The relationship between morphologies and the electrochemical performances of MoS2‐based nanocomposites in the three typical and promising rechargeable systems is also discussed. Finally, perspectives on major challenges and opportunities faced by MoS2‐based materials to address the practical problems of MoS2‐based materials are presented.Typical layered transition‐metal chalcogenide materials, in particular layered molybdenum disulfide (MoS Figure2 materials in recent years. This shows visually that the MoS2 material is attracting increasing interest in the recent four years, based on the quickly increasing numbers of scientific publications. Until now, wide ranging categories of applications based on MoS2 nanomaterials have been investigated in energy storage devices, electronic sensors, and biomedical engineering instruments.2 layer consists of two S layers and one Mo layer (S‐Mo‐S), which is bonded to adjacent MoS2 layers by weak van der Waals force and the corresponding molecular structure scheme visualized by JSmol is shown in Figure 2 makes it an ideal solid lubrication material and it is possible to be used in many specific harsh environments such as space crafts and artificial satellites.Among these transition metals, molybdenum disulfide  have been found, and the typical structures of MoS2 are presented in Figure 2 are existent in nature and both of them have trigonal prismatic coordination . However, the 3R type MoS2, showing three layers per unit in a rhombohedral symmetry, is easily transformed to the more stable 2H type during heating. However, 1T‐MoS2 is a type of artificially prepared structure. Compared to 1T‐MoS2 .4)6Mo7O24·4H2O and thiourea have been utilized as the precursors), showing relative lower crystal quality yet abundant active sites and easy to form morphologies with various substrates . The preparation methods and typical morphologies of MoS2 are summarized in FigureUntil now, the study of the pure MoS2‐based materials and describe their applications in three types of energy storage devices  in detail. We also discuss the relationships between the tuned features and the electrochemical performances of MoS2 materials. Finally, we give a perspective on major challenges and opportunities faced by MoS2‐based materials to address the practical problems of MoS2‐based materials.In this article, we summarize new preparation methods forf MoS22.1−1 in theory) and capacity (755 mAh g−1), lower volume ratio and higher stability , the Li ion battery is regarded as the most promising energy storage system to power millions of portable devices and electric vehicles.With the rapid developments of mobile electric devices and electric vehicles, the demand for high energy batteries with highly efficient energy storage is more urgent. Relying on its high energy density value . The intercalation of MoS2 with lithium ions occurs between 3 V and 1.1 V during the discharge process, which shown in the redox reaction equationThe properties and electrochemical performances of MoS2 electrode is cycled between 3 V and 1.1 V with an intercalation reaction mechanism, contributing a capacity of 167 mAh g−1. Further discharge process lead to the decomposition of LixMoS2 and the formation lithium sulfide as the final product, which might follow the conversion reactions: The lithium ion insertion is fully reversible if the MoSxMoS2 after 1.1 V, contributing a capacity of 669 mAh g−1, there is no solid evidence to support Equation 2 at high rates due to the formation of metallic molybdenum. It is believed that the additional capacity comes from the nanostructured carbon materials of MoS2 composites, and cycling performance of MoS2 can also be enhanced by minimizing particle size and assembling with various carbon‐based materials. This discharge product of Li2S might cause other redox reactions to generate S as the charge product during recharge, which makes it face the shuttle effect of lithium sulfur batteries and arouse irreversible capacity loss.2 composite materials can be divided to three categories based on the dimension of the matrix: low dimensional , two dimensional (monolayered MoS2 nanosheets), and three dimensional .Even though, many reported results support the conversion reactions of Li2.1.12 layer can be arranged to zero dimensional spheres to decrease the mechanical strain generated during the conversion reaction with lithium ions and to convert the inner strain due to robust carbon spheres, which is beneficial to the long‐term cycling performance. The 0D MoS2/C spheres have been demonstrated as good candidates for LIBs that show quite impressive performances.2/carbon spheres (MoS2/C) using a simple hydrothermal method with the hollow carbon spheres as the matrix and SiO2 nanospheres as the sacrificed template. The thin MoS2 nanolayers with a layer distance of 0.65 nm are homogenously wrapped on the surface of hollow carbon nanospheres and show a monodispersed state, which can be seen in Figure2/C electrode retained a high reversible capacity of 750 mAh g−1 after 100 cycles, which was much higher than that of the bare MoS2 electrode (around 140 mAh g−1) as shown in Figure 2 and enhance the surface conductivity and charge transfer.−1 after 50 cycles).2 nanosheets vertically grown on the surface of carbon nanospheres, showing the good mechanical properties of the materials. The reversible capacity of the MoS2 nanocomposites maintained a capacity of 650 mAh g−1 after 300 cycles, which is attributed to the exceptional robust structural stability to buffer the large volume changes during cycles and the reduction of the diffusion energy barrier of Li+ in the lithiation/delithiation processes.With the assistant of carbon spheres, the MoS2 nanosheets is another novel way to prepare a 0D sphere, and plays an important role in providing a special dual effect to prevent the aggregation of MoS2 nanosheets and absorb the discharge byproduct Li2S. For instance, Zheng and Qu et al. loaded MoS2 nanolayers on the monodisperse sulfonated polystyrene (SPS) microspheres (served as both the sacrificed soft template and the carbon source) and in situ converted SPS to amorphous carbon after annealing at 800 °C, finally achieving the monodisperse C@MoS2 microspheres.2 microspheres with a diameter of 540 nm exhibited uniform ultrathin nanosheets and an interlayer spacing of 0.63 nm, slightly larger than the theoretical value of 0.62 nm as shown in Figure 2, C@MoS2 microspheres maintained a capacity of 652 mAh g−1 after 100 cycles in Figure 2 microspheres. The good cycling performance of C@MoS2 microspheres is associated with the interstices between ultrathin MoS2 nanosheets and the quasi‐hollow interior on buffering volume changes, and the porous carbon cores on absorbing the discharge byproduct of lithium sulfide. The well‐preserved morphologies of C@MoS2 microspheres after 100 cycles can fully validate the hypothesis, which also provides a new idea for MoS2‐based materials to achieve good electrochemical performance of LIB. Additionally, for the first tmie, Lou and co‐workers synthesized highly uniform Mo‐glycerate solid spheres via a solvothermal process and the precursor can be easily in situ converted to hierarachical MoS2 hollow spheres (HNS) by the subsequent sulfidation reaction, as shown in Figure 2 HNS exhibited a high capacity of 1100 mAh g−1 at 0.5 A g−1, and long discharge platform in Figure 2 solid nanospheres. That is ascribed to the unique hierarchical hollow structure, high surface area, and uniformly distributed structural strains.Soft polymer as a template for MoS2 sphere also varied with the calcination temperature and the layer distance. According to K. S. Kumar's study, it was found that the MoS2 nanoparticle, prepared using supercritical ethanol with a short reaction time, was more likely to perform as a nanoplate structure and it was quite sensitive to the calcination temperature.2 nanoparticles reached to 1342 mAh g−1 but the electrodes prepared at high temperature (800 °C) exhibited better cycling performance, as shown in Figure Moreover, the electrochemical performance of MoS2 nanotube‐like materials, including carbon nanotubes, carbon nanofibers, and metal oxide nanotubes. The MoS2 nanosheets can be uniformly wrapped on the surface of CNTs and exhibit enhanced electrochemical performance. For example, Lee and co‐workers used a unique microwave irradiation technique to prepare the cylindrical nanostructure of the MoS2‐CNT composite using sulfur and molybdenum chloride as raw materials, and solved the problem of non‐intimate contact between active MoS2 and carbonaceous materials, which can be seen in Figure2‐CNT composite is confirmed using high‐resolution transmission electron microscopy (HRTEM), as shown in Figure −1 to 984 mAh g−1 at 200 mA g−1 after 10 cycles and the capacity is maintained at 670 mAh g−1 at 1600 mA g−1. The good electrochemical performances of MoS2‐CNT composite are ascribed to the synergistic effects between the cylindrical nanostructured MoS2 materials and CNT, and the prevention of aggregation by the unique structure and the high‐conductivity of CNT to improve the conductivity of insulating MoS2 materials for LIBs.2 and CNT (3:1 to 1:2) were discussed andenhanced electrochemical performances are demonstrated.2/CNT hybrid for LIBs. The few layer MoS2 nanosheets were homogenously and vertically grown on the surface of CNTs, thus forming hierarchical nanostructures. The specific capacity of the hybrid reached 1293 mAh g−1 when the current density was 200 mA g−1.Compared to the nanospheres, the 1D materials (nanotubes/nanofibers) have more potential to be applied as electrodes for LIBs because of their high flexibility, one dimensional electron‐transferring channel, and the multiple/interconnected mechanical properties. There are three main matrixes to prepared 1D MoS2 ultrathin nanosheets is recently reported by Lou et al.2 nanohybrid with the tubular structure was prepared using an electrospinning method with the carboxylic group functionalized CNT as the precursor and polyacrylonitrile (PAN) as the sacrificial template after coating an ultrathin MoS2 nanosheets on the outlayer; this can be seen in Figure 2 nanosheets on the surface (interlayer distance of 0.65 nm) and the typical tubular structure with CNT as the core and PAN as the shell. The CNT/MoS2 nanohybrid electrode delivered a high capacity of ≈1100 mAh g−1 at 0.5 A g−1 with no obvious capacity decay in 200 cycles, and exhibited an exceptional rate capability up to 10 A g−1. Moreover, the CNT/MoS2 nanohybrid electrode retained a remarkably reversible capacity of ≈800 mAh g−1 at a current density of 5 A g−1 after 1000 cycles as shown in Figure 2 nanohybrid were associated with the well‐preserved 1D tubular structure and the fast charge transfer of lithium ions via CNT and tubular MoS2, as shown in Figure Another innovation using in situ grown CNTs in hollow carbon nanofibers supporting MoS2 nanocrystals and keep the nanofibers as thin as 100–150 nm in diameter, which can fully encapsulate different active materials for LIBs.2‐carbon nanofiber composites as shown in Figure 2 with a length of 4 nm was homogenously dispersed in the tubular carbon nanofiber. The high conductivity of the carbon nanofiber provided both an electron transfer channel and a well‐connected matrix for monolayer MoS2 nanoplates, which is shown in the 3D TEM model in Figure 2‐carbon nanofiber composites exhibited exceptional rate capacities  in carbon nanofibers is presented in Figure 2/C‐47 maintains the most stable cycling performance and rate capacity.2 nanocrystal and carbon nanofiber showed excellent cycling performance as an anode for LIBs.2, which is fostered in three different kinds of nanofibers: solid carbon nanofibers (SNFs), hollow carbon nanofibers (HNFs), and porous carbon nanofibers (PNFs).2 was grain stacking together on the surface of SNFs and HNFs very severely while this phenomenon seldom existed on the inside surface of PNFs.2 sheets, which are grown on the porous carbon nanofibers, was examined at different current densities. The specific capacity of MoS2/PNFs reached 1210 mAh g−1 at 100 mA g−1, which was higher than the previously reported value (1290 mAh g−1 at 50 mA g−1).2/PNF nanocomposites with MoS2 nanosheets fabricated on nanofibers are an excellent representation as the anode for LIBs. Miao et al. fabricated porous carbon nanofiber‐MoS2 composites with a core–shell structure for the anode of LIB.2 structure, which facilitates the diffusion of lithium ions. The capacity can reach 1532 mAh g−1 at 0.1 C. About 76.5% capacity remained because of the formation of the solid electrolyte surface (SEI) layers and the trapping of lithium in the lattice and defects of MoS2 after the first cycle. Then, the capacity retention ratio reached 100% in the subsequent cycles.2 hybrid nanocomposite.