Because the escalating population and energy production have been speeding up the combustion of fossil fuels 1 around the world, severe environmental issues are gradually appearing. 2 Therefore, the energy storage market is ourishing in terms of energy-saving technology and technical development. 3,4 Furthermore, it is signicant that the obtained energy should be produced from renewable energy resources and be stored [5][6][7] in effective ways. Environmentally friendly energy storage and conversion based on the electrochemical technology [8][9][10][11][12][13] can occupy a position of dominance in the energy market, 14 because it can release/store energy with a considerable conversion efficiency. [15][16][17][18][19][20] Lithium-ion batteries (LIBs) have been widely used in today's society because of their large specic power and high energy density. [21][22][23][24] The widespread use of LIBs in power portable electronic devices and grid-scale energy storage could be impractical, however, because they are limited by their high cost 25,26 and ammable electrolyte. Recently reported aqueous metal ion (monovalent-ion, bivalent-ion, multivalent-ion, etc.) batteries have garnered much attention owing to their highenergy-density, high-safety, and nontoxicity. Rechargeable aqueous zinc-ion batteries (ZIBs) [27][28][29][30] composed of an intercalation/deintercalation cathode, aqueous electrolyte, and zinc metal anode have become one of the most popular research hotspots in the energy and resources sector. Several advantages of ZIBs have already been demonstrated compared to LIBs and other nonaqueous metal ion batteries. First, zinc offers important advantages over lithium due to its natural abundance and low cost, as well as its high theoretical specic capacity (820 mA h g À1 ) and low working potential (À0.78 V vs. the standard hydrogen electrode potential). Second, the processes of zinc deposition and stripping are also facile in an aqueous electrolyte because of its high overpotential. Third, the fewer safety concerns associated with ZIBs could make it possible to scale them up, due to their incombustibility and high volumetric capacity (5854 mA h cm À3 ). 31 Many cathode materials for ZIBs have been reported so far, such as LiV 3 O 8 , 32 Na 3 V 2 (PO 4 ) 3 , 30 MnO 2 , 27 MnS, 33 ZnMn 2 O 4 , 34 Zn 3 V 2 O 7 (OH) 2 $2H 2 -O, 31 Prussian blue, 35 etc. Despite these efforts, further scientic research on such excellent cathode materials with a high capacity and remarkable cycling performance is essential, bearing in mind that ZIBs are just in the developmental stage.
Recalling the whole R&D process for Li-insertion cathodes, most previous studies have focused on the AMn 2 O 4 (A ¼ Zn, Co, Ni, etc.) spinel structure, which can improve the electrochemical performance compared to monovalent alternatives as well as the cycling stability 36,37 because of the multiple Mn valences. Therefore, it is meaningful to explore these structures for Znstorage. For example, the spinel structure of ZnMn 2 O 4 (ref. 34) is maintained during the intercalation/de-intercalation process. Zn(H 2 O) n 2+ (zinc hydrated ions) cannot intercalate into the spinel structure. However, H + can insert into the layered structural materials (such as MnO 2 and NH 4 V 4 O 10 ), and OH À can react with ZnSO 4 and H 2 O in the electrolyte to produce some large ake-like ZnSO 4 [Zn(OH) 2 ] 3 $xH 2 O. 27,31 Therefore, this phenomenon could increase the pH of the electrolyte. It can be concluded that the structure of the spinel systems is much more stable than the layered materials during cycling. These Mnbased materials have common barriers to wider application, however, such as large volume variation and poor electrical conductivity during the phase transitions in the course of cycling. Among these systems (Mn-based electrodes) researched for energy storage, NiMn 2 O 4 , as one of the most interesting spinel oxides, was widely used in lithium-ion batteries (LIBs) and supercapacitors (SCs) with superior electrochemical activity and structural reversibility 38,39 because of its synergetic Ni-Mn structure. Very recently, cation-site doping has been widely used to enhance its structural stability and intrinsic conductivity. NiMn 2 O 4 can provide enough redox active sites 40 for critical electrochemical reactions because of the lattice competition between nickel ions and manganese ions in the NiMn 2 O 4 spinel metal oxide. 41 Furthermore, in contradistinction to other transition metal manganites (AMn 2 O 4 ; A ¼ Mg, Co, Fe, Mn, and Zn), the crystalline structure of the Ni-Mn spinel is a symmetric cubic structure, which shows tetragonal distortion because of the dynamic Jahn-Teller effect 42 related to the existence of trivalent manganese in the octahedral sites. The tetrahedral lattice sites (A-sites) and the octahedral interstices (B-sites) are, respectively, occupied by divalent cations and trivalent cations for a regular spinel. [43][44][45] Nevertheless, NiMn 2 O 4 is not the usual type of spinel but a hybrid spinel, where a fraction y of the divalent nickel ions could leave the A-sites and likely be located on the B-sites in the crystal, which is characterized by cubic close packing of oxygen atoms. 