Synthesis of ultralong ZnFe2O4@polypyrrole nanowires with enhanced electrochemical Li-storage behaviors for lithium-ion batteries

Electrochimica Acta 306 (2019) 198e208

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Synthesis of ultralong ZnFe2O4@polypyrrole nanowires with enhanced electrochemical Li-storage behaviors for lithium-ion batteries Linrui Hou a, *, Ruiqi Bao b, Dienguila kionga Denis a, Xuan Sun a, Jinyang Zhang a, Fakhr uz Zaman a, Changzhou Yuan a, b, ** a b

School of Materials Science & Engineering, University of Jinan, Jinan, 250022, PR China School of Materials Science & Engineering, Anhui University of Technology, Ma'anshan, 243002, PR China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 5 March 2019 Received in revised form 16 March 2019 Accepted 18 March 2019 Available online 21 March 2019

Spinel ZnFe2O4 (ZFO) has been emerging as promising anode material for lithium-ion batteries (LIBs) due to its intrinsic merits. However, inferior electrical conductivity and severe volume change hinder its practical applications. Herein, we developed a bottom-up electrospinning and gas-phase polymerization strategy for purposeful fabrication of one-dimensional (1D) nano-architecture with continuous polypyrrole (PPY) nanoshell coating on each ZFO nano-subunit throughout the whole nanowires (NWs). Benefiting from unique compositional and structural advantages, the resultant 1D hybrid product (denoted as ZFO@PPY) with large electroactive sur-/interfaces and robust structural stability displayed superior Li-storage performance with high initial Coulombic efficiency, large reversible capacities, highrate capability and long-term cycling performance, when evaluated as a competitive anode for low-cost LIBs. Furthermore, In-situ X-ray diffraction measurement provided insights into the underlying Li-storage process of the ZFO@PPY NWs. More significantly, we strongly envision that our integrated electrode design would pave a promisingly feasible avenue for advanced hybrid anodes for advanced LIBs. © 2019 Elsevier Ltd. All rights reserved.

Keywords: Spinel ZnFe2O4 Ultralong nanowires Polypyrrole nanoshell Hybrid anode Li-ion batteries

1. Introduction The increasing energy demands and ongoing depletion of fossil fuels have motivated tremendous research endeavors in exploring advanced energy conversion and storage techniques. Lithium-ion batteries (LIBs) are becoming one of the most extensively used energy storage devices for electric vehicles (EVs), hybrid EVs, smart grids and portable electronic products [1]. The growing energy requirements drive great efforts in developing low-cost LIBs with high energy/power density and long cycle lifetime. Nevertheless, the common commercial graphite anode still suffers from a low theoretical capacity of ~372 mAh g
1, seriously limiting the development of the next-generation LIBs [2,3]. Therefore, the purposeful exploitation of appropriate anodes for advanced LIBs with excellent stability and even higher capacity is highly desired. Since the pioneering work reported by Poizot in 2000 [4],

* Corresponding author. ** Corresponding author. School of Materials Science & Engineering, University of Jinan, Jinan, 250022, PR China. E-mail addresses: [email protected] (L. Hou), [email protected], mse_ [email protected] (C. Yuan). https://doi.org/10.1016/j.electacta.2019.03.121 0013-4686/© 2019 Elsevier Ltd. All rights reserved.

transition-metal oxides (TMOs) with at least twice the capacities of common graphite-based anodes have been intensively investigated as attractive anodes for LIBs. Recently, binary spinel mixed TMOs such as AB2O4 (A ¼ Zn, Ni, Cu; B ¼ Co, Mn, Fe) have been emerging for high-performance LIBs as promising electrodes with superior electrochemical properties, originated from the synergistic effect from their complex components [5e11]. Particularly, the spinel ZnFe2O4 (ZFO) with bivalent Zn2þ and trivalent Fe3þ occupying the tetrahedral and octahedral sites stands out from other mixed spinel TMOs owing to the low cost, earth abundance, environmental benignity of the Zn and Fe species [12,13]. Besides, another intriguing features, i.e., low lithium extraction voltage (~1.5 V vs. Li/ Liþ) and large theoretical capacity of ~1001 mAh g
1 estimated from 9 Liþ ions per formula unit of ZFO, would be significantly conductive to boost the output voltage of a full battery with high energy density [12]. In 2004, NuLi et al. firstly reported the nanocrystalline ZFO and Ag-doped ZFO thin films as anodes for LIBs [14]. Afterwards, different ZFO anodes with various nanostructures were synthesized and investigated in detail as remarkable electrodes for LIBs application [15e20]. Although significant advances have been achieved so far, the ultimate performance of ZFO anodes still cannot meet the commercial demands, owing to its congenital

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disadvantages including poor electrical conductivity, large volume change, and particle aggregation during the charge/discharge process, always resulting in inferior rate capability and fast capacity fading [15]. To effectively overcome these critical obstacles, several strategies have been developed well. First, smart construction of nanoframework, especially, one-dimensional (1D) ultralong nanowires (NWs), would ensure rapid charge transfer of the electrode coupled with small contacting resistance [21e24]. Secondly, mesoporous architectures can be designed to enlarge electrode/electrolyte contact area, shorten Liþ diffusion path, and buffer the strain of volume change during continuous lithiation-delithiation processes, contributing to the improved electrochemical performance [22e24]. Thirdly, the nano-carbon or conducting polymer (CP) coating layer can be constructed as an efficient cushion and/or protecting layer to alleviate volume variation, impede particle aggregation, and passivate the negative effect from the solidelectrolyte-interphase (SEI) film, improving the high-rate cycling stability of the anodes [18,25,26]. Moreover, the conductive coating layer can enhance the electric conductivity of electrodes, which is in favor of the excellent rate performance [18,25,26]. Typically, the polypyrrole (PPY) with simple processability, high electronic conductivity and chemical stability can serve as an excellent coating candidate for high-performance LIBs [27,28]. One especially notes that each solution applied alone usually generates limited enhancement in electrochemical Li-storage performance, despite respective superiorities of the above smart strategies. Therefore, it is highly desired for advanced LIBs by developing an integrated solution to render high-performance ZFO-based anodes. With the above discussions in mind, herein, we first devised an electrospinning technique along with subsequent calcination and gas polymerization to construct 1D ZFO-based multi-componential hybrid ultralong NWs, designed as ZFO@PPY, in which the nanoparticles (NPs)-assembled ultralong ZFO NWs were bestowed with continuous PPY nanoshell coating on each ZFO subunits throughout the whole NWs. The as-synthesized durable ZFO@PPY NWs exhibited outstanding Li-storage performance with a reversible capacity of ~881 mAh g
1 after 220 cycles at a specific current of 1 A g
1, when evaluated as low-cost anode for LIBs, thanks to the continuous PPY network in the mesoporous hybrid NWs. Besides, the understanding in Li-storage behaviors of the ZFO@PPY NWs during charge-discharge processes was proposed through the insitu XRD technology measurement. 2. Experimental 2.1. Preparation of the ZFO@PPY NWs All the chemicals (Sinopharm Chemical Reagent Co., Ltd) used here were of analytical grade. Typically, 2 g of polyvinylpyrrolidone (PVP, K90) was dissolved in 18 mL of absolute ethanol under stirring for 8 h at room temperature (RT). 1 mmol of Zn(CH3COO)$22H2O and 2 mmol of Fe(NO3)$39H2O were dissolved in 10 mL of N, N-Dimethylformamide under stirring for 30 min at RT. Afterwards, the above two clear solutions were mixed together. Then, 0.24 mmol of terephthalic acid was added. After stirred for another 2 h, the final homogeneous solution was loaded into a syringe (10 mL) with a 20gauge blunt tip needle for electrospinning. A voltage of 22 kV was set for electrospinning. The distance between the needle and collection board was 15 cm. The feeding speed was 0.8 mL h
1. The obtained precursor containing zinc/ferric salt (designed as ZnFeprecursor NWs) was obtained after dried in air at 60  C, and further subsequently calcinated at 500  C for 5 h in a tube furnace (GSL-1500X) in air with a ramp of 0.5  C min
1. The obtained sample was denoted as ZFO NWs.