−1 and only a slight capacity decay was found after 100 cycles. However, the conducting polymer composites are not stable during cycling, which is ascribed to the non‐uniform dispersion and the relatively low conductivity of MoS2.The electrospinning technique has the advantages of low price, being easy‐to‐scale‐up approach, easy operability, and being a versatile technique to gain nanofibers for various functions. Through electrospinning, carbon nanofibers can be coated by MoS2 combined with organic materials is a new protocol to develop electrodes for LIBs. A good illustration of MoS2‐PANI nanowires (MoS2 66.7%: PANI 33.1%) composite is shown as a well‐designed hierarchical nanostructure under mild conditions. That delivered a high reversible specific capacity of 1063.9 mAh g−1 at 100 mA g−1 with a capacity retention ratio of 90.2% after 50 cycles.2 nanotube/nanofiber, MoO3 nanorods, or MnCO3, can be applied as a substrate or raw material to prepared 1D MoS2 nanohybrids with enhanced electrochemical performance. For example, compared to carbon nanofibers, mesporous TiO2 nanofibers exhibit outstanding flexibility and plasticity features, which can be used as main bone to support MoS2. For example, Ying et al. used mesoporous TiO2 nanofibers as templates to prepare MoS2‐TiO2 nanofibers by in situ vulcanizing the (NH4)6Mo7O24·4H2O precursor by H2S. The TEM and HRTEM images are shown in Figure2 (011) plane inside layer and MoS2 (002) as the outside layer. The typical charge/discharge profiles of TiO2 and MoS2‐TiO2 nanofibers are shown in Figure 2‐TiO2 nanofibers delivered a higher capacity (301 mAh g−1) than bare TiO2 (247 mAh g−1) and a longer discharge platform. The MoS2‐TiO2 nanofibers also exhibited stable cycling performance for 1000 cycles at 6C, which was much better than bare TiO2 and MoS2 electrodes, as shown in Figure 2 structure and the homogenously distributed MoS2 nanosheets.2‐coated TiO2 nanobelts (TiO2@MoS2) via the hydrothermal method andthe 1D structure is shown in Figure 2 on the TiO2 nanobelt and the layer distance of 0.62 nm. The TiO2@MoS2 hybrid maitains a high capacity of 710 mAh g−1 after 100 cycles at 100 mA g−1, which is four times higher than that of bare TiO2 nanobelt and three times higher than that of MoS2. That is related to the integerity of 1D nanostructure after long‐term cycles and the ultrathin MoS2 nanosheet shortening the diffusion paths of Li+ ions.2@TiO2 nanocomposites with thin MoS2 nanosheets grown on TiO2 nanotubes using a sol×gel method, as shown in Figure 2 nanosheets are homogenously attached to the surface of TiO2 nanotubes, which can be observed in the SEM and HRTEM images in Figure 2 on TiO2 nanotubes. The initial reversible capacity of this material was 931 mAh g−1 and it remained 578 mAh g−1 after 100 cycles, as shown in Figure 2@TiO2 is enhanced in electrochemical performance compared to the bare MoS2 nanosheets. Another example of a 1D MoS2 material is the typical hierachical hollow MoS2 nanotubes that are produced by a solvothermal reaction of Na2MoO4·2H2O, MnCl2·4H2O, and (NH4)2CS, followed by etching the MnCO3 nanotubes.−1 after 100 cycles at a current density of 100 mA g−1. The impressive electrochemical performance is possibly attributed to the special hierarchical surface and hollow tube structure by post‐mortem analysis. Zhang et al. used MoO3 nanorods as the precursor and in situ sulfidation by H2S/H2, and then added an additional carbon layer via a chemical vapor deposition (CVD) method to synthesize carbon‐coated MoS2 (C‐MoS2) nanorods as shown in Figure 2 nanorods. The as‐prepared carbon‐coated MoS2 nanorods exhibited the typical charge/discharge profiles of C‐MoS2 nanorods  and its unique structure gives the convenience of the intercalation and transportation of Li+ ions. To fully magnify this edge, 2D MoS2 is often synthesized with amorphous carbon, conductive polymers, and metal oxide decoration to do the surface‐treatment.2 nanowalls via a two‐step reaction with freshly prepared H2S and hydrothermal post‐treatment. It was also reported that the use of carboxymethyl cellulose (CMC) with a strong interaction for the conversion reaction electrode exhibits a good mechanical binding effect and stability compared to PVDF. The SEM and TEM images in Figure+ ions in the interlayer space at 1.1 V with around 1.8 Li+ intake, the conversion reaction with Li+ ions at 0.6 V with around 4 Li+ ions consumption, and electrolyte decomposition reaction or other reactions. The 2D MoS2 nanowalled electrode with CMC as the binder delivered a high capacity of 1047 mAh g−1 and maintained around 880 mAh g−1 after 50 cycles at 100 mA g−1  is prepared by the oleic acid (OA) pre‐functionalized MoS2 nanosheet reacting with dopamine and annealed at 850 °C for 2 h. The 2D nanocomposites piled by single layer MoS2 and mesoporous carbon was produced as an anode for LIBs. The ideal atomic interface contacts between two kinds of layers promising the maximization of synergistic interaction. The blue shift of peaks at 8.9° and 17.9° in the XRD pattern was attributed to the insertion of the OA molecules. Figure 2/m‐C nanocomposites delivered a high capacity of 1113 mAh g−1 at 400 mA g−1 and maintained 1023 mAh g−1 after 500 cycles. Additionally, the MoS2/m‐C hybrid also showed improved rate capability compared to bare MoS2 and graphene. As far as the electrochemical performance of MoS2/m‐C hybrid at extreme temperature, the hybrid exhibited much better rate capability and cycling performances at high temperature  than in a cool environment (3 °C), as shown in Figure 2/G (or MoS2/G/MoS2) at the interface because of the synergistic interaction as shown in Figure 2 layer is bonded with carbon nanosheets via van der Waals interactions. Wang et al. synthesized highly ordered mesoporous MoS2 using a vacuum‐assisted impregnation method, which has an expanded spacing of the (002) crystal plan that can be used in the LIBs  and good cycling performance over 100 cycles. Moreover, the highly ordered mesoporous MoS2 anode shows good rate capability in Figure −1 at the current density of 10 A g−1), which is much higher than the bulk MoS2 anode.2 is related to the homogenously distributed nanopores, high surface area, and the high pore volume, which are beneficial to the lithium transport and electrolyte diffusion. The intrinsic property of highly ordered mesoporous MoS2 also provides the short crystalline pathways for lithium ions, low activation energy phase, and wide layer distance at (002) plane. The combination of MoS2 sheets and graphene is another approach to avoid materials aggregation and enhance the conductivity of MoS2. Park et al. prepared ultrathin MoS2/graphene composites using the hydrothermal reaction and simultaneously reduced the graphene oxide with NaOH. The SEM and TEM images in Figure 2 in graphene (3%) nanosheets,2 with an interlayer of 0.62 nm at (002) plane and the typical (100) plane with a lattice of 0.194 nm. The comparison of rate capacity and cycling performance in Figure 2/graphene (3%) composite is much better than that of bare MoS2 or graphene, respectively, indicating that the introduction of graphene sheets can effectively absorb the inner mechanical strain, confine the volume changes, and enhance the conductivity of MoS2 nanosheets.As the special lithiation process of MoSs Figure l,m. The 2 monolayers through electrostatic interactions and then in situ turned to graphitic carbon through a pyrolysis process.2@N‐doped nanocomposites with a sandwiched structure can be obtained. Hu synthesized loose carbon‐coated MoS2 sheets via a one‐pot hydrothermal synthesis method, which exhibited a high capacity of 1419 mA h g−1 at 0.1 A g−1 and the capacity remained 80% after 50 cycles.2 nanoplates with a graphene‐like layered structure and a thickness of 30 nm via a simple, scalable, and one‐pot reaction method, showing a high reversible capacity of 700 mAh g−1 at 50 C. Additionally, a demonstration of a full battery (LiCoO2/MoS2) exhibited a stable cycling performance.Inspired by an adhesive protein that is usually found in mussels, biologic materials have been used to prepare various materials with special structures. Polydopamine (PD) has recently shown great possibilities for the surface modification of organic and inorganic materials via polymerization. As a good self‐assembly agent and a carbon precursor, PD can be absorbed on exfoliated MoS2, some metal nanoparticle permeating MoS2 materials are also demonstrated as an effective approach to enhance the electrochemical performance of LIBs. For example, combined with MoS2, the drawbacks of bare tin with a fast capacity fade and poor rate capability can be easily avoided, while its advantages such as low toxicity and high capacity still remain.Figure2/C nanomaterial. The Sn/MoS2/C anode delivered a high capacity of 1050 mA h g−1 at 100 mA g−1, which is attributed to the common contributions of Sn and MoS2 for the achieved capacity. The co‐existence of Sn nanoparticles can promote the growth of stable solid electrolyte interphase films and prevent the rupture of SEI during the charging process. At the same time, MoS2 nanosheets can suppress the aggregation of the Sn nanoparticles and keep them in a homogenous distribution state.2 nanosheets (MoS2‐PPy‐rGO) via the chemical bonding.2‐PPy‐rGO nanocomposites are shown in Figure −1 at a current density of 200 mA g−1 and there was about 1070 mAh g−1 remaining after 400 cycles ‐MoS2/NC/G hybrids shown inthe SEM and HRTEM in Figure2. The MoS2/NC/G hybrids delivered a high rechargeable capacity of 1570.6 mAh g−1 at the condition of 0.1 A g−1 and superior rate capability, which were higher than that of 2D bare MoS2 materials due to the well‐connected and open framework structure of the flower‐like and nitrogen‐doped PPy‐MoS2/NC/G hybrids, as shown in Figure 2 nanosheets using potassium sodium tartrate as an agent using a facile treatment, and the typical worm‐like morphology is shown in Figure 2 nanosheets was 845 mAh g−1 after 50 cycles at 100 mA g−1, as shown in Figure −1 at higher current rate of 500 mA g−1 after 100 cycles, which is associated with the large surface area (providing more exposed active sites and a large contact area), the free space between neighboring nanosheets , and the self‐interconnected structure (avoiding electrode aggregation).The 3D MoS2 nanoarchitecture prepared by a facile P123‐assisted solution‐phase approach and loaded into 3D graphene foam (GF) reported by Lin and Shen et al.,−1 at 200 mA g−1 and retaining 85.8% of the reversible capacity after 60 cycles. The corresponding SEM, TEM, and HRTEM are shown in Figure 2@GF materials are presented in Figure 2@GF contributed to the exceptional rate capabilities and good cycling performances. The synergistic interaction between MoS2 and graphene can prevent the severe expansion of MoS2 layers during discharge.2 can reduce the instinct of wrapping of graphene layers and keep its structured performance.2 in other carbon‐based materials can significantly improve the use efficiency of the active sites. Zhang and co‐workers carried out a specific investigation into the use of graphene and its derivatives and found many new characteristics.2 sheets on insulating substrates.2 and bulk MoS2, the sandwich MoS2@N‐doped carbon nanosheets exhibited enhanced lithium storage capability. The specific capacity of the aforementioned nanocomposites reached 1239 mAh g−1 and remained at 597 mAh g−1 at 10 A g−1.2/graphene composite using polydiallydimethylammonium chloride (PDDA).