42,[46][47][48] This will result in a corresponding 2y of the trivalent manganese ions that is disproportional to divalent manganese ions and tetravalent manganese ions in these B-sites, so that, simultaneously, the corresponding divalent manganese ions could migrate to the A-sites to balance the valence change induced by the transfer of divalent nickel ions. 47 This may result in capacity fading during cycling owing to the dissolution of Mn 2+ for Mn-based materials. A Ni-doped electrode material with Mn 4+ -rich phases can provide a much more exible response to the active sites, however, optimizing the cycling stability and suppressing the destruction of the crystal structure. With these considerations from the above analysis in mind, new highly porous, narrow band-gap, and Mn 4+ -rich Ni x -Mn 3Àx O 4 (x ¼ 0, 0.55, 0.8, 1.0, and 1.2) nano-particles were synthesized in a carbon matrix formed in situ. The details are included in the Experimental section in the ESI. † We report highly reversible Zn 2+ de/intercalation and stable performance for the Zn/Ni x Mn 3Àx O 4 @C system for the rst time. The unique idea for the design could be principally demonstrated as follows: (I) Ni-doping endows Ni x Mn 3Àx O 4 with a low band-gap, improving its electrical conductivity; (II) Ni-doped Mn 4+ -rich Ni x Mn 3Àx O 4 can provide much more exible active sites to preserve the crystal structure; (III) the pores inside Ni x Mn 3Àx O 4 nanoparticles facilitate fast mass/ion transport and penetration in the mild electrolyte; and (IV) Ni x Mn 3Àx O 4 in the Ni x Mn 3Àx -O 4 @C electrode could anchor the coated carbon strongly and conformally, restraining the aggregation and dissolution of Ni x Mn 3Àx O 4 . Based on these synergic effects, the Zn/Ni x -Mn 3Àx O 4 @C (x ¼ 1) system can work well in an aqueous electrolyte (2 M ZnSO 4 + 0.15 M MnSO 4 ), and demonstrate a high specic capacity (133.7 mA h g À1 at 200 mA g À1 ) and long-term cyclability (91.9% capacity retention aer 850 cycles at 400 mA g À1 ) in the range of 1.0-1.85 V. Its electrochemical properties can be improved by controlling the molar ratio of Ni and Mn, owing to the displacement of Ni 2+ from B sites, which can change the electrical properties introduced by the "hopping mechanism" of Mn 3+ and Mn 4+ on the B-sites via localized states. [48][49][50] Moreover, the mechanism of Zn 2+ de/intercalation in the nonstoichiometric Ni x Mn 3Àx O 4 spinel coated in situ with amorphous carbon layers has been deeply studied by ex situ Xray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), scanning electron microscopy (SEM), energy-dispersive spectrometry (EDS), transmission electron microscopy (TEM), rstprinciples calculations, and nitrogen adsorption/desorption experiments.
As shown in Fig. 1a, the synthesis of Ni x Mn 3Àx O 4 @C samples with different Ni : Mn ratios was conrmed by powder XRD analysis. The obtained XRD data as a function of the Ni proportion in a series of prepared samples correspond well to the NiMn 2 O 4 spinel structure (Fd3m), which indicates that the calcination of precursors did not change the crystal structure of Ni x Mn 3Àx O 4 @C. For these Ni x Mn 3Àx O 4 @C hybrids, no NiO impurity was detectable except for the Ni x Mn 3Àx O 4 @C (x ¼ 1.2) composition, demonstrating the successful transformation of Ni x Mn 3Àx O 4 @C (x ¼ 0.55, 0.8, and 1), which was understood by referring to the JCPDS standard. Interestingly, when the Ni content was increased from 0 to 0.55, 0.85 and 1, respectively, Ni x Mn 3Àx O 4 @C (x ¼ 0) with the tetragonal Mn 3 O 4 (JCPDS, 24-0734, I4 1 /amd) structure (as shown in Fig. S1 in the ESI †) was gradually transformed into Ni x Mn 3Àx O 4 @C (x ¼ 1) with the space group Fd3m. Compared with Ni x Mn 3Àx O 4 @C (x ¼ 0.55, 0.8), the strong peaks of the Ni x Mn 3Àx O 4 @C (x ¼ 1) composition at 2q ¼ 18. 3 , 30.15 , 34.95 , 36.86 , 43.52 , 53.54 , 57.90 , 62.95 , 73.99 , and 75.37 correspond to the (111), (220), (311), (222), (400), (442), (511), (440), (533), and (622) reections in the XRD standard of NiMn 2 O 4 (JCPDS, 1-1110, Fd3m). It has been found, however, that the peaks of NiO (JCPDS, 47-1049, Fm 3m (225)) were only observed in the Ni x Mn 3Àx O 4 @C (x ¼ 1.2) sample, which are labeled with blue stars. This is due to the fact that further increase of the Ni content to 1.2 could lead to the segregation of multiphase materials including NiO and Ni 2 MnO 4 , which indicates that the substitution limit of Ni for Mn in Ni x Mn 3Àx O 4 @C (x ¼ 1.2) has been reached. We nd that nearly all of the peaks of the samples are shied to lower 2q values, especially between the dashed-dotted lines, denoting a larger lattice constant with increasing Ni content. The peaks in Fig. 1a are known to shi to lower 2q values with the substitution of large radius foreign ions. 51 The ionic radius of Ni 2+ (69 pm) is larger than that of Mn 2+ (60 pm), Mn 3+ (58 pm), and Mn 4+ (53 pm), consistent with this trend.
