight MTO-CNTs capping layer directly coated onto the surface of sulfide cathode as shown in Figure 4a. The MTO-CNTs capping layer prepared on the sulfur cathode, which is directly contacted with an electrically conducting layer to form a cathodic “sub cell” for capturing and decreasing the polysulfide species. The thickness (~2 μm) and the weight density (0.06 mg cm−2) of the MTO-CNTs capping layer are much lower than other barrier layers reported elsewhere, as the mesoporous carbon and the grapheme layers [47, 48]. Moreover, it is noted that the area of conventional interlayer is higher than the coated capping layer. To understand the layer-by-layer electrode structure and the cathode structure was observed by SEM and energy dispersive X-ray spectroscopy (EDX). The 1D, MTO-CNTs nanostructure and the ultrathin capping layer are formed through self-weaving and firmly coated on the surface of sulfur cathode (Figure 4b). The existence of MTO-CNTs capping layer is estimated by the elemental mapping analysis (Figure 4c,d). The MTO-CNTs capping layer formed on the cathode can trap more effectively, polysulfides within the cathode material, thereby reducing the shuttle effect of polysulfide. The CV curves of electrodes reveal that the electrode with MTO-CNTs capping layer has sharper and more intensive oxidation and reduction peaks exhibited than an electrode without capping layer, indicating that the capping layer efficiently enhanced reaction kinetics of the electrode. The substantially promoted charge transfer is further confirmed by EIS analysis. The sulfur cathode with MTO-CNTs capping layer presents much lower charge-transfer resistance compared with cathode without MTO-CNTs capping layer, indicating that improvement of redox-conversion ability as well as conductivity. These results indicate that the capping layer is not only favorable for adsorption confinement of polysulfides within the electrode, but also increasing charge transfer, accelerating reversibility of polysulfide conversion. The galvanostatic charge-discharge profiles of the device with the capping layer were recorded at different rates as shown in Figure 4e. At the 0.2 C rate reversible capacity of 1212 mAh g−1 has been achieved. The specific capacities of 922 and 606 mAh g−1 were delivered at the high rates of 0.5 and 1 C, respectively. The cycling stability of the device with the capping layer at 0.5 C rate still it retains a capacity of 577 mAh g−1 after 500 cycles with capacity decay rate of 0.07% per cycle (Figure 4f), indicating a good cycling stability. These results indicate that the formation of MTO-CNTs capping layer is convenient route to fabricate high performance Li-S batteries with sulfur host along with commercial carbon materials.

Figure 3.

Schematic diagram of as fabricated Li-S battery using MTO-CNTs interlayer. Reprinted from ref. [41].

Figure 4.

(a) The dripping of MTO-CNTs on the surface of sulfur electrode. (b) SEM images of the sulfur electrode with MTO-CNTs capping layer. (c) Corresponding Ti and (d) S elemental mapping. (e) Charge-discharge curves of sulfur electrode comprises MTO-CNTs capping layer at various C rates. (f) Cycling stability of the MTO-CNTs capping layer on sulfur electrode at 0.5 C. [reprinted from ref. [41].

The cell was disassembled after 100 cycles at 0.2 C, to understand the function of the MTO-CNTs capping layer at a potential of 2.8 V. The dimethyl carbonate solution was used to wash electrodes with capping layer and their structure analyzed by X-ray microtomography (XRM) and SEM. The overall structure and morphology of the MTO-CNTs are similar to that of the original sample as shown in Figure 5a, the SEM images of the capping layer after 100 cycles. The TiO2 volume cannot change due to its robustness and the interaction of the CNTs. The layer-by-layer stacked structure was indicated by a 3D reconstruction of the electrode (Figure 5b). The capping layer is uniformly and strongly anchored on the sulfur cathode surface and worked as a good absorbent layer to keep polysulfide species rather than to diffuse into the lithium anode. The signal of sulfur precipitate is very strong and uniformed in the cathode (Figure 5c), reveals that the polysulfide shuttle behavior retained by the MTO-CNTs capping layer. The charge products are clearly seen in the capping layer execute the recycling of trapped polysulfides. Figure 5d reveals that the capping layer on surface of the sulfur cathode retain as it is without any cracks after cycling indicates that the structural stability. These results indicate that the light MTO-CNTs capping layer coated on the surface of sulfur cathode enhanced battery performances.

Figure 5.

(a) SEM picture and (b) 3D XRM picture and (c, d) partial 2D of the sulfur electrode with MTO-CNTs capping layer at a potential of 2.8 V after 100 cycles at 0.2 C. reprinted from ref: [41].

2.3 Multi (mixed cationic and anionic) redox in chalcogen cathode for Li-ion batteries

In order to improve the metal-ligand covalency by substituting a lower electronegative chalcogen ligand such as sulfur in the cathode. In this case, reversible mixed anionic and cationic redox occurs by the metal d band penetrates into the ligand p band. The partially filled d orbital redox couples like Fe2+/3+ are introduced in the Li-ion conducting phase (Li2SnS3) is the development of a new family of layered structured cathode materials. Investigation of high-resolution transmission electron microscopy and high annular dark field-scanning transmission electron microscopy reveals that the multi-redox structural modifications and nanopore formation on its surface, during cycling process. In this study, Ni, Co ions free cathodes using various functional materials in the chalcogen-based dual anionic and cationic redox cathodes.

2.3.1 Electrochemical kinetic study

The charge-discharge profiles clearly exhibited a large hysteresis (0.511 V at 50% state of charge) starting from the first cycle. A voltage hysteresis plot for the 1st–30th cycles is shown in Figure 6a. Further to understand the hysteresis clearly by analyzed the effect of the upper cutoff potential on the voltage hysteresis as shown in Figure 6b. But these results have no significant effect on the voltage hysteresis, indicating that the hysteresis in the Sn-based chalcogen system may be due to the origin of intrinsic nature of the Li insertion extraction reactions. Therefore galvanostatic intermittent titration technique (GITT) measurements were carried out with current density of 5 mA/g for an interval of 30 min. Then it allowed opening circuit state for 1 hour, to obtain a steady state. Two continuous cycles of the composite electrodes by the GITT measurements are shown in Figure 6c,d. The first charge process represents the Low over potential increment gradually and the discharge process battery delivered a very high potential with different behavior. The 0.2 Fe substituted composite cathode shows the less overpotential with gradual increment is due to the smooth and kinetically faster of Li extraction process compared with the Li insertion process. The huge over potential observed around 2.5–2.1 V during the insertion process reveals that the Li insertion is kinetically limited and slower than the extraction process. In the similar way, the second cycle showed and high over potential during discharge and less over potential during charge condition. Further, the discharge profile region shows the two different slops indicating that due to the multiphase reaction region and the structural modifications. GITT profiles reveals that the inflection point region represents the high potential of ∼187 mV compared to other regions. The electrochemical insertion of Li+ ions is limited in these regions due to the slow Li-ion diffusion and charge transfer resistance.

Figure 6.

Electrochemical kinetic study of the Sn-based chalcogen anion redox cathode (0.2Fe-Li1.33Sn0.67S2). (a) Voltage hysteresis profile of active cathode material at 10 mA/g current density. (b) Voltage profile of the cathode at different upper cutoff potentials. GITT profiles of active cathode material: (c) 1st cycle and (d) 2nd cycle GITT profiles with 30 min pulse and 1 h relaxation. Reprinted from ref. [13].

The Li insertion and extraction was estimated by cycling test for all the compounds in a half-cell configuration at the current density of 10 mA/g. The electrochemical investigation of 0.2 Fe substituted compounds compared with other compositions are shown in Figure 7a. There is a gradual deterioration in all the compounds cyclic performance at a very low current density of 10 mA/g for 50 continuous cycles in terms of capacity fade. Further, a cycling test of the 0.2 Fe substituted cathode done at 50 mA/g, reveals that the high rate cycling stability about 76% retention after 80 cycles as shown in Figure 7b. The multi-redox induced structural transformation is the main reason for capacity degradation and evidenced by the microscopy analysis. The cycling stability of this composite materials is comparable to that of existing chalcogen anion redox cathodes [11, 56]. Further, the nanostructure and surface coatings strategies would enhance the cycle life [57, 58]. In this chalcogen framework the Fe doped composition showed good electronic and ionic conductivity, excellent electrochemical properties with the high loading of 10 mg/cm2 cells at 10 mA/g current density.