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The PPY coating was performed as follows. Typically, 0.15 g of ZFO NWs was immersed in 0.1 mL of FeCl3 solution (0.3 M). After further dried at 60  C for 6 h, the obtained ZFO NWs placed in a vial, together with another vial containing 2 mL of pyrrole monomer, were sealed in a blue cap bottle. Then the bottle was heated at 50  C for 48 h. After washed with water and dried, the hybrid ZFO@PPY NWs were obtained.

2.2. Materials characterization The crystal structure of the samples was analyzed by X-ray diffraction (XRD, Rigaku-TTRШ type, Cu Ka, 40 kV, 300 mA, Japan) with a scanning speed of 2 min
1 in the 2q range from 20 to 80 . The morphology and structure were detected by field-emission scanning electron microscope (FESEM, JEOL S-4800), transmission electron microscopy (TEM), scanning TEM (STEM), high-resolution TEM (HRTEM), selected area electron diffraction (SAED) (JEOL, JEM2100 system). Energy dispersive X-ray (EDX) analysis and corresponding elemental mapping data were taken with the X-ray spectroscopy attached to the TEM instrument. Thermogravimetric (TG) analysis and differential scanning calorimeter (DSC) were performed with a NETZSCH STA 449 PC system TG Analyzer under air flow with a heating rate of 10  C min
1. X-ray photoelectron spectroscopy (XPS) was performed by using a Thermo ESCALAB 250Xi spectrometer with an Al Ka excitation source. N2 sorption isotherms were determined by Brunauer-Emmett-Teller (BET) with the surface area analyzer (Autosorb-iQ). Fourier transform infrared (FT-IR) spectra were recorded on a Nicolet iS50 spectrometer.

2.3. Electrochemical measurements A coin cell (CR2025) was assembled to evaluate the electrochemical performance. The electroactive materials, acetylene black and sodium carboxymethyl cellulose (average Mw: ~250 000) with a mass ratio of 70 : 20: 10 were mixed manually with water to obtain a slurry in an agate mortar. The obtained homogeneous slurry was coated on Cu foil. Then, the coated Cu foil was dried at 100  C for 11 h in a vacuum oven (DZF 6020) as the working electrode. Li foil (Cyprus Foote Mineral, 99.98%, USA) was used as the counter electrode and polypropylene micro porous membrane was applied as the separator. The electrolyte was 1 M LiPF6 in ethylene carbonate/dimethyl carbonate/ethylmethyl carbonate with a volume ratio of 1 : 1: 1. The cells were assembled in an Ar-filled glovebox (MBRAUN UNILAB PRO) with oxygen and moisture contents both under 0.5 ppm. Typical loading of electroactive materials is ~1.0e1.2 mg per cm2. The multi-channel battery testing system (LAND CT2001A) was used to record the galvanostatic charge/ discharge data with a potential range from 0.01 to 3.0 V (vs. Li/Liþ). Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) with an AC signal amplitude of 5 mV over a frequency range from 10
2e105 Hz were performed on an IVIUM electrochemical workstation (the Netherlands). The in-situ analysis upon galvanostatic lithiation and delithiation was performed by XRD (Rigaku-TTRШ type) with an in-situ cell (Rigaku, Japan). The electrode slurry was obtained as mentioned above, and then casted on a Be window (0.1 mm in thickness, Northwest Institute for Non-ferrous Metal Research, China), which was used both as the current collector and window for the X-ray beam. After dried under vacuum overnight, the cell was assembled. And the cell was galvanostatically cycled by a testing system (LAND CT2001A) applying a specific current of 100 mA g
1, and the XRD analysis was performed meanwhile.

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3. Results and discussion 3.1. Physicochemical characterizations The synthesis process of 1D porous ZFO@PPY NWs is schematically illustrated in Fig. 1. Firstly, the primary polymer ZnFeprecursor NWs are electrospun, and then subjected to thermal treatment in air to remove the residual solvent and stabilize the structure. After calcinated at 500  C, the mesoporous ZFO NWs are obtained. The gas-phase polymerization is then conducted to uniformly coat the resultant ZFO NWs. Due to the low vapor pressure of the pyrrole, monomer molecule can easily vaporize, diffuse into, and reach nearly all the surface of porous ZFO NWs. The polymerization reaction of the pyrrole in-situ occurs on the surface of ZFO NWs with the catalysis of the FeCl3, resulting in the formation of 1D hybrid ZFO@PPY NWs. Crystalline structure and phase purity of the samples are examined by X-ray diffraction (XRD) technique. As shown in Fig. 2, the ZFO and ZFO@PPY NWs exhibit quite similar XRD patterns. All the diffraction signals can be indexed to the cubic spinel ZnFe2O4 (JCPDS no. 22-1012) with a space group Fd-3m. And no any reflections belonging to other impurities can be found, indicating the good maintenance of crystalline ZFO NWs after coating the PPY. The tangerine ZFO NWs graduate into black green materials with the polymerization coating, as visualized in the digital photographs (the insets in Fig. 2), fully confirming the successful coating by the PPY. To further support the existence of PPY in the resultant ZFO@PPY NWs, Raman and Fourier transform infrared (FT-IR) spectra are collected accordingly. Fig. 3a exhibits the Raman data of the ZFO and ZFO@PPY NWs. As for the ZFO NWs, the spectrum shows five typical Raman modes at 354.7, 501.7, 648.5, 1113.1 and 1291.5 cm
1, respectively. As reported for the cubic spinel structure, wave numbers above 600 cm
1 should be attributed to the motions of O in tetrahedral AO4 groups [29e31]. Accordingly, the mode at 648.5 cm
1 can be assigned to A1g symmetry, and the other two low-frequency modes (354.7 and 501.7 cm
1) belong to the characteristic peaks of octahedral sites (BO6), corresponding to the F2g mode [29e31]. The two bands located at 1113.1 and 1291.5 cm
1, as observed in Fig. 3a, are suspected to be overtone peaks, which are also found in the previous literature [32]. Distinct from the ZFO NWs, other six prominent peaks related to the PPY emerge along with the weakening peaks of the ZFO, evidently verifying the welldeveloped PPY coating layer. Noteworthy, just a weak peak at 648.5 cm
1 can be easily discerned for the characteristic band of the ZFO. Specifically, the peaks at 922.3 and 990.9 cm
1 are assigned to the ring deformation associated with the bipolaron and polaron. The band at 1046.8 cm
1 can be attributed to the CeH in-