To discern the positions of Ni and Mn ions in the lattice interstices, the intensity ratio R ¼ I 400 /I 220 can be used to study the distribution between them in AB 2 O 4 spinels, 44 and it has been reported that the value of R could change if an increasing proportion of foreign ions is substituted at the A-sites. Furthermore, the XRD pattern of Ni x Mn 3Àx O 4 @C (x ¼ 0) was also obtained, as shown in Fig. S1 †, as a comparison. In the case of Ni x Mn 3Àx O 4 @C (x ¼ 0), the obtained R is 1.237. The values of R measured for Ni x Mn 3Àx O 4 @C (x ¼ 0.55, 0.8, and 1) are 1.228, 1.230 and 1.235, respectively, which indicates that an increasing proportion of Ni 2+ is likely to be occupying B-sites, substitutes that prefer to occupy Mn-sites instead. 47 This will be further studied by means of density functional theory (DFT). In Fig. S1, † however, all characteristic peaks are indexed to the tetragonal spinel structure of Mn 3 O 4 with the space group I4 1 /amd. In this structure, Mn 2+ preferentially occupies the A-sites, while Mn 3+ occupies the B-sites. It is interesting to note that the valence of Mn in Mn 3 O 4 may be not the same as that exhibited in Ni x -Mn 3Àx O 4 @C. The substitution of Ni at the B-sites of Mn 3 O 4 could change the valence state of Mn owing to the electroneutrality principle. 47 A similar phenomenon has been reported in a previous study (ref. 52) for substituted LiM x Mn 2Àx O 4 (M ¼ Cr, Li, Ni, etc.) spinels.
More importantly, Ni cations in spinels enhanced the cycling capacities 21 of Li(Ni y Mn z Co 1ÀxÀy )O 2 cathodes for LIBs, and Mn cations stabilized the spinel structure aer partial substitution by Ni cations at B-sites. It is noteworthy that cations with a larger ionic radius and lower valence can balance the crystal stress and stabilize the spinel structure if they coexist with other cations with a smaller radius and higher valence state at the Bsites. For example, cooperative synergies between nickel cations (+2) and manganese cations (+4) made it possible to increase the stability 53 of the framework of LiNi 1/3 Co 1/3 Mn 1/3 during charge-discharge cycling.
XPS measurements were carried out to evaluate the oxidation states of Mn and Ni cations in Ni x Mn 3Àx O 4 and Ni x Mn 3Àx O 4 @C (x ¼ 1), and the corresponding survey spectra are shown in Fig. S2 and1b. † Fig. 1b demonstrates the existence of signals associated with Ni, Mn, O, and C. The Ni : Mn ratio (1 : 2.002) was also obtained according to the integrated intensity of the corresponding peaks, which was corrected by elemental sensitivity factors. The result from XPS may not be completely dependable, because it is a semi-quantitative method. The obtained Ni/Mn atomic ratio, which is close to the designed value for the experiment, indicates, however, that Ni and Mn cations are homogeneously distributed within the limits of the XPS analytical depth. Ni 2p and Mn 2p 3/2 photoemission spectra are shown in Fig. 1c andd, respectively. In the case of Ni 2p, the spectrum shows the presence of two main peaks at 854.2 eV (2p 3/2 ) and 871.9 eV (2p 1/2 ) along with two satellite peaks at 860.8 eV and 878.8 eV. Interestingly, the peak distance between the satellites and the main peaks is about 7 eV, whereas the main peak difference is 17.7 eV, similar to the value of 17.6 eV observed for Ni 2+ (ref. 54) in NiO. The large full width at half maximum (FWHM) for Mn 2p 3/2 signals of Ni x Mn 3Àx O 4 @C (x ¼ 1) is indicative of multiplet splitting along with the charging effect, 46 owing to the presence of Mn 4+ , Mn 3+ and Mn 2+ ions. The Mn 2p 3/2 spectrum of Ni x Mn 3Àx O 4 @C (x ¼ 1) composites is a superimposition of Mn 4+ , Mn 3+ , and Mn 2+ peaks, as shown in Fig. 1d, demonstrating that Mn exists in three different oxidation states. These peaks correspond to a binding energy of 643.5 eV, 642.4 eV, and 641.5 eV, 46 respectively. It was observed that the peak area of Mn 4+ is the largest among them from the results of XPS investigations of the Ni x Mn 3Àx O 4 @C (x ¼ 1) compound (Table S1, ESI †). Why do we choose it as the example for further studies? Ni x Mn 3Àx O 4 @C (x ¼ 1) exhibits a ne and stable crystal structure according to the physical characterization, shows a highly Mn 4+ -rich structure, and displays a good morphology with uniform nanoparticles, as proven below. Furthermore, the as-prepared Ni x Mn 3Àx O 4 @C (x ¼ 1) has the highest content of Mn 4+ (as shown in Fig. S3 †) in its crystal structure aer the precursor is calcined at the target temperature of 455 C for 2.5 h in an oxygen atmosphere. The O 1s spectrum of Ni x Mn 3Àx O 4 @C (x ¼ 1) is deconvoluted into three peaks as shown in Fig. 1e, where the peak at 529.9 eV (O1) is due to the metal-oxygen bonds, whereas a weaker peak at 531.5 eV (O2) is correlated with the large amount of -OH species adsorbed on the surface. The smallest peak at 533.2 eV (O3) could be due to chemisorbed oxygen or adsorbed water on the surface of the sample. The core-level C 1s XPS spectrum of Ni x Mn 3Àx O 4 @C (x ¼ 1) consists of three components centered at 284.8 eV, 286.1 eV and 288.3 eV, corresponding to the C-C, C-O, and C]O bonds, respectively (Fig. 1f).