−1, the corresponding capacity of the as‐prepared materials was 1100 mAh g−1 with an enhanced cycling stability, and around 856 mAh g−1 remained after 900 cycles.Moreover, the honeycomb‐like structure with the interconnected net and robust mechanical property has recently attracted more attention. For example, the honeycomb‐like MoS2 nanomaterials significantly influence the electrochemical performance. Ma et al. compared the influencevarious types of cationic surfactants on of morphology changes of MoS2 nanomaterials, including dodecyltrimethyl ammoniumbromide, octyltrimethyl ammoniumbromide, and tetrabutyl ammoniumbromide.2 nanomaterials. Assisted by a type of N‐dodecylpropylenediamine γ‐diquaternium bromide, the tile‐like MoS2/graphene nanocomposites were prepared and are shown in Figure 2 (3–6) are homogenously anchored on the surface of the graphene sheets andexhibited high surface area, exposing tremendous active surfaces of MoS2 nanosheets. The rechargeable capacity of the nanotile‐like MoS2/graphene nanocomposites delivered a high reversible capacity of 1127 mAh g−1 at a current density of 100 mA g−1 and showed stable cycling performance in 220 cycles and good rate capacities at different current densities, as shown in Figure 2 nanosheets on the graphene nanosheets.2 nanosheets was reported by Zhao and He et al. , which can substantially improve the electrochemical and mechanical properties of MoS2 materials. The MOF‐derived carbon framework contains a nitrogen‐doped hybrid structure, with high conductivity and additional lithium ion storage. Yu et al. synthesized a C@MoS2 nanobox by fostering the ultrathin MoS2 nanosheets on N‐doped carbon shells and tested it on the LIBs. They thought that the nanosheets could provide more active sites for electrochemical reactions. In the first cycle, the charge and discharge capacity of C@MoS2 nanoboxes were as high as 1966 mAh g−1 and 1164 mAh g−1 and only about 59.2% was retained. The large capacity loss was ascribed to the decomposition of the electrolyte and the irreversible lithiation process, the formation of solid electrolyte interface. FigureThe morphology of MoS2, LiMn0.5Ni0.5O2, or LiFePO4) and exhibits low initial coulombic efficiency, the pre‐treatment of anode materials is an important approach to tackle this problem. Wang et al. synthesized an urchin‐like MoS2 composite using spherical onion‐like carbon core with a solvothermal and subsequent annealing treatment method. The new technique of pre‐lithiation of urchin‐like MoS2 composites by the direct contact with lithium foil is presented in Figure2 turn from 2H with semiconducting characteristics to 1T with metal features.2 composites is shown in Figure 2 powder, the urchin‐like MoS2 composite showed better electrochemical performance. The initial rechargeable capacity of urchin‐like MoS2 was 852 mA g−1 with a high Coulombic efficiency of 71.1% at the current density of 50 mA g−1 and the rate capability of urchin‐like MoS2 composite also demonstrated that the pre‐lithiation method can be applied to enhance the Coulombic efficiency for high‐performance LIBs  and long cyclability CNT also illustrated the enhanced electrochemical performance to the LIB anode. Moreover, as the formation of solid electrolyte interphase in the first few cycles consumes a large amount of lithium ions from the cathode  is much larger than that of Li+ ions (0.76 Å).+ ions, large volume expansion during sodiation, and pulverization of electrodes during cycling all result in the fast capacity decay and low Coulombic efficiency of SIBs. Therefore, many efforts have been devoted to address the problems of SIBs and develop appropriate cathode and anode materials. Graphite is usually applied as anode for commercial LIB, but it was found by Fouletier in 1988 that it is quite difficult to embed Na+ ions in the interlayers of crystalline graphite . MoS2, with its exceptional physical and chemical features, is of great potential for application as an anode for SIBs.Sodium ion batteries (SIBs) have attracted great attention in recent years as an alternative to the LIB rechargeable system because of the abundance of sodium resources, the even geographic distribution, and potentially low cost of the entire battery system. However, there are some disadvantages of SIBs that hinder their vast application. First, sodium ions in nature are almost three times heavier than lithium ions. Second, the radius of Na2 on SIBs, which sheds light on the material design and further industrial application. According the investigation of Ahn et al., the discharge/charge process of MoS2 in SIBs consists of two steps based on the inserted Nax amount . In the first stage, the crystal structure of MoS2 does not change due to the low inserted amount of Na ions. However, when the number of Na ions is in the range of 0.5 < x < 1.1, the layer structure of MoS2 continues expanding along the c‐axis, and causes slight structural distortion by inner strains, which is illustrated by the scheme in Figure2 cannot be fully recovered during charge due to the inner strains and irreversible decomposition during discharge.2 during sodium intercalation was investigated on the atomic scale using in situ XRD and aberration‐corrected scanning transmission electron microscopy by Chen and co‐workers, as shown in Figure + ions first insert in every other interlayer of the MoS2 nanocrystal, and then fully filled the other interlayer during discharge.2, 2H‐Na0.5MoS2, 1T‐Na0.5MoS2, and 1T‐NaMoS2, as well as NaxS, the large strain and high energy barrier cause the phase transition from 2H‐ to 1T‐MoS2 , demonstrating the typical few layer MoS2 nanosheets by TEM, HRTEM, and AFM  compared to the bulk MoS2 materials .2 nanopetals can maintain a capacity of 595 mAh g−1 at 0.1 C after 50 cycles with Na‐alginate as the binder, while the MoS2 nanospheres can only deliver a specific capacity of 520 mAh g−1 at 0.1 C. It is believed that exposed MoS2 edges mainly contribute to the well‐maintained capacity. A similar strategy to achieve monolayer MoS2 was used by Dou and Chou and co‐workers,2 composite with a carbon‐coated layer delivered a high capacity of 400 mAh g−1 at 0.25 C and maintained a long cycling life; most importantly, the carbon‐coated MoS2 nanosheets showed exceptional rate capability at 5 C (290 mAh g−1). Chen et al. prepared MoS2 nanoflowers with expended interlayers for SIB. By the control of cutting off voltage 0.4–3 V, the intercalation mechanism only took place, and the MoS2 nanoflowers delivered both high capacity and good cycling performance.The better understanding of the phase transitions along with the lattice distortion and irreversible structural decomposition are critical to discover the code of the electrochemistry of MoS2 Figure d. The sp2 Figure e,f. Moree Figure g. The thM Figure h–j.133 Ws Figure k. and co2 shows good electrochemical performance, the 2H/3R‐MoS2 with semiconductor properties should be enhanced in conductivity by a highly conductive intermediator, such as carbon nanofibers, carbon nanotubes, or graphene. The combination of MoS2 with graphene and carbon nanotubes can effectively enhance the electrochemical performance of SIBs, and the layered structure of MoS2 can be well‐maintained after long term cycles.2 grown on carbon nanotubes, which exhibited a specific capacity of 495.9 mAh g−1 at a current rate of 0.1 C and maintained 84.8% after 50 cycles. However, the 3D layered MoS2 nanosheets grown on graphene achieved a high capacity of 797 mAh g−1 at the same current rate (0.1 C) and about 99.98% of the initial capacity remained after 600 cycles, which demonstrates that the combination of layered graphene and MoS2 can maximize the advantages of the two materials and maintain a stable morphology after long term cycles.2 nanosheets via hydrothermal method for binder‐free electrode without using a conductive agent.2@C is shown in Figure2 grown on carbon fiber, nanovoids in MoS2 nanosheets, and the typical phase of (002). The MoS2@C electrode with one‐step dip‐coating CMC approach formed MoS2@C‐CMC electrode. The MoS2@C‐CMC electrode delivered a high reversible capacity of 356 mAh g−1 at a current density of 80 mA g−1, which was much higher than the MoS2@C electrode. Moreover, the MoS2@C‐CMC electrode also exhibited superior cycling performance in Figure 2@C‐CMC electrode is associated with the 3D interconnected carbon fiber for high conductivity and the ultrathin MoS2 nanosheets for fast Na+ ions diffusion. Zhu et al. synthesized interlayer‐expanded MoS2‐CNTs composite (MSCNTs) as an anode for SIBs using functionalized CNTs as backbones. The ultrathin MoS2 nanosheets with a space expanded structure were homogenously grown on the surface of CNTs , as shown in in Figure 2‐graphene spheres is ascribed to the 3D porous graphene microspheres offering enough inner voids for big volume changes, and an ultrathin layer of the MoS2 structure for fast Na+ ion diffusion.Although the graphene‐like monolayer MoSs Figure f; the Mo2 is another approach to enhance the electrochemical performance of the MoS2 electrode. Kang et al. synthesized SnO2‐MoO3 yolk–shell microspheres by a one‐pot electrospray method and then successfully vulcanized to the SnS‐MoS2 yolk–shell microspheres as shown in Figure 2, yolk–shell SnS and MoS2, yolk–shell SnS‐MoS2 microspheres exhibited a high capacity of 396 mAh g−1 after 100 cycles, while the capacity of dense‐structured SnS‐MoS2 only maintained 207 mAh g−1 in Figure 2 via an ultrathin and uniform HfO2 layer demonstrated by scheme, elemental mapping and TEM images in Figure 2, was for the first time raised by Ahmed et al.,2 coated MoS2 nanocomposites retained a high capacity of 636 mAh g−1 at the current density of 100 mA g−1, which was near the theoretical capacity of 670 mAh g−1, while the capacity of bare MoS2 only maintained 435 mAh g−1 at the 50th cycle in Figure 2 layer, which acted as the passivation layer and prevented the structural degradation of MoS2 layers. Meanwhile, the amorphous features of HfO2 allowed for the facile diffusion of Na+ ions during the cycles. Therefore, the understanding of the phase transformation, surface modification, and morphology design of MoS2 are of significant importance to enhance the electrochemical performance and shine light for the practical application of SIBs.Surface modification by metal sulfide and metal oxide on MoS2.32, a 2D layered material with ultrahigh surface area and mechanical properties, is regarded as one of the most promising candidates for supercapacitors. Table2. Fan and co‐workers successfully prepared the flower‐like MoS2/C nanospheres using a facile hydrothermal method. The specific capacitance of flower‐like MoS2/C nanospheres reached 201.4 F g−1 at a current density of 0.2 A g−1, demonstrating that layered MoS2 has the natural advantage on the application of supercapacitor.Supercapacitors, divided into double‐layer capacitors, pseudocapacitors, and hybrid capacitors by the energy storage mechanisms, are a new‐type of electrochemical capacitor with high capacitance. They bridge the gap between electrolytic capacitors and high‐energy‐density secondary batteries. Compared to traditional capacitors, supercapacitors have the characteristics of high energy density, high cycling life span, short charging time, high rate of power, and strong resistance to the environmental temperature. These advantages allow them to be applied widely on the portable electronic devices and the state of vehicles.2 nanosheets, and the chemically exfoliated MoS2 nanosheets exhibit outstanding electrochemical performance as the electrode in supercapacitors. For example, Acerce et al. chemically exfoliated MoS2 nanosheets and prepared flexible and free‐standing MoS2 films containing a high concentration of the metallic 1T phase as shown in Figure+, Li+, Na+, and K+) with a high capacitances ranging from 400 to 700 F cm−3 in different types of aqueous electrolytes  operation in supercapacitors with organic electrolytes, and exhibited high volumetric energy and power density values as well as a high Coulombic efficiency of 95% over 5000 cycles  thin‐film supercapacitor .2 are directly bonded with RGO via the aforementioned covalent chemical bonds, which were reversibly measured between 0.25 and 0.8 V in 1 m HClO4 for low concentrations of MoS2 and 0.25 V–0.65 V for medium and high concentration of MoS2 layers on graphene , which illustrated that the supercapacitance can be altered via the changes of MoS2 concentrations in hybrids. Similarly, Byun and co‐workers synthesized a MoS2/GO hybrid film by the addition of GO sheets, which performed as the phase stabilizer and mechanical supporter. The