Fig. 2. Typical XRD patterns and digital images of the ZFO (the upper) and ZFO@PPY (the lower) NWs.

plane bending vibration, while the strong peaks at 1572.1, 1421.4 and 1331.9 cm
1 are corresponded to C]C backbone stretching and ring stretching [33e35]. As observed in the FT-IR spectra (Fig. 3b) of the ZFO@PPY and PPY NWs, both of them exhibit characteristic peaks at 1631.2, 1384.9, 1263.1, 1042.5, 924.9 and 792.4 cm
1. Specifically, the peak at 1631.2 cm
1 is ascribed to C]C stretching vibration, while the ¼ CeH in-plane vibration and CeN stretching vibration are found at 1384.9 and 1263.1 cm
1, respectively. The absorption at 1042.5 cm
1 belongs to the ¼ CeH in-plane vibration. Additionally, the bands at 924.9 and 792.4 cm
1 result from the CeH ring out-of-plane bending vibration [28,36]. Thermogravimetric (TG) measurements are performed in air to quantify the specific content of the PPY in the final NWs. As shown in Fig. 3c, the slight weight loss (~2.2 wt%) in the whole temperature range should be attributed to the combustion of the residual polymers in the ZFO NWs. By contrast, the large mass loss (~11.5 wt%) in the ZFO@PPY NWs should be assigned to the combustion of PPY in air besides the removal of residual polymers in the ZFO NWs, according to the endothermic peak in differential scanning calorimeter (DSC) profile. As a result, the PPY content can be calculated as approximately 9.3 wt% in the hybrid NWs. More detailed information about the surface state of the two NWs is gathered by X-ray photoelectron spectroscopy (XPS) method. The full survey spectra (Fig. 4a) reveal the presence of Zn, Fe, O and C in both samples, while obvious differences can be also observed. Compared with the ZFO NWs, the peaks of the Zn and Fe for the ZFO@PPY NWs are subdued and the ones for the O and C are strengthened. Besides, another peak at about 400 eV, belonging to the elemental N, only appears in the curve of the ZFO@PPY. These

Fig. 1. Schematic illustration for the synthesis of the 1D ZFO@PPY hybrid NWs.

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Fig. 3. (a) Raman spectra of the ZFO and ZFO@PPY NWs, (b) FT-IR spectra of the ZFO@PPY and PPY, and (c) TG (black color) and DSC (blue color) plots of the ZFO (solid line) and ZFO@PPY (dot line) samples as indicated. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

changes should be derived from the PPY coated upon the surface of ZFO NWs. The high-resolution Fe 2p spectra and corresponding fitted curves for the ZFO and ZFO@PPY NWs (Fig. 4b) are contrastively shown with no visible differences in bonding energies between the two. The fitting peaks for the Fe 2p3/2 locating at 710.7 and 712.5 eV are assigned to tetrahedral and octahedral sites, respectively. The other three peaks correspond well to the shake-up satellite and Fe 2p1/2, indicative of the existing Fe3þ in the samples [13,37]. The high-resolution Zn 2p spectra (Fig. 4c) are also finely fitted with two peaks at 1021.3 and 1044.5 eV, which are attributed

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to Zn 2p2/3 and 2p1/2, respectively, revealing the presence of Zn2þ in the composites [13,37]. The high-solution O 1s spectra (Fig. 4d) of the two are deconvoluted into four peaks at 529.5, 530.4, 531.7 and 532.4 eV. The peaks centered at 529.5 and 530.4 eV are related to the typical lattice oxygen in metal (Zn/Fe)-oxygen framework, and other physi-/chemi-adsorbed O species (i.e., OH, H2O or carbonate) on the sample surfaces. The bands locating at 531.7 and 532.4 eV are ascribed to the defect sites with low oxygen co-ordination in ZFO [24,37,38]. The N-bonding configurations in the ZFO@PPY NWs, originating from the PPY, are further studied by the high-resolution N 1s spectrum (Fig. S1, Supporting Information). The main fitted peak at 400 eV can be attributed to the eNH group of pyrrole unit, and the ones locating at 398.5 and 401.5 eV are assigned to C]N defects of the PPY and CeNþ/C]Nþ structures [39]. All the aforementioned results certify the existence of PPY in the ZFO@PPY NWs. Fig. 5a shows the low-magnification field-emission scanning electron microscope (FESEM) image of the resultant ZFO NWs. Typical NWs with the lengths up to dozens of micron are observed, similar to its precursor (Fig. S2). The amplifying images (Fig. 5b and c) obviously show that each NW with a uniform diameter of ~70 nm consists of many NP building blocks. To obtain the interior structure of the ZFO NWs in detail, transmission electron microscopy (TEM) measurement is taken accordingly. As shown in Fig. 5def, numerous NPs with an average size of ~20 nm are assembled into the ZFO NWs. Interestingly, many pores distributed throughout the NWs, which stem from thermal decomposition of PVP and the generation of gaseous substances (CO, CO2, etc.) [3,13,23,24], can be clearly observed. The loosely porous structure is beneficial to the integrated coating of the PPY on the each ZFO subunits with the assistance of the smart gas-phase polymerization. The HRTEM image (Fig. 5g) of the NPs clearly shows the lattice fringes with a spacing of 0.488 nm, which can be indexed well to the (111) crystalline plane of the spinel ZFO. A sequence of concentric circles observed in the selected area electron diffraction (SAED) pattern (Fig. 5h) match well with the crystalline planes of the ZFO, revealing the polycrystalline feature of the ZFO NWs. Fig. 5i displays the energy dispersive X-ray (EDX) elemental mapping images of the resulted ZFO NWs, revealing the co-existence and homogeneous distribution of Fe and O in the whole NWs. Particularly, compared with the elements of Fe and O, a tiny amount of the elemental C still can be detected, which is consistent with the above TG analysis (Fig. 3c). FESEM and TEM measurements are further carried out to study the effect of the PPY coating on the microtopography and structure of the ZFO@PPY. As noted, no obvious change can be seen in respect to the length and diameter of the NWs after coated with PPY (Fig. S3). However, the surface of ZFO@PPY NWs became even smoother when compared with that of the pristine ZFO NWs, as seen in Fig. 6a. Clearly, the pristine ZFO NWs are integrally and uniformly coated by PPY and the thickness of PPY layer is around 2e4 nm. In HRTEM observations (the insets in Fig. 6b), the lattice fringe for the typical inner NPs can be clearly distinguished. Corresponding lattice spacings are calculated as 0.244 nm and 0.298 nm, which can be indexed to (222) and (220) crystal planes of the ZFO. Besides, the SAED with a series of concentric circles (Fig. 6c) suggests the polycrystalline characteristic of the ZFO@PPY NWs, and these diffraction rings can be assigned to the (220), (311), (222), (331) and (511) planes of the ZFO from inside to outside. The EDX characterization (Fig. S4) confirms the co-existence of the Fe, Zn, O, C and N elements in the ZFO@PPY NWs, which is in good agreement with the XPS data above. As seen from the line-scanning profiles (Fig. 6d and e) and surface-scanning element mapping images (Fig. 6f), the Fe and O mainly distribute in center area, while C and N appear in more widely spread area, indicating the