From these XPS results, the valence of Mn ions could change 46,47 during Ni substitution for Mn in the Ni x Mn 3Àx O 4 structure, which is consistent with the previous discussion. Additionally, Mn 4+ displays a higher valence than Mn 3+ but a smaller ionic radius, while Ni 2+ exhibits a lower valence than Mn 3+ but a larger ionic radius. The substitution of Ni species 9 in the crystal structure has been proven to boost the reversible capacity, while Mn species ameliorate the interfacial impedance 1 and enhance the electrochemical performance. Therefore, the synergies of Mn 4+ -Ni 2+ and their substitution at the B-sites balance the valence charge in the crystal lattices, which relieves the crystal stress owing to the mismatching ionic radius and hence promotes the stability of the spinel structure. Furthermore, the Mn 4+ rich phase in Ni x Mn 3Àx O 4 @C (x ¼ 1) can act as protective component to improve the medium to promote zinc ion diffusion and alleviate Mn 2+ dissolution during chargedischarge. According to the above-mentioned research results, it has been observed that the intensive synergic action between Ni 2+ and Mn 4+ at B-sites could be used as the evidence for the reversibility and stability of the active materials during charge/ discharge. Furthermore, more details on the changes to the microstructure and the lattice defects of the obtained samples were also revealed by Raman spectroscopy (Fig. S4 †) and Fourier transform infrared (FTIR) spectroscopy (in Fig. S5 †), which are complementary and provide valuable information on the chemical bonds in Ni-Mn-O systems in the hybrids.
It can be seen from the Raman spectra (in Fig. S4 †) of Ni x -Mn 3Àx O 4 @C hybrids that the intense vibrational modes located around 654 cm À1 and 556 cm À1 are assigned as the A 1g and F 2g modes, respectively. The peak located at 556 cm À1 reects the symmetric Ni-O stretching vibration derived from the F 2g mode 55 in NiMn 2 O 4 . The A 1g active mode originates from the symmetric Mn 4+ -O (MnO 6 ) stretching vibration at octahedral sites. 56 Interestingly, the Mn 2+ -O and Mn 4+ -O vibrations at Asites and B-sites are in accordance with the results obtained from the above XRD and XPS analysis. Additionally, the T 2g (located at 305 cm À1 ) and E g (located at 372 cm À1 ) modes gradually disappear as x increases in the Ni x Mn 3Àx O 4 @C compounds, and the F 2g and Ni-O (locates at 509 cm À1 ) stretching modes appear. These results indicate that Mn 3+ -O is transformed into Mn 4+ -O at the B-sites during Ni substitution for Mn in the Ni x Mn 3Àx O 4 structure. Typical carbon peaks located at 1362 cm À1 (D band) and 1580 cm À1 (G band) are detected from the samples. Furthermore, Fourier transform infrared (FTIR) spectroscopy is a complementary characterization method, which provides valuable information of Ni-Mn-O systems. The absorption peaks at 597 cm À1 and 507 cm À1 are assigned to Mn-O 4 and Mn-O 6 (ref. 56) vibrations (Fig. S5 †), respectively, and the lower absorption bands at 452 cm À1 can be assigned to Ni-O 6 , indicating that Ni can substitute Mn at the Bsites along with the change in the valence of Mn 3+ at the B-sites. It should be noted that Mn 4+ and Ni 2+ at the B-sites take the place of Mn 3+ in the compounds, as evidenced by XRD, XPS, Raman spectroscopy, and FTIR spectroscopy analyses.
The overall typical nanoparticle morphologies of the asprepared Ni x Mn 3Àx O 4 @C compounds and their crystal phases were characterized by using SEM, selected-area electron diffraction (SAED) analysis, energy dispersive X-ray spectroscopy (EDX) elemental mapping, and high resolution TEM (HRTEM) (as shown in Fig. 2 and S6 †). The samples show an irregular particle morphology with a size range of 15-25 nm (Fig. 2a-c), and the morphology is transformed into more irregular aggregated rectangular bulks (Fig. 2d) on a large scale (x ¼ 1.20), indicating that an insufficient or excess content of Ni in the Ni x Mn 3Àx O 4 @C compounds can affect the morphology of the nanoparticles. Therefore, high-resolution TEM and lowresolution TEM images (Fig. 2g and its inset) of Ni x Mn 3Àx -O 4 @C (x ¼ 1.2) reveal various aggregates with mixed phases. The lattice spacings measured in Fig. 2g shows interplanar spacings of 0.241 nm and 0.207 nm corresponding to the (1 1 1) and (2 0 0) planes of crystalline NiO, which are consistent with the XRD (Fig. 1) and SAED (Fig. S6b †) results. Furthermore, the EDX spectrum (Fig. S6a †) demonstrates that the ratio of Ni : Mn is 1 : 2.026 in Ni x Mn 3Àx O 4 @C (x ¼ 1), whereas the ratio of Ni : Mn is 1 : 2.002 according to inductively coupled plasmaatomic emission spectrometry (ICP-AES) results. The differences between these two values and the theoretical ratio are quite small, further demonstrating the successful substitution of Mn by Ni. Elemental mapping images (Fig. 2i-i 4 ) of C, Ni, Mn, and O manifest no distinct phase segregation, revealing a homogeneous distribution of Ni and Mn elements in the Ni x Mn 3Àx O 4 @C (x ¼ 1) hybrid. The diffraction rings obtained from the SAED (inset of Fig. 2e) characterization indicate the polycrystalline nature of this sample. The HRTEM lattice image of this compound is also shown in Fig. 2f, which displays well-dened crystalline planes with d-spacings of 0.483 nm and 0.296 nm, corresponding to the (1 1 1) and (2 2 0) lattice planes of the face-centred cubic spinel (Fd3m) phase, respectively. As seen in Fig. 2e andf, there is a thin amorphous carbon layer ($2 nm) coated on the surface of Ni x Mn 3Àx O 4 @C (x ¼ 1), which can be conducive to restraining the aggregation and volume changes of Ni x Mn 3Àx O 4 @C (x ¼ 1) nanocrystals during charge/ discharge and improving the conductivity of the electrode. The weight percentage of carbon (9.8 wt%) was determined by thermogravimetric (TG) analysis, as shown in Fig. S7. † Additionally, the mesoporous structure of the Ni x Mn 3Àx O 4 @C (x ¼ 1) compound is conrmed by an obvious hysteresis loop (Fig. S6c †) in the N 2 adsorption-desorption isotherms, proving a large surface area of 75.275 m 2 g À1 and a wide pore-size distribution at $4.76 nm (the pore volume is calculated to be $0.39 cm 3 g À1 ).