Figure 7.

Electrochemical study of Sn-based chalcogen anion redox cathodes. (a) Cycling stability of different Fe substituted Li1.33Sn0.67S2compounds. (b) Cycling stability of 0.2Fe- Li1.33Sn0.67S2 cathode at a high current density of 50 mA/g (initial few cycles at 10 mA/g) reprinted from ref. [13].

2.3.2 HAADF-STEM and HR-TEM pictures

The structural evolution due to Li insertion extraction was visualized in a series of samples using HAADF-STEM images (Figure 8). High atomic number elements such as transition metals represented by the bright spots and the light elements such as Li, S, and O represented without bright spots. Figure 8a shows the ordered pristine cathode composition showed bright spots of metal elements and the lattice exhibits without distortions or cracks. The lattice showed severe structural distortions and nanopore formation after the 1st cycle, indicated by the yellow-colored circles and arrows in the Figure 8cf.

Figure 8.

HAADF-STEM investigation of the 0.2Fe-Li1.33Sn0.67S2composite electrode at different states of charge: (a) pristine, (b) 1st charged, (c) 1st discharged, (d, g) 2nd charged, (e, h) 2nd discharged, (f, i) cycled electrodes; In panel (c), the honeycomb ordering was visualized. Reprinted from ref. [13].

The structural distortion was observed in the first discharged cycle like honeycomb and after consequent cycling the ordered crystalline domain lost its crystalline nature as shown in Figure 8c. Further, investigations reveal that complete distorted structure after cycling due to the pores and structural distortions was high and the crystallinity degraded. The degradation of pore and crack formation happened during the initial cycles to extended cycling conditions. The low magnification images of second charged, second discharged, and cycled samples are shown in Figure 8g,h indicating a lot of nanopores present throughout the cycled cathode. Hence, the nanopore created by accelerated Li ion insertion-extraction and also the sulfur loss by degradation. The formation of nanopore/nanovoid due to oxygen loss in the Li-rich oxide anion redox cathode as well as their degradation mechanisms correlated with the fundamental issues of voltage fade, voltage hysteresis, and capacity fade [59]. The sulfide anion is a significant charge contributor during Li extraction and the sulfur loss increased amorphization which is reflected in charge-discharge voltage profiles by means of capacity fade of the cycled cathode materials.

2.3.3 Impedance spectroscopy analysis

Figure 9a shows the impedance plots of cathode material at different discharged cycles. All the EIS measurements showed two semicircle regions except the pristine cell. The semicircle observed at high frequency region indicates the surface film formation on the positive electrodes. Another semicircle was observed in the low frequency region attributed to the charge transfer resistance upon Li+ insertion-extraction. The observed slope line indicates the Warburg diffusion (W) in the bulk electrolyte. An equivalent circuit model (Figure 9b) designed by the fitted EIS data and the resistance values for all the active materials at different cycles noted. In the equivalent ckt Re corresponds to the solution resistance. The Cs and Rs correspond to and capacitance and the surface film resistance of the active sulfide composite cathode, respectively. The Li ion insertion and extraction double layer capacitance and induced charge transfer resistance represents as Rct and Cdl, respectively. From the Figure 9a calculated Re values for all the materials indicated small differences for all the cycles, reveals that the electrolyte is stable during cyclic process. The Rs and Rct values increased proportionally in the initial cycle number during 1st and 2nd cycles, representing that the structural distortion during the electrolyte and electrode interface. This observation consisted with GITT analysis that the very large potential during lithium insertion of initial cycles is due to the kinetic limitation of Li insertion reactions. The most of the surface film formation will occur during initial cycling. The Rct values decreased with increasing cyclic number, indicating that the resistance increased during the initial cycles and decreased by cycling cyclic number. This is due to the sulfur loss by rising cycling number, where the complete amorphization and nanopore are formed [12]. The anion redox indicates sulfur loss by reducing the Li insertion-extraction ability. These indicate the capacity fading occurred during the cycling. Therefore Sn based chalcogen layered structure materials worked as mixed redox cathode for Lithium Ion batteries.

Figure 9.

EIS of 0.2Fe-Li1.33 Sn0.67S2 composite electrode at different states of charge. (a) Nyquist plots of the composite electrode at various cycles. This EIS was recorded after completion of respective cycle, (b the model EIS was fitted to the experimental Nyquist plot. Reprinted from ref. [13].

2.4 Chalcogen substitusion into the polysulfides for batteries

The cyclic voltammograms (CVs) of half cells assembled using [Li2S + 0.1 S/Se/Te] cathodes as shown in Figure 10a. The presence of polyselenosulfides plays a significant reduction in peak separation (∆Ep), indicating decreasing the overpotentials, and helps with increased peak heights at high scan rates (1 ≥ mV s−1) by retaining the canonical redox peaks of sulfur/Li2S. The relationship between the peak current (ip) and scan rate (ν) can be written as: ip = α ν β, where α and β are the fitting parameters [60] Plotting log (ip) versus log (ν) yields β = 0.64 for Se, compared to 0.52 for the control. The value of β changes with the addition of Se represents the shifting of slow diffusion controlled reactions to fast surface-controlled reactions. This improvement of the redox kinetics is not much with the introduction of tellurium compared with selenium. Figure 10b shows the capacities of Li || [Li2S + 0.1 S/Se/Te] half cells at 0.25 A g−1 of Li2S (BC/5). Selenium is electrochemically more active between 2.8 and 1.8 V, whereas Te is inactive at the same voltage. Hence the addition of 0.1 Se enables a considerable improvement in capacities of ̴40% under the control. Therefore the relative dominance of catalytic SeS2* radical presence in polyselenosulfide solutions works as conversion reactions and utilized completely to drives complete electrochemical reaction [61, 62]. This is critical for high capacities obtained under low-electrolyte conditions in a practical Li–S cell [24, 63]. The presence of SeS2* radicals highly react with the various electrolyte components thereby the faster capacity fade observed with polyselenosulfides even after 70 cycles. The sulfur/Li2S final product conductivity also improved due to incorporation of Se atoms. The charging/discharging profiles of Li || Li2S half cells indicate that considerable reduction achieved in overpotentials with selenium compared to sulfur or tellurium. Therefore, considerable improvement in charge-transfer and redox kinetics is occurred with the introduction of selenium rather than tellurium.

Figure 10.

(a) Cyclic voltammograms for [Li2S + 0.1 S/Se/Te] cathodes at scan rate range of 200 to 2000 μv s-1 (b) electrochemical performance of Li || [Li2S + 0.1 S/Se/Te] half cells. Reprinted from ref. [26].