MoS2/GO hybrid film was treated by a post‐annealing process in air to form the high concentration of distorted 1T‐ MoS2 for the high conductivity. Therefore, the MoS2/GO hybrid film as supercapacitor electrodes exhibited enhanced electrochemical performance, such as high volumetric capacitance (≈380 F cm−3 at the scan rate of 10 mV s−1), ultrahigh volumetric power density (≈207 W cm−3), and a good cycling life span (95% retention after 10 000 cycles) in the condition of the aqueous electrolyte.2, and MoS2‐graphene composites) as shown in Figure −1 at a current density of 1 A g−1. The energy density was up to 73.5 Wh kg−1 when the powder density was 19.5 kW kg−1  due to the partial re‐exfoliation of the layered materials via ion intercalation.To investigate the influence of the synergetic effect between MoSe Figure j,k. The 1 Figure n. Good c2 nanocomposites were synthesized by Hu et al. using porous anodic aluminum oxide as the template. When applied as supercapacitor electrode, the tubular C/MoS2 nanocomposites exhibited a high capacitance of 210 F g−1 at a current density of 1A g−1 and superior cycling performance (over 1000 cycles).2 with the 2D graphene‐like structure, the MoS2/MWCNT nanocomposites were successfully prepared as shown in Figure2/MWCNT nanocomposites exhibited a specific capacitance of 452.7 F g−1 at a current density of 1 A g−1, which was much higher than that of bare MWCNTs (69.2 F g−1) and bare MoS2 (149.6 F g−1). Only about 4.2% of the capacitance was observed after 1000 cycles at a current density of 1A g−1 in Figure 2. The flower‐like MoS2 incorporated by carbon aerogel was synthesized by a facile hydrothermal route assisted by l‐cysteine. The MoS2/carbon aerogel hybrid film delivered a high capacitance of 260.0 F g−1 at 1 A g−1 in the condition of 1.0 m Na2SO4 aqueous solution,and exhibited a stable cycling performance .Additionally, tubular C/MoS2, conductive polymers are also used to modify the MoS2 sheets. Huang et al. prepared polyaniline/molybdenum disulfide (PANI/MoS2) using a facile in situ polymerization method, which provided direct paths for electrons via PANI  in Figure 2. The combination of PANI and MoS2 was also demonstrated by Liu et al. via the 3D tubular structure with the MoS2 nanotube as core and PANI nanowire arrays as shell that grew on both sides of MoS2 nanotubes. When the amount of the PANI loading on the MoS2 was 60%, the MoS2/PANI‐60 hybrid electrode exhibited a high capacitance of 552 F g−1 at a current density of 0.5 A g−1, which also showed exceptional rate capability of 82% from 0.5 to 30 A g−1. The MoS2/PANI‐60‐based symmetric supercapacitor illustrated good rate capability and cycling stability. Specifically, the capacitance based on the two electrodes was 124 F g−1 at the current density of 1 A g−1, and around 79% of the initial capacitance was well maintained after 6000 cycles, which demonstrated the stable tubular structure and the synergetic effects of PANI and MoS2 hybrid.In addition to the carbon based materials for modifying the features of MoSI Figure e.152 The2 is in favor of the structural stability of MoS2 and improves the electrochemical performance of supercapacitor electrodes. For example, the Ti/TiO2/MoS2 coaxial fibers with a fabricated TiO2 buffer layer and MoS2 nanosheets grown on Ti nanowires via a hydrothermal method exhibited flexible features and strong mechanical properties . The good electrochemical performance was attributed to the backbone of C/Ni nanorods on increasing rate performances and the interconnected MoS2 nanosheets on providing numerous accessible surfaces and the good contact with the electrolyte.3S4) with MoS2 nanospheres is illustrated by Zhang et al.  by a facile one‐pot process for the supercapacitor. The Ni3S4@amorphous MoS2 nanospheres exhibited a high capacitance of 1440.9 F g−1 when the current density was 2 A g−1 and exceptional cycling performance (90.7% remained capacitance was achieved after 3000 cycles at 10 A g−1) in Figure 3S2@MoS2 composite was synthesized by Wang et al. via a green and scalable one‐step method.3S2@MoS2 composite delivered a capacitance of 848 F g−1, which was almost twice that of pure Ni3S2 (425 F g−1). After 2000 cycles, the capacitance retention of Ni3S2@MoS2 composite was about 97%. This was attributed to the core–shell structure and the high surface area of Ni3S2@MoS2 composite. The metallic oxide nanoparticles endow MoS2 nanosheets with different properties. For example, Wang and co‐workers fabricated MoS2‐Co3O4 composites using laser ablation in liquids and an aging‐induced phase transformation method , which is converted from MOF  will be possible in the near future.In summary, we have introduced the unique properties of monolayer MoS


Here we show that an improved crystallinity and grain size of CH3NH3PbIxBrx3– films could stabilize these materials under one sun illumination, improving both the efficiency and stability of the wide‐bandgap perovskite solar cells. We further increased Br concentration in methylammonium halide precursor up to 64% (or a MAI:MABr weight ratio of 1:1.3)—the optical bandgap of MAPbBrxIx3– was increased to 1.75 eV. We assign the perovskite with this composition as MAPbBr0.8I2.2 based on the lattice parameter derived from the XRD pattern. The device parameters for this wider‐bandgap MAPbBr0.8I2.2 extracted from the red curves in Figure Voc = 1.21 V, Jsc = 15.8 mA cm−2, FF = 77.9%, and PCE = 14.9%.xIx3– devices studied in this work also showed improved photo‐stability, in striking contrast to what was reported previously.J–V curves. In order to find out the mechanism for the improved solar cell PCE and photo‐stability for these devices, we fabricated the control cells with PEDOT:PSS as the hole transport layer and all other layers fabricated with exactly the same protocol. It should be noted that no solvent annealing was applied, which is different from our previous study,J–V curves for these structured devices is shown Figure FF and larger photocurrent‐hysteresis when the scanning direction was flipped. As shown in Figure −2, which agrees with the J–V curve collected after the light‐stability test  of the films in Figure0.8I2.2 grains formed on PTAA are much large than those formed on PEDOT:PSS. The lateral size of the grains grown on PTAA are several times of the film thickness, while the grains grown on PEDOT:PSS are much smaller than the film thickness. It should be noted that the MAPbBr0.8I2.2 films formed here only went through the thermal annealing, rather than the solvent annealing, therefore the MAPbBr0.8I2.2 films on PEDOT:PSS have much smaller grains than what previously reported.2 and MAX. A wetting surface to OTPs, such as PEDOT:PSS, with small contact angle (θ) reduces the Gibbs free energy barrier for nucleation (ΔGhet) by a factor that determines by the contacting angle:Ghom is homogeneous nucleation energy barrier, C0 is the number of atoms per unit volume in the phase; C* is the concentration of critical sized nuclei; k is Boltzmann constant; T is temperature. For instance, a small θ of 10° reduces ΔGhet to be 1.7 × 10−4 ΔGhom, which dramatically promotes the nucleation and forms very dense nuclei on the wetting surface. While on a nonwetting surface where θ approaches 180°, the ΔGhet is comparable to ΔGhom, which suppresses nucleation and results in larger spacing between nucleuses and thus the formation of larger grains by the end of the initial stage of film drying. The followed thermal annealing induces the OTP grain growth. The formation of small grains on non‐wetting surface is not favored from energy point of view, and the grain boundaries tend to be in the out of plane direction to minimize the total grain boundary area. This mechanism is illustrated in Figure 0.8I2.2 perovskite film is as thick as 540 nm, the resulting MAPbBr0.8I2.2 films have reduced crystallinity and smaller grains in the region closer to the bottom of the films, but still large grains formed on the upper level. This may be explained by the fact that the substrate surface energy loses its influence on the grain formation at the upper level area when the film is too thick.In order to find the origin of the improved photo‐stability of MAPbBr0.8I2.2 determines its photo‐stability by comparing the MAPbBr0.8I2.2 films with different thickness on PTAA HTL. As shown in Figure0.8I2.2 film, which has large and crystalline grains, was stable under one sun illumination for 33 min, while the devices with thicker MAPbBr0.8I2.2 active layer showed quickly degraded photocurrent under illumination. The optical bandgaps of the MAPbBr0.8I2.2 films with different thicknesses are the same, as shown by the absorption spectra in Figure 2 in interdiffusion methods.Another piece of evidence was found to support the microstructure of MAPbBr0.8I2.2, the post‐illumination EQE reduced significantly for above bandgap excitation, while increased for below bandgap excitation (720–750 nm), which can be assigned to the photocurrent contribution from the new low bandgap phase. The very low EQE from this new phase indicates the strong trapping effect of it. Another notable finding is the photocurrent dropped more at the shorter wavelength (350 nm to 550 nm) range, which indicates the phase separation predominately occurs in the film region close to the PTAA side.Further optical and material structure study confirmed the influence of microstructure on the grain photo‐stability. The photo‐induced phase separation generates a low‐bandgap phase which should cause photocurrent contribution from this new phase. Figure 0.8I2.2 thin films with X‐ray diffraction (XRD) measurement and photoluminescence (PL) measurement, which were previously employed to identify the phase separation in this type of materials.Figure0.8I2.2 film (320 nm) and the unstable one (530 nm) before exposing them to strong illumination. Although only a small amount of bromide was introduced, a cubic Pm‐3m group is identified in the MAPbBr0.8I2.2 XRD pattern with the characteristic peaks of (100), (110), (111), (200), (210). The peak positions are identical, which agrees with our earlier claim that the composition of these films does not change with the MAPbBr0.8I2.2 polycrystalline microstructure. However, the increased (100) peak intensity compared to other peaks for the stable MAPbBr0.8I2.2 suggests that the thinner MAPbBr0.8I2.2 polycrystalline film is better oriented than the thicker one, which gives new microstructure difference between these two compositional identical thin films in addition to the crystallinity and grain size. The random oriented MAPbBr0.6I2.4 generally contains larger‐angle grain boundaries, which would enhance the halide migration to assist the phase segregation. By looking at the (200) peak pre‐ and post‐illumination of the two films  peak  peak position for MAPbBr0.6I2.4—which agrees with the previous study by McGehee's group.0.8I2.2 composition might have the lowest Gibbs free energy, which drives the formation of this phase automatically assisted by photoexcitation. To exclude the possibility that the peak splitting comes from the perovskite decomposition, we kept the illuminated thick films in the dark for two hours, and found the peak splitting can be reversed , 0, 111),1,0.8I2.20.5I2.5. The light stability of the good wider‐bandgap MAPbBr0.8I2.2 device is demonstrated by a steady photocurrent output at maximum power output point over 30 minute under one sun illumination. Microstructure difference between the photo‐stable and photo‐unstable devices is presented by the cross‐section SEM images of the MAPbBr0.6I2.4 active layer: the spatial homogeneous polycrystalline with large sized grains and the stacking layered polycrystalline with small sized grains, respectively. The PL and EQE spectral change, accompanied with XRD pattern comparison between the MAPbBr0.8I2.2 thin films with two different microstructures, indicate the enhanced crystallinity and grain size are favorable to retain the homogeneous phase for the mixed halide perovskite during the photoexcitation, thus maintain a stable photocurrent output under the device working condition. The phase instability for mixed halide perovskite was studied with the MAPbBrxIx3– infiltrate into porous TiO2 scaffold,xIx3– are crucial to achieve the stable wide bandgap perovskite solar cells. This explains why