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Fig. 4. XPS spectra and fitted data of the ZFO and ZFO@PPY NWs: (a) survey spectra, (b) Fe 2p, (c) Zn 2p, and (d) O 1s.

Fig. 5. (aec) FESEM images, (def) TEM images, (g) HRTEM image, (h) SAED pattern, and (i) STEM and corresponding elemental (Fe, O and C) mapping images of the ZFO NWs.

homogeneously coating throughout the whole NWs. The specific surface area (SSA) and porous nature of the ZFO and ZFO@PPY NWs are investigated by N2 adsorption/desorption measurements (Fig. 7a). Clearly, both the curves can be classified as

type-IV isotherms for the two, along with the distinct H3-type hysteresis loop according to the IUPAC classification. It is worth noting that the isotherm of the ZFO@PPY sample exhibits larger hysteresis loop than that of ZFO NWs, indicating more abundant

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Fig. 6. (a) TEM image, (b) HRTEM image, (c) SAED pattern, (d) TEM image and (e) corresponding element (Fe and C) line scanning, and (f) STEM and corresponding EDX elemental (Fe, O, C and N) mapping images of the ZFO@PPY NWs.

~80.8 m2 g
1, which is 1.9 times higher than that of ZFO NWs (~42.6 m2 g
1). The average pore size of the ZFO@PPY NWs is slightly smaller than the ZFO (4.7 vs. 5.1 nm), and the pore volume of the ZFO@PPY NWs is greatly enhanced with the PPY coating (from ~1.1 to ~2.0 cm3 g
1), further facilitating the electrodeelectrolyte contact and convenient mass transfer [40]. The difference in SSA and pore volume should be related to the contribution of the extra PPY phase in the hybrid NWs.

3.2. Electrochemical evaluation

Fig. 7. (a) N2 sorption isotherms and (b) pore size distributions of the ZFO and ZFO@PPY NWs.

mesoporous structure. The pore size distribution (PSD) is analyzed based on the adsorption branch of the isotherms, presenting the small pore size in the ZFO@PPY NWs. The PSD curves of the samples are contrastively presented (Fig. 7b). The quantitative data are calculated and summarized (Table S1). The SSA of ZFO@PPY NWs is

Considering these appealing structural and compositional advantages above for LIBs, the half-cell electrochemical properties of the ZFO@PPY NWs were investigated by assembling the coin cell. Electrochemical performance of the samples is first evaluated using cyclic voltammetry (CV) technique. Fig. 8a exhibits the typical CV profiles of the ZFO@PPY and ZFO NWs within a potential range from 0.01 to 3.00 V (vs. Li/Liþ) [3,12e20]. For both of the samples, the first cycle is obviously different from the following ones. During the first cathodic process for the ZFO@PPY NWs, two small peaks locating at around 1.53 V (A) and 0.85 V (B) can be observed, followed by a strong reduction peak at about 0.52 V (C). Below 0.3 V, a broad peak (D) can be also identified. These peaks correspond to the different structure transformations and redox reactions occurring in the potential sweep, such as the Li-intercalation into ZFO (A), formation of the crystalline ZnxFeyO (B), further reduction of ZnxFeyO to Zn0/ Fe0 and Li2O (C) and LieZn alloying reaction (D) [41e43]. More indepth discussions will be conducted in the following section about in-situ XRD evaluation. In the subsequent anodic sweep, two broad oxidation peaks around 0.7e0.8 V and 1.6e1.7 V can be observed, related to the de-alloying processes of LieZn alloys and the reoxidation of Zn0 and Fe0, respectively [41e43]. During following cycles, no significant change can be found on these CV profiles, indicating high electrochemical reversibility and structural stability after the initial delithium process. One note that the stable and reversible reduction peak of the ZFO@PPY NWs positively shifts, compared to the ZFO NWs, indicating the enhancement in electrochemical reversibility and the decrease in polarization. Galvanotatic discharge/charge tests are carried out to further evaluate the electrochemical performance of the ZFO and ZFO@PPY NWs at a specific current of 50 mA g
1. Fig. 8b comparatively displays the typical discharge-charge plots of the electrodes. In the initial discharging plots, both of the ZFO and ZFO@PPY NWs show similar

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Fig. 8. (a) Initial CV curves (1.0 mV s
1) and (b) voltage profiles of the galvanostatic discharge-charge plots (50 mA g
1) of the ZFO (the upper) and ZFO@PPY (the lower) NWs.