2.4.1 Impact of Se and Te substituted polysulfides on lithium deposition

Electrochemical performance of Li2S and Li2S2 entirely depends on the lithium plating and stripping effectively. Figure 11a shows capacities for anode-free Ni || [Li2S + 0.1 S/Se/Te] full cells at ~ 1 mA cm−2 (C/5). Selenium showed rapid capacity fade and loss of their peak capacity 50% after the 35 cycles. By the introduction of tellurium exhibited remarkable cycling stability in the anode-free configuration and retain 52% of peak capacity at 265 cycles. The loss rate of lithium per cycle is decreased to 2.14% with Se and 0.24% with Te [64]. The improvement in lithium plating and stripping reversibility reflects by the coulombic efficiencies of the anode-free full cells as shown in Figure 11b. Polytellurosulfides showed a dramatic effect on lithium cycling efficiency by situ formation, kinetic hindrance occurred with tellurium substitution in polysulfides compared with selenium. The formation of polyselenosulfides has no effect on the reversibility of lithium deposition. These improvements were analyzed with symmetric Li || Li cells containing Li2SexSy and Li2TexSy introduced as an electrolyte components. Electrochemical impedance spectroscopy reveals the polyselenosulfides showed high and unstable overpotentials (~100 mV) and polytellurosulfides enable low and stable overpotentials (~10 mV), indicating a thin SEI layer has excellent ionic transport properties. The dense and uniform lithium deposits formed with polytellurosulfides exhibited irreversible loss of lithium. The electrochemical performance of an anode-free Ni || Li2S full cell with 0.05 Se + 0.05 Te additives was observed. There is a synergetic effect realized that the higher initial capacities and cycling stability than that of the pure 0.1 Te and 0.1 Se based cells. Therefore we believe that the presence of SeS2*radicals increased the faster capacity fade with 0.05 Se + 0.05 Te than with 0.1 Te, representing that an electrolyte system might be allowed radical anion of selenium to obtained higher capacities and also retained for longer cycles. These results further implemented impractical, large-area (4.8 × 8.1 cm2), single-layer pouch cells assembled in the anode-free configuration (N/P = 1) with a 164 mg Li2S cathode (4.2 mg cm−2) containing 10 wt% Te0 (Te:Li2S molar ratio = 0.055) and operating under lean-electrolyte conditions (E/Li2S = 4.5 ml mg−1) the results are shown in Figure 12c [65]. Tellurium was replaced with carbon black for control. Figure 11b showed the control cell exhibited a high initial capacity of 77 mAh, but it has very rapid capacity fade 80% retention within 13 cycles. By the addition of Te exceeded 80% of its peak capacity for nearly 150 cycles and retain their cycling capacity without rapid drop until the electrolyte dry-out nearly 300 cycles [66, 67]. The initial rise in capacity can be regarded as ‘activation period’ in which the dissolution of tellurium slowly into polysulfides. The improvement in cycle life with the introduction of tellurium can be attributed to the stabilizing effect of polytellurosulfides on lithium deposition. These results are valid practically relevant to the cell design and testing parameters such as long cycle life and high energy dense, anode-free configuration significantly closer to commercial viability of Li–S system.

Figure 11.

Electrochemical performances – (a) capacity retention and (b) coulombic efficiencies of anode-free Ni || [Li2S + 0.1 S/Se/Te] full cells. (c) Electrochemical performances of large-area (39 cm2) anode-free Ni || Li2S single-layer pouch full cells with 10 wt% tellurium (Te: Li2S molar ratio = 0.04) or 10 wt% carbon black as cathode additives. Reprinted from ref [26].

Figure 12.

(a) Crystal structure of Li2X (X = S, Se, and Te) and the three Li+ ion diffusion pathways marked as purple [100], red [110], and green [111] lines. Migration energy barriers along [110, 111] show a steady reduction from Li2S to Li2Se and to Li2Te. (b) Li+-ion transport pathway in Li2TeS3 along the x-axis and the corresponding energy barrier based on single-ion migration. Reprinted from ref [26].

2.4.2 Lithium-ion transport properties of selenides, tellurides, and thiotellurates

First principles calculations were evaluated to understand reduced species on lithium deposition and their ionic transport properties. Li2S, Li2Se, and Li2Te indicated crystallize in a cubic antifluorite structure with a face-centered cubic anionic framework along with Li+ ions in the tetrahedral sites. Figure 12a showed Li+ can diffuse along the directions of [100], [110], [111]. Climbing image nudge elastic-band (CI-NEB) method is used to find the diffusion barriers along each of these pathways. Barrier energy found to be ̴0.3 eV in the lowest- pathway [68]. Barrier energies are 0.875 eV to 0.748 eV to 0.539 eV, calculated from the transitions Li2S to Li2Se to Li2Te, respectively. Te2 exhibited more polarizable anionic framework compared to those of S2and Se2due to the larger size and lower charge density. Previous report [69] reveals that the larger size of Te2 provide more open channel along [110], [111] in which high diffusion pathways for Li+ ions. Therefore, Li2Te due to its more uniform, homogenous, and dense lithium deposition can provide alternate pathways to facilitate three-dimensional ion transport. Li2TeS3 exhibited monoclinic structure with trigonal pyramidal TeS32−anions arranged in layers. The Li+ ions coordinated with sulfur atoms and occupied their octahedral and tetrahedral sites. Li2TeS3 unit cell consists of eight distinct steps between five adjacent sides and the non-equivalent lithium sites can be found in the migration pathway. The single-ion NEB model was introduced to calculate the corresponding barrier energies and find the most favorable path. It indicates the migration from one tetrahedral site to another tetrahedral site through an intermediate octahedral site in the direction of the x-axis as shown in Figure 12 (b). The migration barriers are found to be 0.378 and 0.250 eV. The barrier energies are found to be 0.4 and 0.6 eV for other migration pathways. Therefore, multiple viable Li+ ion diffusion pathways available in Li2TeS3 in three dimensional paths ways for ion transport due to stable and reversible lithium deposition. All these factors would improve the lithium cyclic efficiencies by the formation of interfacial components with polytellurosulfides.

2.4.3 Interfacial chemistry of Se and Te substituted polysulfides on lithium batteries

In order to understand their effects of modified Li2Sn, Li2SexSy and Li2TexSy species by XPS, in which lithium anode-free full cells analyzed after 20 cycles. Figure 13a shows the S 2p + Se 3p and Li 1 s + Se 3d spectra for the cell with addition of 0.1 Se. The S 2p spectra SO42− species are dominated by oxidized sulfur due to the decomposition of LiTFSI. The dominated peaks at 165 eV and 58.7 eV are appeared in the Se 3p and Se 3d spectra, respectively by the oxidized Se+4 in selenites (SeO32−) and minor components are present due to Reduced sulfur species (Li2S). The presence of oxidized selenium species is due to LiNO3, which is a strong oxidizing agent and oxidizes selenides (Se2) into selenites (SeO32−) [70]. Therefore, polyselenosulfides introduced not only the fundamental alter to the lithium–electrolyte interface, which remains dominated by oxidized sulfur/selenium species. Figure 13b shows the cell with 0.1 Te additive spectra of S 2p and Te 3d. The reduced sulfur species (S2− at 160.6 eV) exhibited dominated S 2p spectra. Likewise, the Te 3d spectra are dominated by sulfurized tellurium species (Te+4 at 574.6 eV). The quantification of the spectra reveals the formation of thiotellurate (TeS32−) species [70]. Thus, the formation of Li2TeS3 as the dominant interfacial component and are reduced on the lithium surface. Some of previous research reveals that the oxidized sulfur species are present as minor components. It is extended that the oxidized tellurium species (TeO32−) made only a minor fraction of tellurium atoms on the lithium surface. Hence introduction of tellurium alters the lithium–electrolyte interface by the reduction of sulfur species (as Li2TeS3). These XPS observations consistent with the time-of-flight secondary ion mass spectrometry (ToF-SIMS). Figure 13c shows the profiles for Li2 (metallic lithium) and SO3 for three-dimensional reconstructions. A thick layer of electrolyte decomposition products is observed on the deposited lithium with polyselenosulfides but not with polytellurosulfides. Depth profiles indicate that the selenium has strong signal for SeO is compared to that for SeS. This is reversed for tellurium, in which TeS exhibited much stronger compared to that for TeO. Thus, majority of tellurium atoms made bond with sulfur and the majority selenium atoms making bond with oxygen. This is due to oxidation of LiNO3 explained by Pearson’s HSAB theory [71]. The soft Lewis acid cations (Te+4) are formed by the Tellurium that prefer soft Lewis bases such as S2 sulfides, while selenium forms hard Lewis acid cations (Se+4) that prefer hard Lewis bases such as O2oxides [72, 73]. The divergent lithium stabilization capabilities of polyselenosulfides and polytellurosulfides explained the differences in lithium interfacial chemistry. The sulfide anionic framework such as Li2TeS3 identified as preferable compared to an oxide anionic framework such as Li2SeO3 or Li2SO3 [74]. The greater size and polarizability of S2compared to those of O2, improves ionic transport properties by reducing Li+ ion diffusion barriers. The varying compositions of tellurium and selenium to get a stable sulfide-rich SEI layer, in the presence of LiNO3 changed the characteristics of lithium deposition [75].