some large bandgap MAPbBr0.8I2.2 built on mesoporous scaffold device were not stable. Further studies on the role of the grain boundary's area and grain orientation in phase separation of mixed halide perovskite may contribute to a deeper understanding. This work demonstrates the potential of mixed halide perovskite to stay a reliable homogeneous phase in photovoltaic working condition as the wide bandgap light absorber in tandem application.In summary, we achieved highly efficient and stable planar heterojunction devices based on the wide bandgap MAPbBr2 and MAI1PbIx–x Precursor PreparationBr: MAI was synthesized using the method described by Michael M. Lee et al. MABr was synthesized by the reaction of methylamine with a concentrated aqueous solution of hydrobromic acid  at 0 °C for 2 h with constant stirring under nitrogen atmosphere followed by a crystalized, purification and dry process which was the same as the preparation of MAI. MAIx1–Brx precursor was prepared by mixing MAI and MABr in 2‐propanol for 1:1.3 weight blend ratio and 62 mg ml−1 concentration. PbI2 precursor was prepared by dissolving PbI2 in DMF with the different concentration 500 mg ml−1, 600 mg ml−1 and 700 mg ml−1.Film and Photovoltaic Device Fabrication: First, hole transport layer was prepared by spin‐coating 0.25 wt% PTAA (toluene) solution doped with 1 wt% F4‐TCNQ at 4000 rpm for 25 sec, and the as‐prepared film thermally annealed at 100 ºC for 10 min. We have already demonstrated that F4‐TCNQ is a suitable dopant to PTAA.x1–Brx films, 500 mg ml−1, 600 mg ml−1, and 700 mg ml−1 PbI2 DMF precursor solutions were spun coated on the hole transport layer at 6000 rpm for 35 sec. Then, the as‐prepared PbI2 film was dried in a hotplate at 100 °C for 5 min, and followed by spin coating the 62 mg ml−1 MAIx1–Brx precursor on top of the PbI2 layer at 6000 rpm for 35 sec with the subsequent 75 °C 15 min thermal annealing. This process allows the bromide inter‐diffuse into the perovskite structure. Then we increased the annealing temperature to 100 °C, lasting 90 min before cooling down to the room temperature. For photovoltaic devices, 2 wt% solution of PCBM in DCB was spun onto the annealed perovskite film and followed by additional 60‐min annealing. Then the device was completed by sequence depositing 20 nm C60, 8 nm BCP and 100 nm Al.Device Characterization: The photocurrent curves were measured under simulated AM 1.5G irradiation (100 mW cm−2) using a Xenon‐lamp‐based solar simulator . A Schott visible‐color glass‐filtered (KG5 color‐filtered) Si diode (Hamamatsu S1133) was used to calibrate the light intensity before photocurrent measurement. The device area was 7.5 mm2. EQE was measured with a Newport QE measurement kit by focusing a monochromatic beam of light onto the devices.Film Characterization: XRD measurements were performed with the PANalytical Empyrean Diffractometer with Bragg–Brentano parafocusing geometry, the 3 kW Cu Kα source, and the PIXcel 3D detector. The scan rate is 0.02 s per step and 0.5 s per step with an angular range of 10–60. The single path absorption was measured using an Evolution 201 UV–visible spectrometer (Thermo Scientific) with the scan rate 1 nm per step and 0.5 s per step in the range of 350–900 nm. The PL spectrum was measured by iHR320 Photoluminescence Spectroscopy at room temperature. A 532 nm green laser with an intensity of 100 mW cm−2 from Laserglow Technologies was used as excitation source in PL measurement. PL integration time is 5 secs and the slit is 0.12 mm. PL intensity was collected every 10 sec continuously for 30 min. The cross‐section SEM images were taken from the Quanta 200 FEG Environmental Scanning Electron Microscope (ESEM) using a field‐emission gun (FEG) electron source to scan the gold coated cross‐section area morphology.


In the case of the soft selenium Lewis acid PhSeCl, sequential activation of the alkyne bonds leads initially to an isocoumarin (1 equiv PhSeCl) and then to a tetracyclic conjugated structure with the isocoumarin subunit fused to a benzoselenopyran (3 equiv PhSeCl). Conversely, the reaction with the hard Lewis acidic borane B(C In these reactions, the nucleophilic center can be a wide variety of atoms including hydrogen,A particular theme within this area is the cooperative addition of FLPs across C−C π‐bonds Scheme , bottom,6F5)3 in a series of rearrangements of aryl‐functionalized diynes to form complex, extensively fused polyaromatic systems such as dibenzopentalenes and derivatives thereof.Other reactivities associated with the use of strong boron Lewis acids involve the synthetically useful carboboration reactions as well as relevant cascade reactions of unsaturated C−C bonds. Recent studies by the Erker and Yamaguchi research groups showcase the use of B(Cn−2 or n−4 oxidation state). Here we compare the reactivity of methyl‐2‐{[2‐(phenylethynyl)phenyl]ethynyl} benzoate (1) with the soft Se‐centered Lewis acid PhSeCl and the hard B‐centered Lewis acid B(C6F5)3 in the pursuit of obtaining novel heterocyclic frameworks such as isocoumarin derivatives. Indeed, isocoumarins are well known for their biological activity and their derivatives have many applications as antimicrobial,Indeed, with a few exceptions,1 was prepared by a palladium‐catalyzed cross‐coupling reaction according to literature procedures.1 in the presence of excess phenylselenyl chloride (in a 1:3 molar ratio) in CDCl3 was monitored by multinuclear NMR spectroscopy. After a few hours at ambient temperature, the 1H NMR spectrum showed the formation of a new compound with only aromatic protons, thus indicating the loss of the methyl ester. This observation was confirmed by slow evaporation of the solvent to yield a crop of orange crystals (37 %), which were characterized by single‐crystal X‐ray diffraction and unequivocally confirmed the product to be the salt 2 ° and 100.86(12)° and with Se−C bonds having single bond character (1.912(2)–1.939(3) Å). These observations indicate that the delocalization of the isocoumarin unit does not fully extend throughout the isoselenochromene fragment. Indeed, the lack of planarity across these two moieties further substantiates this observation. The counteranion [PhSeCl2]− adopts a T‐shaped geometry with the chlorine atoms occupying the axial positions. The Se−Cl bond lengths (2.445(1) and 2.457(1) Å) and Cl−Se−C bond angles of 90.47(10)° and 90.72(10)° are similar to those described for the only previously reported [PhSeCl2]− anion.The tetracyclic core of 1. When a 1:1 molar ratio was used, selective activation of the alkynyl functionality proximal to the ester occurred through a 6‐endo‐dig cyclization to afford the phenylselenyl‐substituted isocoumarin 3 a  was in agreement with previously reported selanyl compounds.The reaction was then performed with 1:1 and 2:1 molar ratios of PhSeCl/3 a with two further equivalents of PhSeCl. These studies revealed 1H NMR signals corresponding to the formation of 2, thus suggesting that 3 a is an intermediate en route to 2. A plausible mechanism for the reaction of 1 to 2 via 3 a is the electrophilic addition of PhSeCl to the alkyne in 1 with formation of a selenirenium cation. This then undergoes an intramolecular ring‐opening reaction by nucleophilic attack from the carbonyl oxygen atom of the ester moiety, thereby resulting in isocoumarin 3 a with a concomitant loss of chloromethane, as observed in the in situ 1H and 13C NMR spectra .3 a to give the zwitterion 2 may be envisaged as an FLP‐type trans‐1,2‐addition of the Lewis basic selane in 3 a and a second equivalent of the Lewis acidic PhSeCl across the second alkyne (again via the three‐membered selenirenium cation). The chloride ion that is eliminated following the 1,2‐addition is intercepted by the third equivalent of PhSeCl to generate the selenium‐containing counterion [PhSeCl2]−.To probe this reactivity further, multinuclear NMR spectroscopic studies were undertaken to follow the reaction of compound 3 a using the harder Lewis acid B(C6F5)3 failed to lead to the tetracyclic structure reminiscent of 2. Instead, this afforded the Lewis acid/base adduct 3. The stoichiometric reaction of B(C6F5)3 with 1 (1:1) in toluene resulted in an immediate color change to a dark‐green solution, which afforded block‐shaped red crystals (84 %) upon standing for several hours. Single‐crystal X‐ray diffraction analysis revealed the product to be the conjugated compound 4 (Scheme 4 from the cyclization of 1 with the Lewis acid B(C6F5)3 is in stark contrast to the formation of 2 from 1 by using PhSeCl. Although bio‐ and biomimetic domino reactions are often acid‐catalyzed, the use of boron compounds in Lewis acid mediated annulation reactions is comparatively rare and remains largely unexplored.To compare the Lewis acidity of PhSeCl with more conventional Lewis acids, we investigated the reaction of 4 consists of an isobenzofuran bicycle connected by a double bond to an indene moiety. Although the isobenzofuran heterocycle and the indene ring are conjugated, they do not lie perfectly in the same plane in the solid state, with the angle between the two rings being 17.99(11)°. The fact that this system does not lie in the same plane clearly reflects the steric constraints in the molecule, which is also revealed in the NMR data. In addition, the phenyl substituent in 4 also rotates out of the plane of the indene ring by 67.87(10)°, thus permitting it to adopt a near coplanar geometry (12.52(9)°) with one of the C6F5 rings of the B(C6F5)3 moiety with a centroid⋅⋅⋅centroid distance of just 3.4785(16) Å. This value is somewhat shorter than that observed for other C6F5⋅⋅⋅C6H5 contacts (3.6421(10) Å).The structure of 11B NMR chemical shift for 4 in solution occurs at δ=−15.8 ppm as a very sharp singlet, which is synonymous with four‐coordinate vinyl borate species.19F NMR spectrum is more complex and indicates chemical inequivalence of the three C6F5 rings, thus suggesting some degree of restricted rotation around the boron–indene bond brought about by the coordination of B(C6F5)3 to the bulky organic fragment in 4. Three discrete resonances are observed in the 19F NMR spectrum for the para‐ and meta‐fluorine atoms: δ=−160.7 ppm (para‐F) and δ=−165.5 (meta‐F). In addition, the ortho‐positions within each C6F5 ring are also inequivalent, thereby resulting in six distinguishable signals between δ=−123.0 and −136.1 ppm. This steric hindrance is also observed in the space‐filling model (see the Supporting Information), which clearly shows a very sterically encumbered boron center in which the arene ring of the indene unit interlocks between two C6F5 rings.The In conclusion, the reactions of soft and hard Lewis acids with dienyl esters show contrasting reactivity that is unprecedented in FLP chemistry. The divergent reactivity of these systems allows novel heterocyclic systems to be generated that incorporate main‐group elements. To date, the addition reactions of FLPs to alkynes have been limited to hard Lewis acids, with the use of softer Lewis acids in FLP‐type chemistry being comparatively rare. Indeed, we are continuing to explore the utility of softer Lewis acids in FLP‐type reactions.As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