We investigated the electrochemical properties of the Ni x -Mn 3Àx O 4 @C samples using typical coin cells, employing a Zn foil anode, glass ber separator, and aqueous electrolyte (2 M ZnSO 4 + 0.15 M MnSO 4 ), and a schematic illustration of the ZIB is presented in Fig. 3a. Enlarged views of the cathode and the interior of Ni x Mn 3Àx O 4 @C (x ¼ 1) are shown in Fig. 3b andc, respectively. Ni x Mn 3Àx O 4 inside the nano-particles generates numerous narrow pores so that zinc ions could transfer more efficiently through these tunnels (Fig. 3d). The uniquely designed Ni x Mn 3Àx O 4 @C (x ¼ 1) with its spinel structure and conductive carbon network provides superior electrochemical performance. Fig. 3e displays the possible reaction mechanism of zinc ion intercalation/deintercalation in Ni x Mn 3Àx O 4 @C (x ¼ 1). Fig. 4a shows the cyclic voltammetry (CV) curves of Ni x -Mn 3Àx O 4 @C (x ¼ 1) measured between 1.0 and 1.85 V within the rst eleven cycles at a scan rate of 0.3 mV s À1 . The proles in the rst and third cycles are slightly shied to higher potentials in comparison with the following cycles, which can be related to the initial activation of the electrode. During cycling, the polarization of the electrode decreases from the third cycle, as demonstrated (in Fig. S8 †) by electrochemical impedance spectroscopy (EIS) analysis. The CV curves (Fig. 4a) prove that there is only one pair of redox peaks for Ni x Mn 3Àx O 4 @C (x ¼ 1) at about 1.34/1.52 V during subsequent cycles, which can be attributed to Zn 2+ deintercalation/intercalation from/into the host, consistent with the voltage plateaus shown in the charge and discharge curves (Fig. 4b). The rst discharge plateau at about 1.30 V corresponds to Zn 2+ intercalation into the cathode, while the rst charge plateau at about 1.55 V originates from Zn 2+ extraction. The pair of redox peaks of Ni x Mn 3Àx O 4 @C (x ¼ 1) at 1.34/1.52 V could be attributed to the reversible conversion of Mn 3+ /Mn 4+ , which can be attributed to the corresponding phase transition during charge/discharge. The charge storage kinetics and phase transition in the Ni x Mn 3Àx O 4 @C (x ¼ 1) host will be discussed below.
As expected, the possible cathodic process may proceed as follows:
Furthermore, the CV data indicate little change in the potential or peak area loss aer the third cycle for all the redox features, proving that Ni x Mn 3Àx O 4 @C (x ¼ 1) can be an appropriate cathode candidate for aqueous ZIBs.