Figure 13.

(a) S 2p + Se 3p and Li 1 s + Se 3d spectra for the lithium surface in an anode-free full cell cycled with polyselenosulfides. (b) S 2p and Te 3d spectra for the lithium surface in an anode-free full cell cycled with polytellurosulfides. (c) 3D reconstructions of ToF-SIMS depth profiles for Li2 (metallic lithium) and SO2 (oxidized sulfur species) secondary ions. Reprinted from ref. [26].

2.5 Freestanding Se1-xSx foamy cathodes for high-performance Li-Se1-xSx batteries

The development of supercritical CO2 synthesis of selenium-sulfur solid solutions (Se1-xSx) are promising new cathodic materials for high-performance secondary lithium batteries due to their high electric conductivity than S and superior theoretical specific capacity than Se. The morphology and microstructure of N-doped carbon framework with three-dimensional (3D) interconnected porous structure (NC@SWCNTs) host are characterized by SEM and TEM pictures as shown in Figure 14. A depicted in Figure 14a shows the NC@SWCNTs host 3D honeycombed structure and interconnected melamine foam framework. The magnification SEM images (Figure 14a,c) reveals that numerous interlaced SWCNTs are covered the surface of melamine foam by the derived carbon skeletons and SWCNTs are formed as thin sheets between carbon skeletons. This structure of NC@ SWCNTs exhibited a highly conductive 3D network to transport the electron or ion, but also increases the mechanical strength as well as flexibility of NC@SWCNTs host. TEM results (Figure 14d ) reveals that SWCNTs, 3D network structure are crisscrossed in carbon skeletons. By EDS analysis the main elements found to be in the NC@SWCNTs are C, O and N, which are uniformly distributed as shown in Figure 14e. N signal is derived from melamine foam because of it contain high N. 3D network structure of NC@SWCNTs is made of by composing SWCNTs-coated N-doped carbon skeleton melamine foam and wafery sheets of SWCNTs. The NC@SWCNTs consists of pores and layer gaps are favorable for loading of Se1-xSx active conductive materials. The 3D conductive network promotes not only redox kinetics, but also endow NC@ SWCNTs host with strong buffering in volume during cycling. Further, N doped is also beneficial for the adsorption of intermediates, after Se1-xSx impregnation, compared to NC@SWCNTs host. NC@SWCNTs@ Se1-xSx composites retain their original morphology of NC@SWCNTs and no discernible Se1-xSx particles can be found. According to EDS results, the C, N, Se and Se signals are overlapped well, suggesting Se1-xSx composites are uniformly permeated into the pores and layer gaps of NC@SWCNTs host with the assistance of SC-CO2 due to the good permeability, excellent diffusivity and high solubility of SC-CO2 [76].

Figure 14.

(a–c) Shows the SEM images, (d) represents the TEM image (e) indicated STEM image of NC@SWCNTs and the corresponding mapping images. Reprinted from ref. [76].

Electrochemical performance of NC@ SWCNTs@ Se1-xSx composites based cathodes evaluated in Li- Se1-xSx batteries using carbonate-based electrolyte (LiPF6-EC/DMC). Figure 15a show initial three cyclic voltammetry (CV) curves of NC@SWCNTs@ Se1-xSx cathodes with scanning rate of 0.1 mV s−1 in the potential range from 1.0 to 3.0 V versus Li/Li+. Initially, a sharp reduction peak at ∼1.38 V, a small reduction peak at ∼2.37 V and a broadened oxidation peak at ∼2.14 V appeared. The sharp reduction peak at ∼1.38 V shifts to ∼1.7 V and the small reduction peak at ∼2.37 V was disappeared after the first scan. Initially, during the lithiation the peak shift due the activation process and the polarization is also further reduced [77] CV curves overlapped after the first scan reveals that the good cyclability and reversibility of NC@SWCNTs@ Se0.2S0.8 cathode [78]. Notably, CV curves of NC@SWCNTs@ Se1-xSx cathodes obtained differently from S cathode, representing the change of electrochemical reaction of S by Se and it is more conducive and stable with carbonate-based electrolytes. Figure 15b shows the galvanostatic charge-discharge curves of NC@SWCNTs@ Se1-xSx cathodes are consistent with the result of CV. There are two plateaus are observed (1) extremely short plateau at ∼2.38 V, and (2) a long plateau at ∼1.75 V in the first cycle. Further, subsequent cycles, long plateau at ∼1.75 V becomes a little steeper and shifts to ∼1.88 V and the short plateau appeared at ∼2.38 V. The short plateau appeared at ∼2.38 V is attributed to the transformation of Se0.2S0.8 into polysulfides/polyselenides. The short plateau is disappeared due to dissolution of intermediates into the electrolyte [79]. The long plateau at 1.75–1.88 V is attributed to the conversion of polysulfides/polyselenides to Li2S/Li2Se [77]. There is only one sloping plateau appeared during the charge process, at ∼2.12 V due to the conversion of Li2Se/Li2S to Se0.2S0.8. The cyclic performance of NC@SWCNTs@ Se1-xSx cathodes at a current density of 0.2 A g−1 with different Se/S ratios as shown in Figure 15c. As prepared NC@SWCNTs@Se0.2S0.8 cathode delivers the highest initial discharge capacity (2398.5 mA h g−1) among all the samples. Discharge capacity exceeds the theoretical capacity at initial stage may be due to side reactions and the formation of SEI layer on the surface of electrode [80]. Electrochemical characteristics of NC@SWCNTs@ Se0.2S0.8 cathode exhibit the superior cyclic stability. Figure 15d showed the rate capabilities of NC@SWCNTs@ Se1-xSx cathodes at various current densities. Among all the samples, NC@SWCNTs@ Se0.2S0.8 cathode showed the best rate performance. At the various current densities of 0.2, 0.5, 0.8, 1.0 and 2.0 A g−1 the reversible rate capacities of NC@SWCNTs@ Se0.2S0.8 cathode are found to be 998.4, 723.7, 606.8, 506.1, and 415.0 mA h g−1, respectively. The reversible discharge capacity of NC@SWCNTs@ Se0.2S0.8 cathode reverts to the initial value, when the current density switches back to 0.5 A g−1. NC@SWCNTs@ Se0.2S0.8 cathode with Se loading of as high as 4.4 mg cm−2 exhibited areal capacity of as high as 2.78 mA h cm−2 is the best candidate most reported Se1-xSx cathodes in literature [77, 81, 82, 83]. The electrochemical performance of NC@SWCNTs@Se0.2S0.8 cathode is more effective due to the following reasons: 1) Se and S in Se0.2S0.8 solid solution play various roles: Se can improve more electrical conductivity, whereas the S can raise its capacity. 2) N-doped 3D porous carbon matrix and interlaced SWCNTs can provide storage and the structural stability; thereby promote the cycling stability of NC@SWCNTs@ Se1-xSx cathodes. NC@SWCNTs@ Se0.2S0.8 cathode exhibits good cycling stability (632 mA h g−1 at 0.2 A g−1 at 200 cycle) and high rate performance (415 mA h g−1 at 2 A g−1) due to well-designed structure as well as optimized chemical composition with in carbonate-based electrolyte. Hence these developments of high-performance Se1-xSx cathodes suitable for advanced Li- Se1-xSx batteries.

Figure 15.

(A) CV curves of NC@SWCNTs@Se0.2S0.8 cathode. (B) Charge/discharge curves of the NC@SWCNTs@ Se0.2S0.8 cathode at 0.2 a g−1. (C) Cycle performances and (D) rate performances of NC@SWCNTs@Se1-xSx cathodes. Reprinted from ref [76].