By using a common protocol over a wide range of experimental conditions, including very high magnetic fields and very high radio‐frequency (RF) nutation rates, decoupling performance is observed to degrade consistently with increasing magnetic field. Inhomogeneity of the RF field is found to have a significant impact on T2'values, with differences of about 20 % observed between probes with different coil geometries. Increasing RF nutation rates dramatically improve robustness with respect to RF offset, but the performance of phase‐modulated sequences degrades at the very high nutation rates achievable in microcoils as a result of RF transients. The insights gained provide better understanding of the factors limiting decoupling performance under different conditions, and the high values of T2'observed  provide reference points for experiments involving spin magnetisation refocussing, such as 2D correlation spectra and measuring small spin couplings.Factors affecting the performance of  Such experiments are central to the use of NMR spectroscopy for establishing molecular structure and dynamics. Moreover, measurement of small J couplings, such as those across N−H⋯N hydrogen bonds, provides direct information on molecular assembly.1H spins leads to significant magnetisation losses during these periods, directly affecting the viability of experiments. For example, the refocussed INADEQUATE experiment13C spectra of testosterone solid forms13C labelling. Isotopic enrichment was also used when probing biopolymers in secondary plant cell walls15n experiments),Effective decoupling of 13C magnetisation decays as a result of imperfect decoupling. Although considerable progress has been made in developing approaches to decoupling and understanding how they work,quantitative reproduction of experimental data through simulation is intrinsically difficult owing to the rapid population of high‐order coherences. Existing experimental studies focus on the important goals of improving spectral linewidths, often through new sequences, including studies of how to choose between sequences.The viability of experiments exploiting J couplings in organic molecules is directly related to the rate at which T2', is chosen as the primary experimental metric because it is unaffected by inhomogeneous contributions to the spectral linewidth, such as shimming, anisotropic bulk magnetic susceptibility or sample inhomogeneity.T2'continues increasing as the decoupling efficiency increases even though the spectral resolution has plateaued.T2'values tend to vary significantly even though the differences in spectral linewidths at the same conditions are small. So while optimising spectral resolution is relevant for many applications, T2'provides more insight when trying to understand the factors determining decoupling performance, and is directly relevant to the challenging experiments discussed above. The T2'values observed here suggest that J couplings as small as a few Hz are measurable and that J‐based 13C correlation experiments should be viable for most systems without the need for isotopic enrichment.As in our earlier study,13C to decouple, ensuring that decoupling that is effective on methylene sites will also be generally effective. Sample data sets acquired by using the methine group of l‐alanine showed similar trends (see Table S1 in the Supporting Information for data sets available), and the conclusions are thus expected to be generally applicable to organic solids.In common with several previous studies, the methylene group of glycine was used as a model system. The strong dipolar coupling network, both homonuclear and heteronuclear, in methylene groups makes them the hardest type of T2'were performed on polycrystalline samples of glycine‐2–13C,15n  and l‐alanine‐2–13C,15n  purchased from CortecNet. The glycine sample was confirmed to be α‐glycine based on the 13C carbonyl peak at 176.5 ppm, which is sensitive to polymorphic changes.Experimental measurements of 13C magnetisation was created by using cross‐polarisation (CP), ramped on the 1H nutation frequencyνr=12kHz, and through the p=+1zero‐quantum sideband (ν1H-ν1C=pνr) for νr=25kHz. At νr=62.5kHz, CP matching was done on the p=+1double‐quantum sideband (ν1H+ν1C=pνr), as this required much lower RF powers compared with the p=±1zero‐quantum sideband sidebands commonly employed at slower MAS, but gave comparable signal enhancement.13C magnetisation owing to the effects of RF inhomogeneity.T2'decays when using direct excitation compared with CP, presumably associated with a poorly decoupled sample towards the coil ends, but the long‐term decay and overall fitted T2'values were not significantly different. Although low‐power CP at high MAS rates has been reported to selectively excite the 13C spectrum,13C transmitter on the methylene resonance. Relaxation delays were always 4 s, whereas the CP contact times were optimised for each set of experiments, varying in the range 0.8–2.7 ms, as noted in the figure captions. The magnetisation was then measured after a spin‐echo period, τ–π–τ, during which either CW, two‐pulse phase‐modulated (TPPM),ϕ, or additionally optimising the angles αand β.Table 1H decoupling was used in both spin‐echo and acquisition periods. The only exception was early measurements using the microcoil probe, hardware configuration 8, for which optimised TPPM decoupling at ν1=105kHzwas used during acquisition. As discussed in more detail in Ref. T2'values through the orientation dependence of decoupling efficiency.1H transmitter frequency was positioned on the maximum of the unresolved 1H spectrum at slow MAS. Inversion of 13C magnetisation used a π‐pulse with duration between 5 and 6 μs.Generally, the same 1H decoupling nutation rate, ν1, was measured by using the same sequence with a zero spin‐echo period. The initial 1H pulse width was incrementally increased in at least 100 steps of approximately 1/(2.5ν1)to acquire a 1H nutation spectrum and the peak position was taken as the nominal ν1. TPPM and SPINAL‐64 pulse widths are expressed as a flip angle, θ=τpν1360∘, whereas XiX pulse widths are expressed as a fraction of the rotor period, τp/τr.The ν1=nνr, where nis an integer. Although the PISSARRO‐5Combinations of MAS and decoupling RF nutation rate were chosen to avoid rotary resonance recoupling conditions at 2τ, linearly in 30–40 steps from zero to approximately twice the maximum expected T2'. The free induction decays were zero‐filled and Fourier transformed (without apodisation) using matNMR.13C peak height as a function of 2τwas fitted to a decaying exponential to obtain T2'using MATLAB.T2'values were inferred from a pair of experiments at 2τ=0and 2τ≈T2,max'by assuming a mono‐exponential decay of the peak height between these points. As previously discussed in Ref. T2'values obtained by this quick, but approximate, approach, and those obtained from full decays were corrected by scaling the approximate values to coincide with the more accurate values obtained from full decays. The experiments also provided data for estimating “T2*”, the effective time constant describing the linewidth, calculated from 1/(πFWHM), where FWHM is the full width at half‐maximum.Full decay curves were obtained at selected decoupling conditions by incrementing the evolution time, 1H Larmor frequencies of ν0H=300, 600 and 850 MHz, at 12 kHz MAS rate and 105 kHz 1H nutation rate, using the same 2.5 mm o.d. rotor and similar probe designs. Additional data was collected at ν0H=500MHzunder the same MAS and decoupling nutation rates using hardware configuration 2. For TPPM and SPINAL‐64, full T2'parameter maps were first acquired as a function of both pulse width and phase to locate the positions of optima, then a detailed parameter cross‐section at a fixed pulse phase was recorded through the region of peak decoupling, which was ϕ=6∘for both sequences under these conditions. The XiX performance was first characterised over a wide range of pulse lengths, and then in more detail over the region of peak decoupling to ensure the narrow optima were well defined.Experiments to assess the dependence of the decoupling performance on magnetic field were performed at Decoupling experiments using very high RF decoupling fields, beyond the reach of commercially available probes, were performed using a piggyback μMAS design equipped with a 450 μm inner diameter coil, as described in Ref. B1homogeneity. The same sample rotor, MAS rate, and spectrometer equipment were used in both sets of experiments, and care was taken to adjust the power levels for each probe such that the peak nutation frequencies were the same.Experiments to assess the influence of RF field inhomogeneities on decoupling were performed using two 3.2 mm MAS probes with different coil geometries: an Agilent T3 MAS probe with standard solenoid coil geometry, and an Agilent BioMAS probe, whose scroll coil geometry exhibits significantly better 13C magnetisation in the presence of 1H decoupling were performed using pNMRsim,n, with nindicating the number of protons coupled to the central carbon atom. The decays were fitted to a single‐exponential function to derive computed dephasing time constants, T2c, describing the loss of the 13C single‐quantum coherence. Note that spin‐echoes are not included in the simulation as there are no inhomogeneous components of the decay to refocus, and computations of T2'are much more sensitive to finite‐sized spin systems.Numerical simulations of the decay of 2T2*and T2'. As these were acquired under a uniform set of conditions, these data should be a useful resource for further study. A summary of the data sets available is given in the Supporting Information.Determining the optimal decoupling parameters across a range of experimental set‐ups involved acquiring a large number of detailed parameter maps for both T2'.νr=25to 62.5 kHz under ν1=170 kHz decoupling. As illustrated by Figures S9–S11 in the Supporting Information, the parameter maps became fragmented by multiple destructive resonance conditions, requiring finer parameter grid increments. For TPPM and SPINAL‐64, the optimum pulse phase moved away from the commonly prescribed ϕ=7∘and ϕ=6∘,An initial comparison of the data highlights that ease of optimisation varies significantly between sequences and experimental conditions. These aspects of optimisation have previously been discussed for peak heightT2*parameter map qualitatively correspond with those of the T2'map.T2'optima were, however, generally narrower, and optima with similar T2*values tended to have different relative T2'values. In other words, a T2*map cannot be relied upon to provide the best sequence parameters for T2'.As has been previously observed, the local minima and maxima across a 2.1T2'.1H offsets ±10kHzabout the optimum, thus the offset dependence could be determined independently using the same optimal sequence parameters. Figure 1/(πT2*)and 1/(πT2')for TPPM across a range of experimental parameters. CW and SPINAL‐64  exhibit the same behaviour. With the exception of XiX decoupling, an increase in the B0field is accompanied by a consistent increase in linewidth for both 1/πT2*and 1/πT2', that is, a vertical offset. XiX has a more complex offset dependence, being more broadband at larger B0.The detrimental effects of off‐resonance irradiation on decoupling efficiency are well‐established for spectral linewidth,13C linewidths on transmitter offset under CW decoupling was shown to be described by a parabola, and was justified theoretically by VanderHart et al.,1/(πT2')under the same conditions. The bparameter shown in Figure bparameter. The 1/(πT2')curves do not fit well to simple parabolas, but show qualitatively similar trends. Although these observations are interesting and worth further investigation, they are not directly relevant to the complex interplay of experimental conditions and spin system dynamics at the heart of the decoupling problem. The following results assume that the 1H transmitter offset is close to the optimal conditions.The dependence of  al.,1/πT'under t2.21H nutation rates (12 kHz and 105 kHz respectively) were acquired at several magnetic fields. Figure B0fields studied. The decoupling optima marked for each decoupling sequence in Figure 1/ν0H. The trend in Figure T2'values with increasing B0for CW and TPPM decoupling is consistent with their decoupling performance being dominated by second‐order cross‐terms between the heteronuclear dipolar couplings and 1H chemical shift anisotropy (CSA) tensor, which increase proportionately with B0.B0dependence. Although the detrimental effects of increasing B0upon T2'can be inferred from previously published dataT2'at 600 MHz compared with 300 MHz have been noted for XiX and RS‐HEPT decoupling,T2'values in the limit of infinite magnetic field are likely to be an artefact of increasing off‐resonance effects at higher field; this behaviour is not borne out in the wider collated data of Figure S6. Moreover, unlike TPPM, the SPINAL‐64 parameter map, Figure B0, and so it may be more difficult to observe consistent trends for more complex sequences. This may explain, for example, why SPINAL‐64 appeared to perform better at 700 MHz compared with 500 MHz in a previous study,Parameter maps of decoupling performance at moderate MAS and T2'values for all four B0fields, although the peak performance at high field is measurably poorer. The position of the XiX optimum is very sensitive to ν1under this combination of νrand ν1, in the same way as for low‐power XiX decoupling.