Fig. 4b presents the discharge/charge curves of the Ni x -Mn 3Àx O 4 @C (x ¼ 1) electrode cycled at a current density of 150 mA g À1 within the voltage window of 1.0-1.85 V. The active material delivers initial specic discharge and charge capacities of 115 mA h g À1 and 119 mA h g À1 , respectively. Aer 100 cycles, the electrode sustains a stable capacity of around 130 mA h g À1 with a high coulombic efficiency (CE) of around 100%. The average discharge plateau is 1.37 V during cycling, giving an excellent energy density of $178.1 W h kg À1 according to the active mass. The Ni x Mn 3Àx O 4 @C (x ¼ 1) nanoparticles can provide a calculated energy density of 67 W h kg À1 under the condition that the content of cathode materials accounts for one third of the whole weight of the battery, which is much higher than that for Ni-Cd (about 50 W h kg À1 ) and Pb-acid (about 30 W h kg À1 ) batteries. In comparison, a Ni x Mn 3Àx O 4 + C (x ¼ 1) mixture (9.8 wt% carbon) displays an obviously lower capacity of 44 mA h g À1 at 100 m A g À1 aer 100 cycles (Fig. S9 †). To evaluate the capacity from the carbon layer coated in situ on the surface of Ni x Mn 3Àx O 4 @C (x ¼ 1), galvanostatic charge/ discharge curves (Fig. S10 †) of the carbon matrix synthesized by the same method were collected between 1.0 and 1.85 V at 15 mA g À1 and indicate that the obtained capacity ($0.5 mA h g À1 ) can be neglected. Fig. 4d presents the cycling performance of the obtained samples at 250 mA g À1 over 400 cycles. All the samples clearly display a gradual capacity increase before 53 cycles owing to the process of electrochemical activation. The capacity of Ni x Mn 3Àx O 4 @C (x ¼ 0 and 0.55) began to drop quickly aer electrochemical activation, however, and its capacity retention ratios are 45% and 48% aer 400 cycles, respectively, representing bad electrochemical performance. The Ni x Mn 3Àx O 4 @C (x ¼ 1, 1.2) composites present better cycling stability than other samples owing to the sufficient substitution of Ni, which also can affect the crystalline structure of Ni x Mn 3Àx O 4 , as discussed in Section of structure characterization. The discharge capacity of Ni x Mn 3Àx O 4 @C (x ¼ 1) is the highest (130 mA h g À1 aer 400 cycles) among them with a shortened activation process, indicating enhanced Zn 2+ deintercalation/intercalation kinetics. This excellent cycling performance could be principally due to several factors. First, the characteristic inverse spinel structure with abundant pores and Mn 4+ -rich crystalline phases could provide countless bonding sites and shorten the Zn 2+ ion diffusion paths. Second, the moderate amounts of Mn 2+ in the electrolyte efficiently suppress the structural collapse of Ni x Mn 3Àx O 4 , which can enhance the cycling stability and reduce the deterioration of the electrochemical performance (as shown in Fig. S11 †). Third, the homogeneous in situ carbon coating can enhance the ionic and electronic conductivities and increase cycling stability. The impressive zinc ion storage performance of Ni x Mn 3Àx O 4 @C (x ¼ 1) is further emphasized by its excellent rate capability. Fig. 4e presents the rate performance at different current densities, and corresponding cycle proles are shown in Fig. 4c. The cell displays a reversible discharge capacity of 130 mA h g À1 at a current density of 400 mA g À1 , which is close to the achieved discharge capacity (139 mA h g À1 ) at a lower current density of 50 mA g À1 . It should be noted that the obtained capacities from 50 mA g À1 to 100 mA g À1 are nearly invariable. This similarity in capacity demonstrate that 139.7 mA h g À1 is the limit for the deintercalation/intercalation of zinc ions from/into the Ni x -Mn 3Àx O 4 @C (x ¼ 1) hybrids. As the current density increases to 600 mA g À1 , 800 mA g À1 , 1000 mA g À1 , and 1200 mA g À1 , the reversible capacity progressively decreases to 90%, 85%, 79%, and 68%, respectively. The capacity fading (in Fig. 4e) is not irreversible when the cell is cycled at 1200 mA g À1 , because it can recover the previous capacity when subsequently cycled at a current density of 200 mA g À1 . The rate performance of Ni x -Mn 3Àx O 4 @C (x ¼ 1) is superior to that of the other samples, as conrmed in Fig. S12. † Moreover, the Ni x Mn 3Àx O 4 @C (x ¼ 1) electrode could operate for more than 200 cycles with a high reversible capacity (nearly $97% recovery of the capacity) of 131 mA h g À1 aer the rate test. The long cycling stability of Ni x Mn 3Àx O 4 @C (x ¼ 1) was further investigated at 400 mA g À1 over 850 cycles. As can be seen in Fig. 4f, the reversible capacity is 129.5 mA h g À1 aer dozens of cycles of activation and remains stable in the subsequent cycles, demonstrating superior cycling stability. Furthermore, the coulombic efficiency was increased signicantly to 97.8% aer about 35 cycles, and then it reaches almost 100% subsequently. The typical plateaus and shapes of the discharge/charge proles of Ni x Mn 3Àx O 4 @C (x ¼ 1) at 400 mA g À1 were also recorded aer different cycles (Fig. S13 †), further manifesting its superb cycling stability. Compared with the reported cathode materials (in Table S2 †) and other Mn-based materials (in Fig. S14 †), it shows remarkable cycling stability.