profiles with four distinct stages, corresponding to the A/B/C/D peaks on the CV curves (Fig. 8a). As for the ZFO NWs, the first discharge/charge capacities are ~1041 and ~858 mAh g
1, giving an initial Coulombic efficiency (CE) of ~82.4%. Compared with the ZFO NWs, the first discharge capacity of the ZFO@PPY anode increases up to ~1108 mAh g
1, while the charge capacity is ~827 mAh g
1, resulting in a reduced CE of ~74.6%. This phenomenon should be attributed to the large SSA of the ZFO@PPY NWs, which results in more production of SEI film and consumption of Liþ, eventually leading to the higher discharge capacities and lower CE of the ZFO@PPY anode [44,45]. During the following charge-discharge processes, the CE values of the two both rapid reach ~100%, suggesting the excellent reversibility of the materials. In order to better understand specific electrochemical processes of the ZFO@PPY NWs for LIBs applications, the in-situ XRD measurements coupled with galvanostatic charge/discharge tests (100 mA g
1) were systematically carried out. An overview of the evolution of diffraction reflections is displayed in Fig. 9aec. Typically, Fig. 9a exhibits the contour map of in-situ XRD patterns obtained during the initial lithiation and delithiation processes. The color variations in the contour map intuitively show the transformation of crystal structure during charge/discharge processes. The waterfall diagram (Fig. 9b) of the in-situ patterns and corresponding galvanostatic discharge-charge profiles (Fig. 9c) are profiled to reveal the variations of the ZFO@PPY anode with the voltage change, where several distinct regions with substantial change are labeled accordingly. The initial discharge and charge plots, as shown in Fig. 9c, can be typically divided into four (denoted as A/B/ C/D) and three (denoted as E/F/G) voltage stages, respectively. The XRD signals (Fig. 9b), corresponding to the different chargedischarge states of the ZFO@PPY, are distinguished with distinct colors. In the A/B/C stages of the discharge process, the representative XRD patterns show apparent differences. The shift of the reflection position represents the change in crystal lattice parameters, which is triggered by the continuous lithiation process. The disappearance of these characteristic peaks and appearance of other new characteristic peaks corroborate the occurrence of phase transformation. When discharged at the stage D, the intensity of characteristic reflections gradually become weak and finally disappear, which means that the electroactive material transforms from the crystalline to amorphous state. In the following charge process, no obvious new diffraction peak appears. This phenomenon combined with the CV observation (Fig. 8a) authenticates the fact that amorphous state should be always maintained in the subsequent charge-discharge cycles.

For even better visualization of the labeled areas, several regions with the substantial changes in Fig. 9b are magnified, and further collected in Fig. 9d. As evidenced, the variation tendency of the characteristic peaks is indicated by the blue arrow. In the patterns for the regions I/II/III/IV, the red curves synchronized with the discharge stage A (corresponding to the reduction peak at 1.53 V in Fig. 8a) exhibit four reflections at ~35.3, ~42.8, ~56.6 and ~62.2 , respectively, which can be indexed well to the lattice planes (311), (400), (511) and (440) of the cubic ZFO (JCPDS no. 22-1012), respectively. With the end of the discharge platform A, the four characteristic reflections slightly shift to the lower angles at ~35.2 , ~42.7, ~56.5 and ~62.1, respectively, which perfectly match with the crystal faces of the (311), (400), (333) and (440) of the Li0$5ZnFe2O4 (JCPDS no. 40-1116). It therefore can be inferred that the lithium ions embed into the interstitial void of the spinel ZFO, resulting in the formation of Li0$5ZnFe2O4 during the discharge stage A. In the process, the lithiation reaction does not break the crystal structure of cubic ZFO at all, but just changes the lattice parameters of the spinel structure. The blue plots in the regions I/II/III/IV are recorded in the discharge process of the stage B, corresponding to the peak at 0.85 V in CV curves. In the region I, the reflection intensity at ~35.2 continuously decreases, while the reflection at ~36.8 gradually shifts to lower angle (~0.45 ) and the intensity is simultaneously enhanced following the finish of the stage B. In the region II, the reflection at ~42.7 slightly moves to lower angle (~0.53 ), along with the intensity increasing. The diffraction intensity of the reflection locating in the region III exhibits a rapid decrease until vanishes at the end of the stage B. In the region IV, the diffraction reflection continuously shifts to a lower 2q value (~0.35 ), while the intensity almost does not change any more. According to the above results and analysis, the evolution of diffraction signals in the stage B is more complicated than that in the stage A, which suggests the phase and/or structure of the Li0$5ZnFe2O4 have greatly changed. Although no any standard cards in the JCPDS database can be found matching well with the new diffraction patterns obtained after the stage B, the diffraction reflections are similar to the ZnxFeyO representing some oxides with low valence states of Zn/Fe and/or the ZnFe-alloy. Therefore, it can be rationally deduced from the aforementioned results that the further lithiation of the Li0$5ZnFe2O4 in discharge process of the stage B results in the reduction of highvalence Zn/Fe species and crystal structure transformation, eventually resulting in the formation of the ZnxFeyO. Unlike the previous report [46], a newly formed Li1$45ZnFe2O4 with further lithiation cannot be detected in our case before the transformation from

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Fig. 9. (a) Contour maps of in-situ diffraction patterns of the ZFO@PPY anode within 2q ranges of 27e45.6 , 64e63.5 and 72e79 , (b) waterfall diagram showing the evolution of the in-situ patterns at different lithiation and/or delithiation states indicated by various colors, corresponding to the regions from A to G, respectively, as indicated in (c) voltage profile of the 1st discharge-charge cycle, and (d) enlarged XRD patterns corresponding to the regions (I
VII) in the panel (a). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

Li0$5ZnFe2O4 to the ZnxFeyO phase. And the observation here is the same to the Li-storage processes of other ZFO-based anodes reported before [3,6,11,12,14e20,22e24]. The XRD profiles in the regions V/VI/VII are collected in the main discharge plateau of the stage C, corresponding to the distinct reduction peak at ~0.52 V (vs. Li/Liþ) in the CV plots (Fig. 8a). It should be noted that a specific capacity of ~583 mAh g
1 is obtained during this long voltage plateau, which means that the reduction reaction of electroactive material with Liþ has occurred. A gradual decrease in reflection intensity for the ZnxFeyO can be observed in the XRD patterns of the regions V/VI/VII, and finally these reflections vanish at the end of the stage C in discharge process, while no any new characteristic reflections emerge during this process. According to the reported literature [41,46,47], the products of the reduction reaction of the ZnxFeyO should be Zn0, Fe0 and Li2O. The above analysis indicates that the crystalline ZnxFeyO material transformed into amorphous or quasi-amorphous (i.e., the crystalline grains are too small to be detected by XRD technique) Zn0, Fe0 and Li2O. This transformation process also can be clearly verified in three-dimensional in-situ XRD patterns of the ZFO@PPY NWs (Fig. S5). It is well known that the potential region of the stage D corresponds to the formation process of LiZn-alloy [14,41]. But in