2.6 Nanoreactors for metal-chalcogen batteries

Porous hollow nanoreactors are investigated widely for lithium selenium and tellurium batteries. The mesoporous material exhibited considerable porosity (0.2 cm3 g−1) and a large surface area of 462 m2 g−1, which allowed for uniform distribution of Se8. The Se8/C based lithium selenium batteries showed a high reversible capacity of 480 mA h g−1 at 0.25C (1C = 678 mA g−1) without loss of its capacity after 1000 cycles [84]. Further development of the Se/porous carbon cathode battery showed a high volumetric capacity of 3150 mA h cm−3 with excellent rate capability about 1850 mA h cm−3 at 20C. Therefore, it will be used for future commercialization of LSeBs [85]. Single-atom Co decorated hollow porous carbon also works as a nanoreactor with superior catalytic activity to polyselenides. These (Se@CoSA-HC) cathodes based batteries exhibited high discharge capacity, superior cycling stability,an excellent rate capability [86]. Metal or heteroatom doping (N, S, and Co) is also another alternative approach to enhance the utilization of Se or Te [87, 88, 89]. He et al. synthesized a nanoporous Co and N-co-doped carbon nanoreactor (C–Co–N) provide a high Te loading (77.2 wt%) provide ultrahigh capacity of 2615.2 mA h cm−3 and superior rate performance of 894.8 mA h cm−3 at 20C as shown in Figure 16 [90]. Design structure and micro-environmental of Te-based nanoreactors could provide high electrochemical performance. In conclusion, the development of hollow porous nanoreactors not only provide a suitable specified space for chalcogens (S, Se, and Te), but also load active species for the regulation of the microenvironment in the electrode. Further development of nanoreactors, it is necessary to design the new methodologies at the molecular level to regulate the microenvironment of the catalyst.

Figure 16.

(a) CV curves of the three dimensional rGo/tellurium (3DGT) aerogel at a scan rate of 0.1 mV s−1. (b) Discharge curves of the 3DGT cathode at a 0.2 C. (c) Cyclic performance of the 3DGT cathode at 0.2 C for 200 cycles. (d) Rate performance at various crates for the 3DGT cathode. (e) Cyclic stability of 3DGT cathode at 1 C for 500 cycles. Reprinted from ref. [90].

2.7 Flexible cholcogen lithium-ion batteries

2.7.1 Flexible sulfur cathodes

Zhang and his coworkers developed 1D/3D hybrid flexible sulfur electrodes with good flexibility and exhibited improved electrochemical performance [91, 92]. They used sulfur-infiltrated 3D nanostructure porous carbon materials with various sizes nanometers to ten micrometers representing with high versatility and applicability for constructing flexible electrodes. These materials not sustain without support, therefore by incorporating ultralong CNT scaffolds, very robust films are obtained without sacrificing mechanical flexibility compared to ultralong CNT/MWCNT film. Such a materials exhibited tremendous specific surface area, high micro or mesoporosity, and surface functionalities than MWCNTs. Hence, this strategy is an ideal generic and versatile host to facilitate flexible sulfur cathodes. The use of graphene in 2D/3D hybridized structure is essential to alternative of CNTs in the 1D/3D, which provides required mechanical adhesion and good electrical conduction into 3D carbon constituents, but that lack of flexibility. Therefore, Wu et al. [93] demonstrated freestanding graphene based hierarchical porous carbon (GPC) films for flexible sulfur cathodes batteries as shown in Figure 17. Graphene-based microporous carbon (GMC) sheets are obtained by thin layers of microporous carbon were coated on both sides of GO after hydrothermal carbonization and KOH activation. The small sulfur molecules are stored in rich micropores of GMC, provides stronger physical confinement than normal graphene. Therefore, GPC files based batteries showed excellent cycling performance with stabilized capacities of 1030(422) and 626(357) mA h gsul(ele)−1 at 0.2C with the sulfur content as 41 and 57 wt%, respectively. In general, graphene-based film electrodes showed rapid decay in their capacity due to their polysulfide dissolution. Furthermore, the GPC-S cathode films used in flexible Li-S batteries by attaching the tape to pack the material, displaying comparable electrochemical performance in both flat and bent states as shown in Figure 18(a) and (b). Ni et al. [94] reported a facile route for synthesizing ultrathin and flexible composite films based on rGOwrapped sulfur particles with the help of sodium alginate (SA) aqueous binder, which worked as a surfactant and an adhesive agent. The SA-glued electrode battery exhibited a high reversible capacity of 1341(818) mA h gsul(ele)−1 at 0.1C and retained its capacity 823(502) mA h gsul(ele)−1 at 0.5C after 100 cycles, which are more better compared to physically mixed rGO/S film. Therefore, in order to improve electrical conductivity and their mechanical stiffness, researchers made hybride by mixing SA with polyaniline and used as glue for rGO/Mn3O4/S nanocomposite particles electrode films prepared. They exhibited a high capacity of 1015(538) mA h gsul(ele)−1 at 5.0 A gsul−1 (B3.0C) and capacity retention of 71% after 500 cycles [95].

Figure 17.

Preparation of free-standing graphene-based porous carbon (GPC) films 1) impregnation of sulfur into the micropores of 2D graphene-based microporous carbon (GMC) sheets; 2) non-covalent functionalization of carbon (GMC)-sulfur sheets by CTAB. 3) assembly of positively charged GMC-sulfur sheets and negatively charged graphene oxide.Reprinted from ref. [93].

Figure 18.

(a) The second charge-discharge profiles of the GPC-sulfur cathode films at the bent and flat states at 0.5C. (b) the cycle performance for the GPC-sulfur cathode films at 0.5C and 1C, and inset showing that a bent cell is encapsulated in the glass bottle filled with argon. Reprinted from ref. [93].

2.7.2 Flexible Li: Se batteries

Selenium has several major merits for serving as cathode materials over to sulfur: (1) The magnitude higher electrical conductivity is an approximately 1024 times higher (2) More stable at room temperature, chain-like allotrope h-Se is more electro-active and more easily stabilized via spatial confinement (3) Selenium has more compatibility with conventional, cheap carbonate-based LIB electrolytes [96]. Therefore, selenium exhibits a better utilization rate, cyclic stability, and rate capability than sulfur. The volumetric specific capacity of h-Se is 3265 mA h cm−3 comparable to sulfur, 3461 mA h cm−3, therefore it is more suitable for portable electronic devices and electrical vehicles due to its volume sensitive. Theoretically, the Li–Se battery utilizing h-Selenium as cathode lithium metal as anode, respectively at average voltage of 2.0 V, affords high gravimetric and volumetric energy densities of 1155 W h kg−1 and 2528 W h L−1, respectively. Han et al. [68] introduced the mesoporous carbon nanoparticles (MCNs) with smaller size of 50 nm and favorable mesopore dominance, efficiently eliminated agglomeration in the bulk selenium. The electroactive selenium chains were stabilized in smaller micropores or mesopores, enabling high utilization and good cycling stability according to previous reports of Se–micro−/mesoporous carbon composite cathodes [84, 97]. The flexible Se/MCN–rGO cathodes demonstrated an ultrahigh selenium utilization of 97% at 0.1C, i.e., 655(406) mA h gsel(ele)−1. They exhibited good long cycling life with 89% capacity retention after 1300 cycles at 1.0C. This work is one of the most remarkable achievements for flexible Li–Se batteries by considering the high content of selenium. Similarly, Yu and Zhu’s group prepared the composite PCNFs are represented as f-PCNFs, and they maintained good flexibility after selenization as shown in Figure 19ad [98]. Very less crystalline selenium was present in PCNFs than in f-PCNFs leads to a remarkable improved capacity and initial Coulombic efficiency as shown in Figure 19e. The capacity and initial Coulombic efficiencies are 643/322 mA h gsel/ele−1 and 56.9% for Se@PCNFs at 0.05 A gsel−1, while 405/203 mA h gsel/ele−1 and 34.9% for Se@f-PCNFs. This is attributed due to the suppression of side reactions between free polyselenides produced from bulk selenium and carbonate electrolytes. Additionally, owing excellent encapsulation of selenium in the 1D conductive porous skeleton, flexible Se@PCNF cathode also exhibited non fading cycling performance with a capacity of 516(270) mA h gsel(ele)−1 retained after 900 cycles at 1.0 A gsel−1 (B1.5C). The same electrospun PCNF–CNT also demonstrated in flexible Li–S batteries as like flexible selenium PCNF–CNT fabricated, battery exhibited reversible capacity of 638(223) mA h gsel(ele)−1 after 80 cycles at 0.05 A gsel−1 (B0.074C) [99]. The utilization of conductive selenium (94%) was much higher than that of sulfur (38%) for the same PCNF–CNT conductive backbone demonstrated in cathodes.