Structural studies and electrochemical investigations were used to further understand the deintercalation/intercalation mechanism and storage behavior of the Ni x Mn 3Àx O 4 @C electrodes. Ex situ XPS (Fig. S15a andb †) and XRD measurements (Fig. 5a andb) were carried out to analyze the mechanism of zinc ions deintercalation/intercalation from/into the Ni x Mn 3Àx O 4 @C (x ¼ 1) host. The ex situ XRD study was conducted in the range from 15 to 77 during the rst cycle in ZIBs. All the peaks could be indexed to the Ni x Mn 3Àx O 4 structure for the pristine electrode. The (1 1 1), (2 2 0), (5 1 1), and (4 4 0) peaks in the patterns display an obvious shi to lower 2q angles during discharge and return to their original positions in the following charge, which can be attributed to more zinc ions being extracted/inserted within the tested voltage range. Furthermore, some additional peaks, except for those of the current collector, emerge in the subsequent deintercalation/intercalation, in the potential region where the Ni x Mn 3Àx O 4 @C (x ¼ 1) host was intercalated by plenty of zinc ions. During the rst discharge platform above 1.35 V, only two sets of peaks exist, which are related to pristine Ni x Mn 3Àx O 4 @C (x ¼ 1) and the steel current collector. Subsequently, some new peaks rose gradually and the peaks of Ni x -Mn 3Àx O 4 @C (x ¼ 1) decreased during the following discharge process. The new peaks began to diminish and disappear little by little, while the peaks of Ni x Mn 3Àx O 4 @C (x ¼ 1) recovered to the original pattern during the subsequent charge process, demonstrating the excellent structural stability of the Ni x -Mn 3Àx O 4 @C (x ¼ 1) nano-composite. We note that a small amount of Mn ions shied to the A-sites from the B-sites, leading to some Mn-associated vacancies, which may induce some zinc ion diffusion along with the consequent structural transformation in the voltage range of about 1.4-1.0 V during discharge/charge. The crystal structure of the active materials can be preserved, however, because of the synergetic effects of Mn 2+ and Zn 2+ in the electrolyte, indicating the signicantly reversible behavior of Ni x Mn 3Àx O 4 @C (x ¼ 1). Good reversibility of zinc stripping/plating (Fig. S16a andb †) can be observed for the Zn 2+ electrolyte with the Mn 2+ additive, which showed a decrease of the overpotential during cycling. In addition, the peak positions for the intermediate transition phase (rich in zinc ions) can be indexed to the nickel zinc manganite spinel structure, (Zn 0.988 -Mn 0.012 )(Mn 1.326 Zn 0.024 Ni 0.65 )O 4 (JCPDS no. 04-01-070-6709). As shown in Fig. S15b, † the intensity of Zn 2p in the discharge states is far above that in the charge states, further manifesting the intercalation/deintercalation of zinc ions into/from the Mn 4+ -rich host. The zinc signal in Fig. S15b † did not disappear completely, which can be attributed to the residual Zn 2+ (on the surface of the host) from the electrolyte. Moreover, the CV curves at various scan rates aer activation at a scan rate of 0.2 mV s À1 for three cycles are presented in Fig. 5c. Obviously, the peaks gradually grow broad and high with increasing scan rate, but the shapes of the obtained CV curves remain consistent. As assumed, only one pair of redox peaks could be detected at a scan rate from 0.4 to 1.8 mV s À1 within the range of 1.0-1.85 V for each CV curve, which ought to be attributed to the Mn 4+ /Mn 3+ redox couple reactions, together with step by step Zn 2+ insertion/extraction into/from the tetrahedral sites of the Ni x Mn 3Àx O 4 @C (x ¼ 1) framework. XPS spectra of the O 1s and Mn 3s orbitals for Ni x Mn 3Àx O 4 @C (x ¼ 1) were recorded aer discharge/charge. As can be seen from Fig. S17a,† the O 1s XPS spectra for Ni x Mn 3Àx O 4 @C (x ¼ 1) are different between the charge state and the discharge state, while the positive shi of the stronger peak at 530.1 eV can be attributed to the intercalation of Zn 2+ into the sample during discharge. The Mn 3s spectra (Fig. S17b †) contain two peaks, which are separated by 5.2 eV aer discharge, 57 indicating the reduction of Mn 4+ to Mn 3+ . The Mn 3s spectra also contain two peaks that are separated by 4.7 eV aer charge, which reects Mn 3+ to Mn 4+ conversion (ref. 58) and the extraction of Zn 2+ from the sample. Therefore, it is reasonable to assume that the spinel host can support a zinc-rich and zinc-depleted state without obvious polarization during the electrochemical processes for the electrode between 1.0 and 1.85 V, indicating good reversibility and excellent cycling stability. To obtain further insight into the zinc-storage mechanisms of the electrode, kinetic analysis according to the scan rate of CV curves was carried out to research the contributions to the capacity of the diffusion-controlled and capacitive controlled processes. The relationships between the scan rate (v) and the measured peak current (i) are shown in Fig. 5d, which can be depicted as in the following equation:
where a and b are adjustable parameters, and the b-values can be obtained from the calculation of the slope of the plot (log i versus log v), which can be described as follows:
In particular, the b-value can vary between 0.5 and 1.0 under well-dened conditions. A b-value of 0.5 demonstrates a completely diffusion-dominated intercalation behavior, while a b-value of 1.0 indicates a typical capacitive process. The bvalues of 0.53 for anodic peaks and 0.55 for cathodic peaks can be matched at a scan rate from 0.4 to 1.8 mV s À1 , manifesting the kinetics of a diffusion-controlled process. In addition, the charge-transfer resistance ($109 U) of Ni x Mn 3Àx O 4 @C (x ¼ 1) is somewhat lower than for the other samples aer 100 cycles at 250 mA g À1 , as shown in Fig. S18, † indicating higher electrical conductivity. CV measurements using a three-electrode system were conducted. Notably, the aqueous electrolyte (2 M ZnSO 4 + 0.15 M MnSO 4 ) supports a wide electrochemical window as shown in Fig. S19, † and O 2 evolution was suppressed signicantly up to 2.15 V. In order to further prove the relationship between electrical conductivity and the amount of Ni substitution in the B-sites, the projected density of states (PDOS) of Ni x Mn 3Àx O 4 was calculated. Clearly, Fig. 5e , respectively. Very large supercells are required to make the numbers in these chemical formulas integers, which is computationally demanding and complicated. Therefore, three tuned chemical formulas were selected for simplication, which are NiMn 5 O 8 (x ¼ 0.5), Ni 2 Mn 4 O 8 (x ¼ 1.0), and Ni 3 Mn 3 O 8 (x ¼ 1.5). Note that all the Ni atoms occupy the B-sites. Therefore, the preference of Ni 2+ for B-sites is also proven by our computational results, where the B-site occupation in Ni x -Mn 3Àx O 4 is energetically more favorable (Fig. S20 †), which is similar to the results of our previous XRD analysis. The DOS demonstrates that pure Mn 3 O 4 is a semiconductor and that its band gap is about 1.20 eV (in Fig. 5e 1-1 ). Besides the metallic conductivity, NiMn 5 O 8 (x ¼ 0.5) and Ni 2 Mn 4 O 8 (x ¼ 1.5) are socalled half-metals, meaning that only spin up electrons are conductive, as shown in Fig. 5e 2-2 and e 4-4 . Whereas Ni 2 Mn 4 O 8 (x ¼ 1.0) differs slightly around the Fermi level, with the Fermi level crossing the spin-up and spin-down bands (in Fig. 5e 3-3 ), which suggests that ferrimagnetism (x ¼ 1) with a similar x value (around 1.0) could make this sample more conductive than the other samples.