our case, no any XRD signals associated to the crystalline Zn and/or LiZn-alloy can be discerned in the diffraction profiles obtained in voltage stage D, which further approves the amorphous nature of the products in the process. Based on the above analysis, the produced amorphous Zn0 continues to react with the Liþ, generating the amorphous ZnLi-alloy in the subsequent discharge process of the stage D. During the delithiation process, as shown in Fig. 9b and c, the first charge plot can be divided into three voltage stages (E/F/G), and corresponding in-situ diffraction patterns are also collected. One particularly note that the electroactive phases always remain the amorphous state during the following charge process, that is, the resultant ZnLi-alloy is first oxidized to amorphous Zn0, then these amorphous Zn0 and Fe0 are further oxidized to the amorphous ZnO and Fe2O3, which are highly electrochemical reversible in the subsequent discharge-charge cycles. The superior rate performance of the anodes is one of the crucial factors for their practical application. To evaluate the rate properties of the ZFO and ZFO@PPY NWs, the assembled half-cells are tested at gradually changed specific currents ranging from 0.05 to 10 A g
1. As plotted in Fig. 10a, both of the two electrodes demonstrate outstanding cyclic stabilities at the low specific current of

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0.05 A g
1, and the ZFO NWs even display a higher average specific capacity than ZFO@PPY NWs (i.e., ~919 vs. ~793 mAh g
1). When the specific current rises to 0.1 A g
1, the capacity of the ZFO electrode shows a fast decay, while the ZFO@PPY counterpart still presents excellent stability. The capacities of ZFO NWs are lower than those for the ZFO@PPY NWs since the 20th cycle. At an ultrahigh specific current of 10.0 A g
1, the ZFO NWs only deliver a capacity of ~51 mAh g
1, which is ~5.1% of the theoretical capacity of the ZFO, while the ZFO@PPY NWs even give a large capacity of ~287 mAh g
1, which is as high as ~28.7% of the theoretical capacity of the ZFO. After 80 cycles with the gradually increased currents from 0.05 to 10 A g
1, the specific current is returned to 0.05 A g
1 again. As for the ZFO NWs, the capacity does not display timely response, which first recovers a low capacity of ~264 mAh g
1, and then further increases to ~569 mAh g
1. While for the ZFO@PPY NWs, a large capacity of ~822 mAh g
1, similar to the initial value of ~793 mAh g
1, is immediately achieved after returning the specific current to 0.05 A g
1. Whereafter, the current stepwise gets back to 0.2 A g
1, the capacities of both electrodes gradually increase with the cycle proceeding. Finally, the capacities of the ZFO and ZFO@PPY NWs reach up to ~917 and ~1235 mAh g
1 in the 300th cycle, respectively. This phenomenon should be attributed to the reversible formation of electrochemical active polymeric gel-like films deriving from kinetic activation in the electrode [3,11,46,48e50]. The above results affirm that the rate property of the ZFO@PPY NWs is much superior to that of the ZFO NWs. In addition to large reversible capacities and superior rate performance, long-term cycling durability is another critical aspect for high-performance LIBs application. Fig. 10b contrastively displays the cycling performance of the ZFO and ZFO @PPY NWs at a specific current of 0.2 A g
1 in the voltage range of 0.01e3.00 V (vs. Li/Liþ). During the initial dozens of cycles, the ZFO NWs exhibit rapid capacity decay. The capacity of the ZFO NWs decreases from ~949 (the 2nd cycle) to ~607 mAh g
1 (the 54th cycle), and the capacity retention is calculated as low as ~56.6%. During the subsequent cycles, the capacity of ZFO NWs rises again and maintains at ~802 mAh g
1 after 240 cycles, however, the initial rapid capacity decay and final capacity fluctuation reveal the poor stability of the ZFO NWs. By contrast, the cycling profile of ZFO@PPY NWs does not show visible capacity loss in the initial cycles. The lowest capacity value is ~767 mAh g
1, corresponding to the 4th cycle. Since then,

Fig. 10. (a) Rate capabilities with the specific currents from 0.01 to 10 A g
1 and (b) cycling performance (0.2 A g
1) of the ZFO and ZFO@PPY NWs.

the capacity of the ZFO@PPY anode gradually increases with the cycle going on, and reaches approximately 1346 mAh g
1 after 240 consecutive cycles. The increased reversible capacity and excellent stability of the ZFO@PPY electrode should result from the following aspects. First, the superior conductivity of the PPY coating layer can be beneficial to improve the electrode dynamics and activate the material. Next, the integrated PPY coating layer can act as a cushion to buffer volume expansion and structural stress produced in the lithiation-delithiation process. Finally, the reversible formation of a polymeric layer on the surface of NWs would also contribute to the increasing capacity with cycling [3,4,46,48e50]. Particularly, it is worthy of noting that the CE values of the ZFO@PPY anode are basically higher than those for ZFO NWs (Fig. S6) over the consecutive cycle, which indicates that the PPY coating, to some extent, can restrain the irreversible formation of extra SEI over the subsequent charge-discharge cycles. The large capacity and long-cycle performance especially at large specific currents play a critical role in the EVs and/or hybrid EVs requiring fast charge or discharge. And the high-rate cycling performance of the ZFO@PPY anode is tested at 1 A g
1, as depicted in Fig. 11a. Although the capacity of the hybrid NWs decreases in the early fifteen cycles, it still delivers a capacity of ~622 mAh g
1. And whereafter, the capacity gradually increases with the chargedischarge cycle proceeding, which may be attributed to the continuous activation of the electrode and formation of the polymeric gel-like films [3,4,46,48e50]. Encouragingly, the ZFO@PPY hybrid NWs can render a large reversible capacity of ~881 mAh g
1 and maintain a CE of ~100% after 220 cycles, indicating its excellent

Fig. 11. (a) Cycling performance of the ZFO@PPY NWs over 220 cycles at 1 A g
1, and (b, c) TEM, (d) HRTEM, STEM and corresponding elemental (Fe, C and N) mapping images (the upper left inset in panel d) for the cycled ZFO@PPY electrode.