Figure 19.

Schematic representation of flexible selenium cathodes. (a)&(b) represents the synthesis of selenium (Se)@PCNF electrodes. (c) and d) picture of flexible Se@PCNF electrode. Cyclic performance of flexible Se@PCNF and Se@f-PCNF cathodes in (e) Li–Se and (f) Na–Se batteries at 0.05 a gsel−1. [reprinted from ref. [98].

2.7.3 Flexible Li–Te batteries

Tellurium is the last nonradioactive element in the chalcogen family, which exhibit highest electrical conductivity of 2.5 Scm−1 compared to all nonmetallic materials. Te shows low gravimetric specific capacity of 420 mA h g−1 due to its heavy atomic weight, it exhibited comparable volumetric specific capacity of 2621 mA h cm−3 to that of sulfur or selenium. Te is an electrically conducting active material required less carbon to prepare electrode. The decreasing weight of the carbon favors both volumetric and gravimetric specific capacities. A Li–Te battery exhibited theoretical gravimetric and volumetric energy densities of 682 W h kg−1 and 2078 W h L−1, respectively with an estimated output voltage of 1.8 V. Li–Te battery was first demonstrated by Wang’s group [100], in which tellurium/porous carbon composite cathode and a carbonate electrolyte as the components of battery. The Li–Te battery showed an average voltage of 1.5 V and a reversible capacity of 224 mA h gtel−1 at 0.05 A gtel−1. They observed that 87% retention after 1000 cycles. Considering the relatively low voltage and promising volumetric capacity, Guo and coworkers demonstrated tellurium/carbon composites as anode materials for LIBs, indicating that extremely high tellurium utilization of 98% and a long-term cycling stability [101]. Particularly, Te is quite interest for flexible electrode materials due to its two most favorable features: (1) high electrical conductivity compared to carbon and (2) The formation of 1D Te nanostructures along the c-axis, i.e., [001] due to its inherent chirality of helical Te chains in the h-Te crystal [102]. Hence, freestanding films consisting of ultralong Te NWs used directly used as an electrodes. Freestanding Te mat via vacuum filtration of Te NWs with a diameter of 7 nm grown in the [001] direction developed by Ding et al. as shown in Figure 20a-i–iii [104]. Such a high anisotropic 1D Te nanostructure exhibited fully Te zig-zag chains to lithium ions transport and a showed high electrical conductivity of 6.7 Scm−1 in the direction perpendicular to the c-axis as shown in Figure 20a(iv–vi). The flexible tellurium cathode comprises the Te NWs along with the new electrolyte exhibited a desirable capacity of B144 mA h gtel/ele-1 at 0.1 A gtel−1 (0.24C). The volumetric energy density of 1800 W h L−1 observed after 80 cycles as shown in the Figure 20a-vii. Further, He and Chen’s et al. [103] demonstrated a flexible tellurium cathode prepared from a 3D hierarchical aerogel with Te NWs wrapped homogeneously by rGO as shown Figure 20b. The synthetic method was adopted from the previous report on 3DCG–Li2S. The rGO/Te NW electrode made of 63 wt% tellurium delivers high capacities of 418(263) and 174(110) mAh gtel(ele)−1 at 0.2 and 10C, respectively and excellent long-cyclic performance at a high rate of 1.0C. Therefore, as prepared flexible Te-NWs electrodes are quite attractive over the sulfur and selenium counterparts due to their distinguish features.

Figure 20.

Flexible tellurium NWs cathodes. (a) Flexible, carbon-free TeNW mat: (i) photograph; (ii) SEM and (iii) TEM morphology images; (iv, v) nanoscale crystalline structure of freestanding TeNW mat; (vi) simulated crystal structure of h-Te; (vii) cyclic performance of TeNW mat at 0.1 a gtel−1. Reprited from ref. [84] (b) 3D rGO/TeNW aerogel: (i) fabrication of 3D rGO/TeNW aerogel and its derived flexible electrodes; photographs of (ii) 3D rGO/TeNW aerogel and (iii) flexible 3D rGO/TeNW electrode. Reprinted from ref. [27, 103].


3. Conclusions

Al-S batteries prepared based on S/HPCK cathode delivered a capacity of 1027 mA h g−1 at 0.2 A g−1 for 50 cycles and exhibited excellent cyclic ability 405 mA h g−1 at 1 A g−1 for 700 cycles due to large porous structure with high surface area by adding of carbon powder. The Al-Se battery based on MCFs material with pore size of 7.1 nm exhibited a good capacity of 366 mA h g−1. The chloroaluminate ion diffusivity greatly improved, which enhances the electrochemical behavior of Al-Se batteries. The rGO materials are introduced in the Al-Te batteries exhibited excellent cyclic ability and its initial capacity of 935.5 mA h g−1 and showed 467.5 mA h g−1 after 150 cycles with the Te loading of 70 wt% due to their excellent encapsulation.

The Interfacial layer of S@ spinel Ni-Co oxide double-shelled microspheres (NCO-HS) prepared under high sulfur loading exhibited minimum capacity fading rate of 0.045% per cycle over 800 cycles with high areal capacity of 6.3 mAh cm−2 and superior rate capability up to 5 C. As prepared capping layer of MTO-CNTs on surface of the sulfur cathode exhibited reversible capacity of 1212 mAh g−1 at the 0.2 C. The cycling stability of the device with the capping layer at 0.5 C rate still it retains a capacity of 577 mAh g−1 after 500 cycles with capacity decay rate of 0.07% per cycle, indicating a good cycling stability.

There is a gradual deterioration in all the compounds in the case of mixed cationic and anionic cyclic performance at a very low current density of 10 mA/g for 50 continuous cycles in terms of capacity fade. Cycling test of the 0.2 Fe substituted cathode done at 50 mA/g, reveals that the high rate cycling stability about 76% retention after 80 cycles. In the case of polysulfides on lithium deposition, the addition of Te exceeds 80% of its peak capacity for nearly 150 cycles and retains their cycling capacity without rapid drop until the electrolyte dryout nearly 300 cycles.

The developments of high-performance Se1-xSx cathodes such as NC@SWCNTs@ Se0.2S0.8 cathode exhibits good cycling stability (632 mA h g−1 at 0.2 A g−1 at 200 cycle) and high rate performance (415 mA h g−1 at 2 A g−1) due to well-designed structure as well as optimized chemical composition with in carbonate-based electrolyte. Synthesized a nanoporous Co and N-co-doped carbon nanoreactor (C–Co–N) provide a high Te loading (77.2 wt%) provide ultrahigh capacity of 2615.2 mA h cm−3 and superior rate performance of 894.8 mA h cm−3 at 20C. Design structure and micro-environmental of Te-based nanoreactors could provide high electrochemical performance.

In the case of flexible batteries, in order to improve electrical conductivity and their mechanical stiffness, researchers made hybride by mixing SA with polyaniline and used as glue for rGO/Mn3O4/S nanocomposite particles electrode films, exhibited a high capacity of 1015(538) mA h gsul(ele)−1 at 5.0 A gsul−1 (B3.0C) and capacity retention of 71% after 500 cycles. The electrospun PCNF–CNT also demonstrated in flexible Li–S batteries as like flexible selenium PCNF–CNT fabricated, battery exhibited reversible capacity of 638(223) mA h gsel(ele)−1 after 80 cycles at 0.05 A gsel−1 (B0.074C). The rGO/Te NW electrode made of 63 wt% tellurium delivers high capacities of 418(263) and 174(110) mAh gtel(ele)−1 at 0.2 and 10C, respectively and excellent long-cyclic performance at a high rate of 1.0C. Therefore, as prepared flexible Te-NWs electrodes are quite attractive over the sulfur and selenium counterparts due to their distinguish features.