τp/ν1=8, corresponding to τp/τr≈0.91; owing to their proximity to destructive resonances, the XiX peak T2'values are much more sensitive to ν1mis‐adjustments than TPPM or SPINAL‐64. Therefore, the peak XiX T2'values may not be robust, especially at ν0H=850MHz, where the RF inhomogeneity is noticeable poorer . The consensus of these and previous results is that achievable XiX performance does decrease as the magnetic field increases.The linewidth under XiX decoupling, and related sequences such as PISSARRO‐5, is dominated by proximity to resonance conditions as well as second‐order cross‐terms between the heteronuclear and homonuclear dipolar couplings,T2'values at two magnetic fields (corresponding to ν0H=500and 850 MHz) and two MAS frequencies (νr=25and 62.5 kHz) for which complete data sets were obtained for all the decoupling sequences used. It can be seen that the trends of decreasing T2'with B0at νr=25 kHz (solid lines) are similar to those at νr=12 kHz in Figure B0as the other sequences. However, under fast MAS (dashed lines), when the two frequencies, νrand ν1, become more comparable, the direct impact of B0on T2'is less clear. This can be attributed to the increased significance of resonance conditions on the parameter space, complicating interpretations of T2'at optima based on a single dominant mechanism by making the optima narrow and very sensitive to small changes in νrand ν1. Unlike the other sequences considered, the XiX T2'either improves or stays unchanged with increasing MAS rate.Figure νrfor slower spinning frequencies,T2'values for the TPPM and SPINAL‐64 decoupling sequences were observed to increase up to 20 kHz MAS before falling off, whereas the performance of XiX and CW systematically increased and decreased, respectively, with fast spinning . Although the qualitative picture that emerges is consistent, it is worth noting the actual values of T2'vary markedly between studies (with the exception of simple CW decoupling). The dependence of T2'values on the optimisation protocol, and potentially, hardware details, illustrates the difficulty of drawing conclusions based on individual studies.These results complement previous studies of decoupling performance as a function of 2.3B1field inhomogeneities are potentially significant for phase‐modulated decoupling sequences, such as TPPM, where the optimum pulse length is strongly dependent on the 1H nutation frequency.B1fields across a sample produce a distribution of scaling factors, dramatically degrading resolution,B1, such as SDROOPYf‐TPPM.B1inhomo‐ geneity on T2'values has, however, not been explicitly investigated.The effects of 1H nutation spectra measured through 13C for two 3.2 mm MAS probes with different coil geometries. The measured nutation spectrum of the T3 probe  clearly shows a peak at ν1=93kHzand a long tail of much‐reduced RF, which is known from B1inhomogeneity imaging experiments (on a Bruker 4 mm probe) to originate from the sample at the ends of the rotor.B1field across the sample. Owing to the additional effects of CP selectivity, determined by the quality of RF matching between the two channels,13C CP/MAS spectra. The nutation spectra show additional peaks appearing at frequencies nνrand ν1±nνr, where nis an integer and νr=12kHz. These features were observed with other combinations of probes and experimental conditions used and are likely to result from a time‐dependence of the magnitude and direction of the effective field owing to anisotropic interactions, which are modulated by the MAS frequency.Figure ν1=93kHzin terms of both pulse duration and phase, and a detailed parameter map cross‐section then recorded at the optimum phase  and y(quadrature) components of the magnetic field, B1(t), following a phase change at t=0are [Eq. RF transients can be measured experimentally by pick‐up coils placed close to the NMR sample. At least in the case of the dipolar recoupling experiment considered by Carravetta et al.,are Eq. :(1)Bx(t)B1=2πν1/γis the amplitude of the driving RF, and Δϕis the phase change relative to the initial xphase. λtransis the rate constant for the transient response of the tuned circuit, which is largely determined by the Qof the probe, and λQparameterises the amplitude of the quadrature component of the transient response. Although λtransis essentially fixed by the probe, λQis largely determined by the mismatch between the frequency of the driving RF and the resonant frequency of the tuned circuit. Bx(t)and By(t)are readily converted to an overall RF amplitude and an instantaneous phase as a function of time. Illustrative examples of RF transient profiles can be found in recent literature reports.where τstep, typically of 1/2λtrans. Calculating propagators for each τstepmaking up an individual pulse would be extremely time‐consuming, and so the response is only modelled over the first 5 to 6 time constants (i.e. t≤6/λtrans), and the set phase and amplitude used for the remaining pulse duration. The convergence of the free induction decay (FID) with respect to both parameters was checked on a case‐by‐case basis. The pulses were always sufficiently long compared with 6/λtransso that overlap of the transient responses did not occur. Using a low‐Qpick‐up coil, values of λtrans=4μs-1and λQ=0.8μs-1were obtained by fitting the experimental transient response of a 2.5 mm 400 MHz DR Bruker probe to a ϕ=90∘,-90∘pulse pair.The transient response is modelled in the simulations by dividing the evolution into short time steps, νr=12kHzand a range of nutation frequencies up to ν1=500 kHz. Under conditions of ν1≫νr, the TPPM pulse width and phase were easily optimised, as shown by the relative sparsity of the resonance conditions across the parameter maps in Figure τp180/cos(ϕ/1.1289), close to the τp180/cos(ϕ)predicted by Floquet analysis.The effects of transients are expected to be larger at the high nutation rates routinely available when using a microcoil MAS probe.T2'values as a function of nutation frequency for TPPM, SPINAL‐64 and CW. Above v1≃300 kHz, the peak T2'values for the phase‐modulated sequences decrease sharply towards those of simple CW decoupling. Other experimental measurements using this probe have shown a similar drop in peak T2'for both TPPM and SPINAL‐64 above v1≃250 kHz.Figure λtrans=ω0/2Q=28μs-1is expected for the microcoil probe based on its Q1H). The value of λQdepends on the exact probe tuning and, in lieu of precise measurements, the λQ=0.8μs-1value measured on the Bruker 2.5 mm probe was assumed to be representative. To find the optimum T2cat a given nutation frequency, care was taken to ensure that the phase and pulse width simulation step sizes were small enough to show smooth evolution of dephasing times. Figure T2cat ν1≥200 kHz is due to the combined effects of in‐phase and quadrature transients. The in‐phase transients have little impact on their own, at least for TPPM with ν1≫νr. This agrees with our understanding that only a small reduction in dephasing time, owing to a minor reduction in the average RF amplitude, will be observed if the effective nutation axis of the sequence remains in the x-yplane. This is true in the presence of solely amplitude transients. If the nutation axis tilts out of the x-yplane as a result of the presence of both amplitude and quadrature transients, however, then a larger impact on dephasing times is expected, especially at high RF where decoupling optima are narrow.Simulations were performed with and without transients over a range of RF nutation frequencies. A value of 3T2'values observed here with previous literature results.Although heteronuclear decoupling has been widely investigated, the diversity of previous studies has made it difficult to make quantitative comparisons. The difficulty of optimising multi‐parameter decoupling sequences, particularly in regions where the decoupling optima are narrow, means that the results presented in the prior literature are often inconsistent. Measuring decoupling performance by using a well‐defined protocol and a wide variety of experimental conditions has provided large data sets that can be mined to address different questions. Here, we focus on understanding what are the limiting factors determining decoupling performance under different experimental conditions. Table 1H chemical shift anisotropy, but is less expected for sequences such as XiX decoupling (which is primarily limited by purely dipolar terms). This is likely to reflect higher‐order terms involving the 1H CSA and also offsets in 1H NMR frequencies from the decoupling frequency. In practice, the greater magnetisation losses at higher field will be largely offset by the intrinsically greater sensitivity and resolution of spectra obtained at higher field. Note that large 1H CSAs, such as those often observed in amide groups,The decrease in decoupling performance with increasing magnetic field was expected for decoupling sequences, such as CW and TPPM, which are primarily limited by cross‐terms between the heteronuclear dipolar coupling and the 13C linewidths, decoupling performance measured by T2'values is systematically worse at MAS rates above 60 kHz compared with 25 kHz. This reflects the increased number of “resonance” conditions in the area of parameter space where decoupling is typically optimal when the MAS rate is of the order of the 1H nutation frequency. “Low power” decoupling offers distinct advantages in these regimes; indeed, an impressive T2'of 200 ms has been observed on a similar test sample (glycine ethyl ester) using a modified form of XiX decoupling at 90 kHz MAS.T2'values increase with spinning rates below approximately 22 kHz MAS.As has previously been observed with 1H nutation rates. At nutation rates above 300kHz, however, simulations and experiments show that phase transients have an increasing impact, particularly from their quadrature components. At nutation rates of 500 kHz or more, achievable in microcoils, the performance of phase‐modulated sequences decreases dramatically towards that of simple CW decoupling. These problems can be addressed by careful tune‐up to minimise quadrature transients and/or development of sequences that are robust with respect to transients.RF transients associated with phase switches are not found to have a significant impact at typical T2', essentially by “smoothing off” peak decoupling conditions. Probes with flatter homogeneity profiles produce T2'decay curves that are closer to exponential and with measurably longer T2'values. Such factors contribute to the difficulty of reproducing T2'quantitatively in simulation,λ/4approaches the dimensions of the transmitter coil. The option of restricting the sample to regions with a more uniform RF profile is relatively unattractive, as improvements in sensitivity owing to longer T2'will be more than offset by the overall loss of signal. Coil geometries with flatter RF profiles, such as the end‐compensated coils introduced by Yannoni and co‐workers,T2'values obtained.Inhomogeneity of the radio‐frequency field has a significant impact on T2', of Cαin glycine summarised in Table T2'can be readily estimated from the reduction of signal intensity relative to a reference experiment without the refocussing period, as used here to acquire parameter maps efficiently. These reference values can be used to judge the scope for further optimisation.The peak values of the time constant for nuclear spin decay, rCWf‐TPPMf‐TPPM‐sc,T2'of 100 ms allows couplings as small as a few Hz to be measured and means that challenging correlation experiments, such as refocussed INADEQUATE, are viable, even at natural isotopic abundance.Although the best results were mostly obtained here using straightforward two‐pulse phase‐modulated (TPPM) decoupling, “good enough” decoupling for a given application may be achievable by using more easily optimised sequences. For example, As a service to our authors and readers, this journal provides supporting information supplied by the authors. Such materials are peer reviewed and may be re‐organized for online delivery, but are not copy‐edited or typeset. Technical support issues arising from supporting information (other than missing files) should be addressed to the authors.SupplementaryClick here for additional data file.