L. Hou et al. / Electrochimica Acta 306 (2019) 198e208

cycling durability and large capability at high rates. To further confirm the structure stability of the ZFO@PPY anode, corresponding TEM technique is performed after 220 cycles at 1 A g
1. As depicted in Fig. 11b and c, the cycled ZFO@PPY electrode still remains the pristine morphology of 1D NWs, and porous structure also can be clearly observed. The evidence confirms the superior structure stability of the ZFO@PPY NWs over cycling, which should benefit from the protection of the integrated PPY buffer layer. But beyond that, deeper changes are further detected by the HRTEM, as displayed in Fig. 11d. In the HRTEM image of the cycled electrode, no any lattice fringe can be observed, indicating that the pristine crystalline ZFO@PPY NWs transform to the amorphous state during the sustaining charge-discharge processes, which is consistent with the in-situ XRD measurements above. Owing to the amorphous feature of the ZFO NWs after cycling, it is very difficult to distinguish the distinct ZFO-PPY interface from the HRTEM image (Fig. 11d). However, it is particularly worthy of noting that the uniform distributions of elemental Fe, C and N in the cycled ZFO@PPY NWs, as evidenced from the STEM and corresponding EDX mapping images (the inset in Fig. 11d), still can be observed, which strongly verifies that the continuous PPY network still exists throughout the hybrid NWs. Out of question, the desirable structural stability of the ZFO@PPY anode over cycling will enormously favors for its enhanced electrochemical properties, as discussed above. To reveal the effect of the PPY coating on electrode conductivity, electrochemical impedance spectroscopy (EIS) tests are performed. As plotted in Nyquist files (Fig. 12) of the fresh ZFO and ZFO@PPY anodes, both of the plots show typical characteristics of a distorted semicircle in the high-frequency region and a slant line in the lowfrequency region. The EIS results are fitted according to the fitting equivalent-circuit model (Fig. S7), and corresponding fitted data are summarized (Table S2), in which Rs, Rct and W represent the solution resistance (the summary of electrolyte resistance, intrinsic resistances of separators and electroactive materials themselves, and contacting resistance between electroactive materials and collector), charge transfer resistance on the electrode/electrolyte interface, and Warburg impedance [37,47]. Evidently, the values of Rs and Rct for the ZFO@PPY NWs (~0.8 and ~98 U) are both much smaller than those for the ZFO NWs (~1.3 and ~261.3 U), representing reduced electronic and charge transfer resistances after the PPY coating with the same cell assemble technique. In addition, Warburg impedance is also decreased from 0.00419 to 0.00107U,

207

indicating the more fast Liþ diffusion in the ZFO@PPY hybrid electrode. These results confirm that the PPY coating layer effectively enhances the conductivity of the ZFO@PPY NWs, which renders a significant improvement in comprehensive electrochemical performance. As an advanced anode material, encouragingly, our as-designed ZFO@PPY anode exhibits the comparable first CE, large reversible capacities and superior cycling stability at the same case. The remarkable lithium storage performance of the ZFO@PPY NWs is ascribed to the following aspects. First, the 1D ultralong NW structure can supply unique conductive pathways for efficient charge transfer in the electrochemical processes [21e24,46]. Second, the integrated PPY network not only improves the conductivity of the whole electrode but acts as a cushion to alleviate volume expansion and structural stain produced in the lithiationdelithiation process and maintain structural stability of the electrode. Finally, the porous structure shortens the Liþ transport length and provides more electrochemical active sites and additional free void to alleviate the structural stress. It is with these structural and compositional advantages above that the ZFO@PPY NWs simultaneously exhibit high reversible capacities, long-term stability and superior rate performance. 4. Conclusions In summary, we have successfully designed and fabricated 1D ultralong ZFO@PPY NWs by combining scalable electrospinning technique with efficient gas-phase polymerization approach. When utilized as the anode material for LIBs, the hybrid ZFO@PPY NWs displayed remarkable lithium storage performance, including large reversible capacity, long-term durability and superior rate performance. Corresponding In-situ XRD and ex-situ TEM measurements revealed the crystalline structure and phase transformation processes in the repeated lithiation-delithiation process. The outstanding Li-storage performance can be attributed to the unique composition and structure advantages in terms of 1D ultralong NW architecture, porous feature, and the integrated PPY coating, which endowed the ZFO@PPY NWs with high accessibility, high conductivity, large active sur-/interfaces and structural stability. More promisingly, our integrated design strategy can be extended to other electrode materials for next-generation LIBs. Acknowledgements The authors acknowledge the financial support from the National Natural Science Foundation of China (No. 51572005, 51772127, 51772131), the Taishan Scholars (No. ts201712050), Major Program of Shandong Province Natural Science Foundation (ZR2018ZB0317), and Shandong Collaborative Innovation Center of Technology and Equipements for Biological Diagnosis and Therapy. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.electacta.2019.03.121. References

Fig. 12. EIS plots and corresponding fitted profiles of the fresh ZFO and ZFO@PPY NWs. The inset for the enlarged EIS files in high-medium frequency region.

[1] F. Cheng, J. Liang, Z. Tao, J. Chen, Adv. Mater. 23 (2011) 1695. [2] J.B. Goodenough, K.S. Park, J. Am. Chem. Soc. 135 (2013) 1167. [3] L.R. Hou, L. Lian, L.H. Zhang, G. Pang, C.Z. Yuan, X.G. Zhang, Adv. Funct. Mater. 25 (2015) 238. [4] P. Poizot, S. Laruelle, S. Grugeon, L. Dupon, J.M. Tarascon, Nature 407 (2000) 496. [5] G.Q. Zhang, L. Yu, H.B. Wu, H.E. Hoster, X.W. Lou, Adv. Mater. 24 (2014) 4609. [6] J.H. Kong, X.Y. Yao, Y.F. Wei, C.Y. Zhao, J.M. Ang, X.H. Lu, RSC Adv. 5 (2015) 13315.