  1. 1. Hao J, Xiong K, Zhou J, Rao AM, Wang X, Liu H, et al. Energy & Environmental Materials. 2021;5(1):261-269
  2. 2. Wang X-X, Denga N-P, Wei L-Y, Yang Q , Xiang H-Y, Wang M, et al. Chemistry, an Asian Journal. 2021;16(19):2852-2870
  3. 3. Zhang Q , Cheng X, Wang C, Rao AM, Lu B. Energy & Environmental Science. 2021;14:965
  4. 4. Liu Q , Rao AM, Han X, Lu B. Advancement of Science. 2021;8:2003639
  5. 5. He S, Zhang D, Zhang X, Liu S-Q , Chu W-Q , Yu H-J. Advanced Energy Materials. 2021;11(29):2100769
  6. 6. Zhang X, Jiao S, Tu J, Song W-L, Xiao X, Li S, et al. Energy & Environmental Science. 1918;2019:12
  7. 7. Zhang Y, Liu S, Ji Y, Ma J, Yu H. Advanced Materials. 2018;30:1706310
  8. 8. Boyjoo Y, Shi H-D, Tian Q , Liu S-M, Liang J, Wu Z-S, et al. Energy & Environmental Science. 2021;14:540575
  9. 9. Huang C, Sun T, Shu H, Chen M, Liang Q , Zhou Y, et al. Electrochimica Acta. 2020;334:135658
  10. 10. Saha S, Assat G, Sougrati MT, Foix D, Li H, Vergnet J, et al. Nature Energy. 2019;4(11):977-987
  11. 11. Hansen CJ, Zak JJ, Martinolich AJ, Ko JS, Bashian NH, Kaboudvand F, et al. Journal of the American Chemical Society. 2020;142(14):6737-6749
  12. 12. Flamary-Mespoulie F, Boulineau A, Martinez H, Suchomel MR, Delmas C, Pecquenard B, et al. Energy Storage Materials. 2020;26:213-222
  13. 13. Nagarajan S, Hwang S-Y, Balasubramanian M, Thangavel NK, Reddy Arava LM. Journal of the American Chemical Society. 2021;143:15732-15744
  14. 14. Li Z, Liu J, Huo X, Li J, Kang F, Appl ACS. Materials Interfaces. 2019;11:45709
  15. 15. Lin S, Chen Y, Wang Y, Cai Z, Xiao J, Muhmood T, et al. Materials & Interfaces. 2021;13:9955-9964
  16. 16. Zhang W, Li S, Wang L, Wang X, Xie J. Sustain. Energy & Fuels. 2020;4:3588-3596
  17. 17. Xu Q-T, Xue H-G, Guo S-P. Inorganic Chemistry Frontiers. 2019;6:1326-1340
  18. 18. Du H, Feng S, Luo W, Zhou L, Mai L. Journal of Materials Science and Technology. 2020;55:1-15
  19. 19. Sun J, Du Z, Liu Y, Ai W, Wang K, Wang T. Advanced Materials. 2021;33:2003845
  20. 20. Jiao H, Tian D, Li S, Fu C, Jiao S, Appl ACS. Energy & Materials. 2018;1:4924
  21. 21. Faegh E, Ng B, Hayman D, Mustain WE. Nature Energy. 2021;6:21
  22. 22. Dong S, Yu D, Yang J, Jiang L, Wang J-W, Cheng L-W, et al. Advanced Materials. 2020;32(23):1908027
  23. 23. Gupta A, Bhargav A, Manthiram A. Advanced Energy Materials. 2019;9:1803096
  24. 24. Shin H, Baek M, Gupta A, Char K, Manthiram A, Choi JW. Advanced Energy Materials. 2020;10:2001456
  25. 25. Cheng X-B, Huang J-Q , Zhang Q. Journal of the Electrochemical Society. 2018;165:A6058-A6072
  26. 26. Nanda S, Bhargav A, Jiang Z, Zhao X-H, Liu Y-Y, Manthiram A. Energy & Environmental Science. 2021;14:5423-5432
  27. 27. Peng H-J, Huang J-Q , Zhang Q. Chemical Society Reviews. 2017;46:5237-5288
  28. 28. Chung SH, Manthiram A. Advanced Materials. 2019;31:1901125
  29. 29. Ding H, Zhou J, Rao AM, Lu B. National Science Review. 2021;8(9):nwaa276. DOI: 10.1093/nsr/nwaa276
  30. 30. Guo Y, Jin H, Qi Z, Hu Z, Ji H, Wan L-J. Advanced Functional Materials. 2019;29:1807676
  31. 31. Zhang X, Wang B, He S, Liu S, Tang M, Yu H. Journal of Materials Chemistry A. 2021;9:8966
  32. 32. Kong Y, Nanjundan AK, Liu Y, Song H, Huang X, Yu C. Small. 2019;15:1904310
  33. 33. Zhang X, Wang M, Tu J, Jiao S. Journal of Energy Chemistry. 2021;57:378
  34. 34. Li F, Liu Q , Hu J, Feng Y, He P, Ma J. Nanoscale. 2019;11:15418-15439
  35. 35. Liu T, Hu H, Ding X, Yuan H, Jin C, Nai J, et al. Energy Storage Materials. 2020;30:346-366
  36. 36. Park JS, Kim JK, Hong JH, Cho JS, Park SK, Kang YC. Nanoscale. 2019;11:19012-19057
  37. 37. Cao B, Huang J, Mo Y, Xu C, Chen Y, Fang H, et al. Materials & Interfaces. 2019;11:14035-14043
  38. 38. Liu Y, Kou W, Li X, Huang C, Shui R, He G. Small. 2019;15:1902431
  39. 39. Zhao C, Ge F, Shi L, Wang H, Liu J, Zhang J, et al. International Journal of Hydrogen Energy. 2019;44:30478-30485
  40. 40. Li F, Tao J, Zou Z, Li C, Hou Z, Zhao J. ChemSusChem. 2020;13:2761-2768
  41. 41. Li N, Chen Z, Chen F, Hu G, Wang S, Sun Z, et al. Carbon. 2019;143:523-530
  42. 42. Li W, Chen Z, Wang D, Gong Z, Mao C, Liu J, et al. Journal of Power Sources. 2019;435:226778
  43. 43. Mussa Y, Arsalan M, Alsharaeh E. Energy & Fuels. 2021;35:8365-8377
  44. 44. Ma L, Zhu G, Zhang W, Zhao P, Hu Y, Wang Y, et al. Nano Research. 2018;11:6436-6446
  45. 45. Du P, Wei W, Dong Y, Liu D, Wang Q , Peng Y, et al. Nanoscale. 2019;11:10097-10105
  46. 46. Huang Y, Field R, Chen Q , Peng Y, Walczak MS, Zhao H, et al. Communications & Chemistry. 2019;2:138
  47. 47. Xiang Y, Wang Z, Qiu W, Guo Z, Liu D, Qu D, et al. Journal of Membrane Science. 2018;563:380-387
  48. 48. Shi H, Sun Z, Lv W, Xiao S, Yang H, Shi Y, et al. Journal of Energy Chemistry. 2020;45:135-141
  49. 49. Wang N, Chen B, Qin K, Liu E, Shi C, He C, et al. Nano Energy. 2019;60:332-339
  50. 50. Li Y, Wang W, Liu X, Mao E, Wang M, Li G, et al. Energy Storage Materials. 2019;23:261-268
  51. 51. Wang J, Yang G, Chen J, Liu Y, Wang Y, Lao CY, et al. Advanced Energy Materials. 2019;9:1902001
  52. 52. Yan X, Zhang H, Huang M, Qu M, Wei Z. ChemSusChem. 2019;12:2263-2270