The morphology, phase composition and microwave absorbing properties of as-prepared samples were characterized by scanning electron microscopy (SEM), X-ray diffraction (XRD), vector network analyzer (VNA), and so on. The results revealed that the 3D network structure was configurated with smooth spherical walls composed of NiFe2O4 nanocrystals and their pore diameters being in the range of 80–250 nm. The microwave absorption properties of the 3D network structure NiFe2O4 were crucially determined by the special structure. The synergy of intrinsic magnetic loss of magnetic NiFe2O4 and the interfacial polarization enhanced by 3D network structure and the interaction of multiple mechanisms endowed the sample with the feature of strong absorption, broad bandwidth and lightweight. There is more than one valley in the reflection loss curves and the maximum reflection loss is 27.5 dB with a bandwidth of 4 GHz. Moreover, the 3D network structure NiFe2O4 show a greater reflection loss with the same thickness comparing to the ordinary NiFe2O4 nanoparticles, which could achieve the feature of lightweight of the microwave absorbing materials.3D network structure NiFe Very recently, electromagnetic interference (EMI) as a new way of pollution has attracted people’s great attention, which results from the explosive growth in application of electronic devices, radar systems and local area networks1et al. proposed a honeycomb SnO2 microwave absorbing material with an excellent dielectric loss and the optimal reflection loss (RL) is −37.6 dB obtained at 17.1 GHz with thin thickness of 2.0 mm18et al. introduced a mesoporous NiFe2O4 material that had high surface areas, tunable pore sizes, and periodic pore arrangements, showing optimal reflection loss of −22.5 dB with a band width less than 2 GHzNumerous microwave absorbing materials have been explored during the latest years and the microwave absorbing materials are now expected to possess the feature of strong absorption, broad band and lightweight. It is well known that the absorbing property is closely related to the electromagnetic parameter, impedance match and the microstructures371012142O4 materials with the configuration of ordered meshes exhibit an excellent microwave absorption performances due to the match of magnetic loss and dielectric loss.In this work, a templated salt precipitation method was adopted to prepare the 3D network structure nickel ferrite. The obtained 3D network structure NiFe2O4 NPs. Therefore, we didn’t discuss this part anymore.It has been studied that metal precursor involving polar solvents can infiltrate the colloidal crystals templates more completely, and introducing fewer structural defects during the composite formationTemplates based on neat array are the basis of the integrated 3D network structure. x gas. Simultaneously, the correlative exothermic peak on the DSC curve is related to the formation of glyoxylate2O4. According to TG-DSC measurements, it could be determined that the minimum feasible template-removing temperature is 400 °C.In order to investigate the thermal behavior of PMMA CCTs filled with EG solution, the calcination process was analyzed by TG-DSC curves in 2O4 at least 500 °C (2O4 NPs. Subsequently in 2O4.The obvious characteristic peaks are perfectly indexed to the nickel ferrite spinel phase (JCPDS card No. 54–0964), which confirms the formation of NiFet 500 °C . Additio2O4 (311) and (220) lattice planes, respectively .With the increasing of the calcination temperature, some nanocrystals aggregate destroying the orderliness in regions and when the crystallite size exceeds ca. 40 nm, the 3D network structure breaks down (21ectively . From thperature . The texaks down  This res311) and 0 lattice2O4 in 2O4 samples. Both the PMMA microspheres and meshes of 3D network show a FCC structure in three-dimensional space  is calculated by using equations (1)9To further investigate the microwave absorbing properties of the 3D network structure NiFein is the input impedance of the absorber and the normalized input Zin was calculated by equations (2).In which Zεr and μr is the complex relative permittivity and permeability, respectively, f is the frequency, d is the thickness of the absorption layer and c is the propagation velocity of the electromagnetic wave in vacuum. Synchronously, 2O4 with a thickness of 4.0 mm, which is stronger and wider than the ordinary NiFe2O4 NPs. For better comparison, the minimum reflection loss and the absorption bandwidth of some microwave absorbing materials are listed in 2O4 (RL = −27.5 dB) and the absorption bandwidth (4 dB) for this work are better than those of the similar microwave absorbing materials which implies that the as-prepared 3D network structure NiFe2O4 can be potentially used in the microwave absorbing field.Where 2O4 show a greater reflection loss with the same thickness comparing to the ordinary NiFe2O4 nanoparticles. In other words, to meet the same performance, less 3D network structure NiFe2O4 will be required in the applications, which demonstrates its feature of lightweight. Most often, reflection loss less than 10, 20 decibel (dB) is equivalent to 90% and 99% attenuation, respectively. The multiple valleys can be attributed to the special 3D network structure and the reasons are introduced in following discussions.Moreover, the 3D network structure NiFeεr) and complex permeability (μr) of ordinary NiFe2O4 NPs  represent the storage capability of electric and magnetic energy, whereas the imaginary parts (ε″ and μ″) stand for the inner dissipation abilities, respectivelyμ″) exhibits one major peak at around 5 GHz and a minor peak at f = 13 GHz, however the corresponding real part (μ′) decreases in the same region. By contrast, both ε′ and μ′ of the NiFe2O4 NPs is larger than the 3D network structure NiFe2O4 on value, but ε″ and μ″ of the NiFe2O4 NPs and the 3D network structure NiFe2O4 are equal. 2O4 NPs and the 3D network structure NiFe2O4 samples are equipped with the weak magnetic loss.It is well known that the microwave absorption property is mainly determined by the electromagnetic parameter, the impedance match and the microstructure. As a consequence, the frequency dependence complex permittivity  and electric conductivity (σ), which can be expressed by the following equationThe magnetic properties of the 3D network structure NiFeis loops . The satμ0 is the permeability of vacuum. As a deformation formula:In Eq. (3), μ″(μ′)−2f−1 should be constant when the frequency is changed. The C0 value is shown in f > 13.0 GHz, the section is approximately constant, which is ascribed to eddy current effect. That may issue from the agglomeration of small NiFe2O4 nanocrystals.According to the skin-effect criterion, if the magnetic loss only originates from the eddy current loss, the values of 2O4 due to the small NiFe2O4 nanoparticles, which can be considered as dipoles. Additionally, numerous voids existing in the 3D network structure NiFe2O4 samples cause multiple reflections and scatter in the material, at the meantime, provide more interfaces between paraffin, NiFe2O4 nanocrystallines and air, which induce interactions of dipoles and interfacial polarization which contribute to the dielectric loss significantlyIn addition to the magnetic loss, the dielectric loss is crucially determined by the configuration of the 3D network structure. It is easy to understand that dipolar polarization takes place in 3D network structure NiFe2O4 are constituted by nanocrystalline domains possessing different orientations, it could be proposed that the energy levels of 3D network structure NiFe2O4 are discontinuous according to quantum effectsMultiple mechanisms work together changing the position of reflection loss peak as shown in 2O4 are concluded as follows: firstly, magnetic loss was attributed to the natural resonance mechanism (intrinsic magnetic loss) of NiFe2O4. Secondly, dipole polarizations and interfacial polarization were presented due to the interfaces between NiFe2O4 nanoparticles, air and paraffin wax. The complementarities between multiple mechanisms also played an important role in improving the microwave absorption properties. In consequence, 3D network structure NiFe2O4 studied here are promising candidates for application as microwave absorber and these 3D network structure can also be used not limited to ferrites.The reasons for the excellent microwave absorption properties of 3D network structure NiFe2O4 was successfully synthesized by using PMMA colloidal crystal templates for its better wettability with polar solvents. From the thermal behavior of PMMA/NiFe2O4 composites, the optimum temperature to remove CCTs was determined no less than 400 °C. Anyhow, excessively high temperature must be avoided because nanocrystalline with size more than 40 nm will make the 3D network structure collapse. As-prepared NiFe2O4 samples presented a perfect 3D network structure that built by nanocrystalline and the interconnection meshes involved many layers. Intrinsic magnetic loss of NiFe2O4, interfacial polarization and the multi-reflection caused by 3D network structure made the sample show an excellent microwave absorbing properties and the reflection loss curves had three valleys. The minimum reflection loss is −27.5 dB with a thickness of 4.0 mm. Therefore, the prepared 3D network structure NiFe2O4 has potential applications in microwave absorbing field.3D network structure NiFeThe reagents mentioned were all of analytical grade and used without further purification except methyl methacrylate (MMA). The MMA monomer was previously distilled under the reduced pressure to remove the inhibitor. All the water of the experiments mentioned before was deionized.2 gas atmosphere protection. After 30 min, 0.1 g KPS of initiator was added. Five hours later, monodisperse PMMA spheres were obtained. The resultant PMMA microspheres were rinsed for several times with the deionized water and close-packed into colloidal crystals by the vertical deposition before use.Non-cross-linked, monodisperse PMMA spheres with various diameters ranged from 100 to 350 nm were synthesized using emulsifier-free emulsion polymerization according to literatures2O4 could be described as a templated salt precipitation method. In detail, Ni(NO3)2·6H2O and Fe(NO3)3·9H2O were dissolved in ethylene glycol with a cation ratio of 1:2 by continuously stirring at least 8 h, given the viscidity of ethylene glycol. And the total ion concentration was 1.8 M. Then, the pre-prepared PMMA CCTs were soaked in precursor solution completely for 2 h with the subsequent vacuum filtration operation. After an overnight drying step, NiFe2O4 with the 3D network structure was prepared by calcining at a template-removing appropriate temperature for 2 h in the air. And a slow heating rate  was imperative to keep the structure as ordered as possible, especially at the stages when the polymer template gasifies or combusts.The preparation strategy of 3D network NiFe2 and O2. The phase composition and purity were analyzed by X-ray diffractometer  using Cu Kα radiation of wavelength λ = 0.15418 nm. Besides, magnetization measurements were performed on a vibrating sample magnetometer  and the magnetic hysteresis loop measurements were performed at 8 kOe at 300 K. The microwave absorbing properties of the material was estimated based on the coaxial transmission/reflection method. The electromagnetic parameters of the 3D network structure NiFe2O4 samples were recorded by vector network analyzer (VNA) with the samples uniformly mixed with paraffin wax in a mass ratio of 7:3, forming ferrite/paraffin composites  and then pressed into a toroidal shape at a preesure of 0.6 MPa . The total mass of sample used in test is a fixed value of 0.13 g.The morphologies, geometric parameters and lattice fringes were characterized by field emission scanning electron microscope  and transmission electron microscope . The thermogravimetry (TG) and differential scanning calorimetry (DSC) were recorded on a thermal analyzer  with a heating rate of 10 °C/min in a flowing mixed gas of NHow to cite this article: Wang, F. et al. Lightweight NiFe2O4 with controllable 3D network structure and enhanced microwave absorbing properties. Sci. Rep.6, 37892; doi: 10.1038/srep37892 (2016).Publisher's note: Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.