208

L. Hou et al. / Electrochimica Acta 306 (2019) 198e208

[7] C.F. Zhang, J.S. Yu, Chem. Eur J. 22 (2016) 4422. [8] G.X. Gao, S.Y. Lu, B.T. Dong, W. Yan, W. Wang, T. Zhao, C.Y. Lao, K. Xi, R.V. Kumar, S.J. Ding, J. Mater. Chem. A 4 (2016) 10419. [9] S. Permien, S. Indris, U. Schürmann, L. Kienle, S. Zander, S. Doyle, W. Bensch, Chem. Mater. 28 (2016) 434. [10] L. Wang, D.C. Bock, J. Li, E.A. Stach, A.C. Marschilok, K.J. Takeuchi, E.S. Takeuchi, ACS Appl. Mater. Interfaces 10 (2018) 8770. [11] L.R. Hou, H. Hua, L. Lian, H. Cao, S.Q. Zhu, C.Z. Yuan, Chem. Eur J. 21 (2015) 13012. [12] C.Z. Yuan, H. Cao, S.Q. Zhu, H. Hua, L.R. Hou, J. Mater. Chem. A 3 (2015) 20389. [13] Y. Ma, Y.J. Ma, D. Geiger, U. Kaiser, H. Zhang, G. TaeKim, T. Diemant, R.J. Behm, A. Varzi, S. Passerini, Nanomater. Energy 42 (2017) 341. [14] Y.N. NuLi, Y.Q. Chu, Q.Z. Qin, J. Electrochem. Soc. 151 (2004) A1077. [15] F. Martinez-Julian, A. Guerrero, M. Haro, J. Bisquert, D. Bresser, E. Paillard, S. Passerini, G. Garcia-Belmonte, J. Phys. Chem. C 118 (2014) 6069. [16] K.T. Lee, X.F. Chuah, Y.C. Cheng, S.Y. Lu, J. Mater. Chem. A 36 (2015) 18578. [17] M. Yu, Y. Huang, K. Wang, X. Han, M. Wang, Y. Zhu, L. Liu, Appl. Surf. Sci. 462 (2018) 955. [18] L.R. Hou, R.Q. Bao, Y.R. Zhang, X. Sun, J.Y. Zhang, H. Dou, X.G. Zhang, C.Z. Yuan, J. Mater. Chem. A 6 (2018) 17947. [19] Z. Xing, Z.C. Ju, J. Yang, H.Y. Xu, Y.T. Qian, Nano Res 5 (2012) 77. [20] J.N. Wang, G.R. Yang, L. Wang, W. Yan, Electrochim. Acta 222 (2016) 1176. [21] Z.L. Wang, Z.W. Zhao, Y.R. Zhang, X.P. Yang, X. Sun, J.Y. Zhang, L.R. Hou, C.Z. Yuan, Chem. Eur J. 25 (2019) 863. [22] H. Qiao, Y.Q. Fei, K. Chen, R.R. Cui, Q.F. Wei, Ionics 22 (2016) 967. [23] H. Qiao, Z.K. Xia, Y.Q. Fei, L.Y. Cai, R.R. Cui, Y.B. Cai, Q.F. Wei, Q.X. Yao, Ceram. Int. 43 (2017) 2136. [24] L.C. Wu, T. Wu, M.L. Mao, M. Zhang, T.H. Wang, Electrochim. Acta 194 (2016) 357. [25] G.X. Gao, L. Shi, S.Y. Lu, T. Gao, Z.Y. Li, Y.Y. Gao, S.J. Ding, Dalton Trans. 47 (2018) 3521. [26] H.Y. Yue, T. Du, Q.X. Wang, Z.P. Shi, H.Y. Dong, Z.X. Cao, Y. Qiao, Y.H. Yin, R.M. Xing, S.T. Yang, ACS Omega 3 (2018) 2699. [27] R.C. Jin, Y.F. Meng, Y.Q. Ma, H.H. Li, Y.X. Sun, G. Chen, Electrochim. Acta 209 (2016) 163. [28] X.R. Sun, H.W. Zhang, L. Zhou, X.D. Huang, C.Z. Yu, Small 12 (2016) 3732. [29] X.B. Zhong, X.X. Wang, H.Y. Wang, Z.Z. Yang, Y.X. Jiang, J.F. Li, Z.Q. Tian, Nano Res 11 (2018) 3814. [30] Z.W. Wang, D. Schiferl, Y.S. Zhao, H.S.C. Neill, J. Phys. Chem. Solids 64 (2003)

2517. [31] A.A. Nada, M. Nasr, R. Viter, P. Miele, S. Roualdes, M. Bechelany, J. Phys. Chem. C 121 (2017) 24669. [32] L. Cabo-Fernandez, F. Mueller, S. Passerini, L.J. Hardwick, Chem. Commun. 52 (2016) 3970. [33] Y. Hou, L. Zhang, L.Y. Chen, P. Liu, A. Hirata, M.W. Chen, Phys. Chem. Chem. Phys. 16 (2014) 3523. [34] D.J. Kim, J.W. Yoon, C.S. Lee, Y.S. Bae, J.H. Kim, Appl. Surf. Sci. 439 (2018) 833. [35] S.U. Muhamad, N. HayatiIdris, H.M. Yusoff, M.F. MdDin, S.R. Majid, Electrochim. Acta 249 (2017) 9. [36] Y.B. Li, R. Yi, A.G. Yan, L.W. Deng, K.C. Zhou, X.H. Liu, Solid State Sci. 11 (2009) 1319. [37] L.H. Zhang, T. Wei, J.M. Yue, L.Z. Sheng, Z.M. Jiang, D. Yang, L.B. Yuan, Z.J. Fan, J. Mater. Chem. A 5 (2017) 11188. [38] C.Z. Yuan, J.Y. Li, L.R. Hou, X.G. Zhang, L.F. Shen, X.W. Lou, Adv. Funct. Mater. 22 (2012) 4592. , M. Mi [39] J. Taba ciarova cusík, P. Fedorko, M. Omastov a, Polym. Degrad. Stabil. 120 (2015) 392. [40] S.G. Wang, Z.T. Cui, J.W. Qin, M.H. Cao, Nano Res 9 (2016) 2270. [41] X.W. Guo, X. Lu, X.P. Fang, Y. Mao, Z.X. Wang, L.Q. Chen, X.X. Xu, H. Yang, Y.N. Liu, Electrochem. Commun. 12 (2010) 847. [42] P.F. Teh, Y. Sharma, S.S. Pramana, M. Srinivasan, J. Mater. Chem. 21 (2011) 14999. [43] Y. Sharma, N. Sharma, G.V.S. Rao, B.V.R. Chowdari, Electrochim. Acta 53 (2008) 2380. [44] Q.X. Wang, H.Y. Yue, T. Du, W.L. Zhang, Y. Qiao, H.Y. Dong, Y.H. Yin, S.T. Yang, RSC Adv. 6 (2016) 35110. [45] R.R. Li, S.G. Wang, W. Wang, M.H. Cao, Phys. Chem. Chem. Phys. 17 (2015) 24803. [46] D. Bresser, E. Paillard, R. Kloepsch, S. Krueger, M. Fiedler, R. Schmitz, D. Baither, M. Winter, S. Passerini, Adv. Energy Mater. 3 (2013) 513. [47] J.H. Sui, C. Zhang, D. Hong, J. Li, Q. Cheng, Z.G. Li, W. Cai, J. Mater. Chem. 22 (2012) 13674. [48] L. Liu, R. Gao, L.M. Sun, S.B. Han, D.F. Chen, Z.B. Hua, X.F. Liu, New J. Chem. 40 (2016) 3139. [49] X.Y. Yao, C.Y. Zhao, J.H. Kong, H.Q. Wu, D. Zhou, X.H. Lu, Chem. Commun. 50 (2014) 14597. [50] F. Zou, X.L. Hu, Z. Li, L. Qie, C.C. Hu, R. Zeng, Y. Jiang, Y.H. Huang, Adv. Mater. 26 (2014) 6622.