  53. 53. Hu Y, Zhu X, Wang L. ChemSusChem. 2020;13:1366-1378
  54. 54. Gao TJ, Xu DP, Yu ZH, Huang ZH, Cheng J, Yang Y. Journal of Alloys and Compounds. 2021;865:11
  55. 55. Chen YT, Abbas SA, Kaisar N, Wu SH, Chen HA, Boopathi KM, et al. Materials & Interfaces. 2019;11:2060-2070
  56. 56. Martinolich AJ, Zak JJ, Agyeman-Budu DN, Kim SS, Bashian NH, Irshad A, et al. Chemistry of Materials. 2021;33(1):378-391
  57. 57. Ramakrishnan S, Park B, Wu J, Yang W, McCloskey BD. Journal of the American Chemical Society. 2020;142(18):8522-8531
  58. 58. Zhu Z, Yu D, Yang Y, Su C, Huang Y, Dong Y, et al. Nature Energy. 2019;4(12):1049-1058
  59. 59. Yan P, Zheng J, Tang Z-K, Devaraj A, Chen G, Amine K, et al. Nature Nanotechnology. 2019;14(6):602-608
  60. 60. Augustyn V, Come J, Lowe MA, Kim JW, Taberna PL, Tolbert SH, et al. Nature Materials. 2013;12:518-522
  61. 61. Wujcik KH, Wang DR, Raghunathan A, Drake M, Pascal TA, Prendergast D, et al. Journal of Physical Chemistry C. 2016;120:18403-18410
  62. 62. Zhang G, Peng H, Zhao C, Chen X, Zhao L, Li P, et al. Angewandte Chemie, International Edition. 2018;57:16732-16736
  63. 63. Baek M, Shin H, Char K, Choi JW. Advanced Materials. 2020;32:2005022
  64. 64. Nanda S, Manthiram A. Energy & Environmental Science. 2020;13:2501-2514
  65. 65. Bhargav A, He J, Gupta A, Manthiram A. Joule. 2020;4:285-291
  66. 66. Cheng X-B, Yan C, Huang J-Q , Li P, Zhu L, Zhao L, et al. Energy Storage Mater. 2017;6:18-25
  67. 67. Chen S, Niu C, Lee H, Li Q , Yu L, Xu W, et al. Joule. 2019;3:1094-1105
  68. 68. Han K, Liu Z, Ye HQ , Dai F. Journal of Power Sources. 2014;263:85-89
  69. 69. Gupta MK, Singh B, Goel P, Mittal R, Rols S, Chaplot SL. Physical Review B: Condensed Matter and Materials Physics. 2019;99:224304
  70. 70. Zhang L, Ling M, Feng J, Mai L, Liu G, Guo J. Energy Storage Materials. 2018;11:24-29
  71. 71. Nanda S, Bhargav A, Manthiram A. Joule. 2020;4:1121-1135
  72. 72. Pearson RG. Journal of Chemical Education. 1968;45:581-587
  73. 73. Sahu G, Lin Z, Li J, Liu Z, Dudney N, Liang C. Energy & Environmental Science. 2014;7:1053-1058
  74. 74. Wang Y, Lu X, Zheng C, Liu X, Chen Z, Yang W, et al. Angewandte Chemie. 2019;131:7755-7759
  75. 75. Lau J, DeBlock RH, Butts DM, Ashby DS, Choi CS, Dunn BS. Advanced Energy Materials. 2018;8:1800933
  76. 76. Lu C-W, Fang R-Y, Wang K, Xiao Z, Kumar GG, Gan Y-P, et al. Frontiers in Chemistry. 2021;9:738977
  77. 77. Zhu T, Pang Y, Wang Y, Wang C, Xia Y. Electrochimica Acta. 2018;281:789-795
  78. 78. Guo B, Yang T, Du W, Ma Q , Zhang L-Z, Bao S-J. Journal of Materials Chemistry A. 2019;7:12276-12282
  79. 79. Li X, Liang J, Zhang K, Hou Z, Zhang W, Zhu Y. Energy & Environmental Science. 2015;8:3181-3186
  80. 80. Luo C, Zhu Y, Wen Y, Wang J, Wang C. Advanced Functional Materials. 2014;2014(24):4082-4089
  81. 81. Li Z, Zhang J, Lu Y, Lou XW. Science Advances. 2018;4:1687
  82. 82. Li Z, Zhang J, Wu HB, Lou XWD. Advances Energy Materials. 2017;7:1700281
  83. 83. Yao Y, Zeng L, Hu S, Jiang Y, Yuan B, Yu Y. Small. 2017;13:1603513
  84. 84. Luo C, Xu Y, Zhu Y, Liu Y, Zheng S, Liu Y, et al. ACS Nano. 2013;7:8003-8010
  85. 85. Li Z, Yuan L, Yi Z, Liu Y, Huang Y. Nano Energy. 2014;9:229-236
  86. 86. Tian H, Tian H, Wang S, Chen S, Zhang F, Song L, et al. Nature Communications. 2020;11:5025
  87. 87. Li Y, Zhang Y, Xu Q , Hu L, Shen B, Liu H, et al. ChemSusChem. 2019;12:1196-1202
  88. 88. Xu J, Ma J, Fan Q , Guo S, Dou S. Advanced Materials. 2017;29:1606454
  89. 89. Zeng L-C, Li W-H, Jiang Y, Yu Y. Rare Metals. 2017;36:339-364
  90. 90. He J, Lv W, Chen Y, Wen K, Xu C, Zhang W, et al. ACS Nano. 2017;11:8144-8152
  91. 91. Zhai PY, Huang JQ , Zhu L, Shi JL, Zhu W, Zhang Q. Carbon. 2017;111:493-501
  92. 92. Shi JL, Peng HJ, Zhu L, Zhu WC, Zhang Q. Carbon. 2015;92:96-105
  93. 93. Wu C, Fu LJ, Maier J, Yu Y. Journal of Materials Chemistry A. 2015;3:9438-9445
  94. 94. Ni W, Cheng JL, Li XD, Guan Q , Qu GX, Wang ZY, et al. RSC Advances. 2016;6:9320-9327
  95. 95. Ghosh A, Manjunatha R, Kumar R, Mitra S, Appl ACS. Materials & Interfaces. 2016;8:33775-33785
  96. 96. Yang CP, Yin YX, Guo YG. Journal of Physical Chemistry Letters. 2015;6:256-266
  97. 97. Yang CP, Xin S, Yin YX, Ye H, Zhang J, Guo YG. Angewandte Chemie, International Edition. 2013;52:8363-8367
  98. 98. Zeng LC, Zeng WC, Jiang Y, Wei X, Li WH, Yang CL, et al. Advanced Energy Materials. 2015;5:1401377
  99. 99. Zeng LC, Wei X, Wang JQ , Jiang Y, Li WH, Yu Y. Journal of Power Sources. 2015;281:461-469
  100. 100. Liu Y, Wang JW, Xu YH, Zhu YJ, Bigio D, Wang CS. Journal of Materials Chemistry A. 2014;2:12201-12207
  101. 101. Zhang J, Yin YX, You Y, Yan Y, Guo YG. Energy Technology. 2014;2:757-762
  102. 102. He Z, Yang Y, Liu JW, Yu SH. Chemical Society Reviews. 2017;46:2732-2753
  103. 103. He JR, Chen YF, Lv WQ , Wen KC, Wang ZG, Zhang WL, et al. ACS Nano. 2016;10:8837-8842
  104. 104. Ding N, Chen SF, Geng DS, Chien SW, An T, Hor TSA, et al. Advanced Energy Materials. 2015;5:1401999

Written By

Varishetty Madhu Mohan, Madhavi Jonnalagadda and VishnuBhotla Prasad

Submitted: 18 December 2021 Reviewed: 03 February 2022 Published: 09 May 2022