In a typical synthesis, a mixture of Ni(Ac)$4H 2 O and Mn(Ac)$ 4H 2 O was dissolved in 95 mL of distilled water. The mixture was incubated in a water-bath at a constant temperature of 45 C, and aer magnetic stirring for 90 min, a light blue transparent liquid was obtained. Meanwhile, citric acid (CA) (0.25 M, 95 mL) was then added to the above mixed solution dropwise, which was generally controlled at about one drop per second. Aer 2 hours, the temperature was increased from 45 C to 85 C. Next, the temperature was maintained, and the mixture was stirred overnight until all the water in the container was evaporated. The wet sample was washed and ltered with acetone and ethanol. Then it was dried by using a supercritical dryer, which started at 15 C, with the temperature subsequently increased to 38 C while owing CO 2 was injected into the autoclave (corresponding pressure of about 20 MPa and lasting for 5 h). The as-synthesized sample was ground and further annealed with a heating rate of 8 C min À1 up to the target temperature of 455 C, which was maintained for 2.5 h in an oxygen atmosphere in a tube furnace. Finally, a calcination procedure is performed with a given heating rate of 5 C min À1 up to the target temperature of 650 C, which was maintained for 9 h in a tube furnace under an Ar atmosphere with a ow of 80 sccm. In order to change the proportion of nickel (x) in the obtained nanoparticles, different Ni : Mn ratios were designed, yielding porous spinel nanoparticles (Ni x Mn 3Àx O 4 @C) with x ¼ 0.55, 0.80, 1.00, and 1.20.
For comparison, Ni x Mn 3Àx O 4 @C (x ¼ 0) was obtained by a similar method to Ni x Mn 3Àx O 4 @C (x ¼ 1). In order to research whether Mn 4+ -rich Ni x Mn 3Àx O 4 @C exists in the compounds, Ni x Mn 3Àx O 4 @C (x ¼ 1) was selected as the study subject (the reason is given in the text). The precursor of Ni x Mn 3Àx O 4 @C (x ¼ 1) was annealed with a heating rate of 8 C min À1 at target temperatures of 365, 410, and 500 C, which were maintained for 2.5 h in an oxygen atmosphere in a tube furnace, and the ongoing processes are the same as in the above operations. All of the related chemical reagents are of analytical pure grade and used as received without further purication for all the experiments. Inductively coupled plasma-atomic emission spectroscopy (ICP-AES) was used to conrm the contents of Mn and Ni in the as-prepared samples.
In conclusion, highly porous Mn 4+ -rich Ni x Mn 3Àx O 4 @C (x ¼ 1) nano-particles wrapped with an in situ formed carbon matrix have been successfully fabricated by a facile and eco-friendly method, and the composite has been evaluated as a cathode material for an aqueous rechargeable ZIB. In this resulting composite, the substitution of Mn by Ni at the B-sites of Mn 3 O 4 was studied by ICP-AES, XPS, XRD, and DFT calculations, which demonstrate that cation-site Ni substitution combined with a high enough doping concentration can decrease the band gap and effectively improve the electronic conductivity in the Ni x -Mn 3Àx O 4 crystal structures. Moreover, the highly porous Mn 4+rich structure and amorphous carbon shell lead to fast electron transport and short Zn 2+ diffusion paths in a mild aqueous electrolyte. On the basis of the above unique structural characterization, the Ni x Mn 3Àx O 4 @C (x ¼ 1) hybrid displays excellent electrochemical properties with outstanding cycling stability, good rate capability and improved zinc storage performance in comparison to most of the reported cathode materials for ZIBs, with high discharge capacities of 139.7 mA h g À1 and 98.5 mA h g À1 at 50 and 1200 mA g À1 respectively, combined with a nearly 93.5% capacity retention at 400 mA g À1 aer 800 cycles. These exhibited results indicate that the as-synthesized Ni x Mn 3Àx O 4 @C (x ¼ 1) would be a prospective candidate as remarkable ZIB cathode materials. Furthermore, this work provides some ideas for not only achieving a good understanding of the intercalation process for divalent metal ions (Zn 2+ , Mg 2+ , and Ca 2+ ), but also facilitating the development of aqueous rechargeable alkali ion (K + , Na + , and Li + ) and their aqueous rechargeable hybrid batteries.
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