Facile plasma treated β-MnO2@C hybrids for durable cycling cathodes in aqueous Zn-ion batteries

Journal of Alloys and Compounds 827 (2020) 154273

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Facile plasma treated b-MnO2@C hybrids for durable cycling cathodes in aqueous Zn-ion batteries Wanwei Jiang a, b, Xijun Xu b, Yuxuan Liu b, Liang Tan b, Fengchen Zhou b, Zhiwei Xu a, **, Renzong Hu b, * a

State Key Laboratory of Separation Membranes and Membrane Processes, School of Textiles, Tianjin Polytechnic University, Tianjin, 300387, China Guangdong Provincial Key Laboratory of Advanced Energy Storage Materials, School of Materials Science and Engineering, South China University of Technology, Guangzhou, 510640, China

b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 23 November 2019 Received in revised form 5 February 2020 Accepted 9 February 2020 Available online xxx

Aqueous zinc-ion batteries have emerged as prospective energy storage devices to partly replace organic ion batteries due to their high safety and eco-friendliness. Providing multifold synthesis methods of cathode materials is essential for Zn-ion battery development. Here, we demonstrated a practical strategy for the large-scale fabrication of high performance b-MnO2@C hybrid cathode materials by plasma assisted milling (P-milling). After P-milling for 10 h, the porous hybrid microparticles consisted of MnO2 nanocrystallites, which combined and wrapped with the thin carbon layer derived from expanded graphite. The pores among the b-MnO2@C particles facilitated electrolyte infiltration during continuous cycling, while combining with carbon greatly enhanced the conductivity of the hybrids and helped to alleviate MnO2 dissolution. Therefore, the b-MnO2@C hybrids delivered excellent cycle stability, with a high capacity of 130 mAh g
1 for 400 cycles at a current rate of 300 mA g
1 in an aqueous Zn(CF3SO3)2 electrolyte. This capacity retention was amongst the highest reported so far for MnO2-based cathode materials for Zn-ion batteries. © 2020 Elsevier B.V. All rights reserved.

Keywords: Zinc-ion battery Manganese dioxide Plasma milling Aqueous electrolyte

1. Introduction Lithium-ion batteries have been widely applied to a variety of portable electronic devices by virtue of their high energy densities and long lifetimes [1e5]. However, the safety issues associated with flammable electrolytes and the growing concerns over price and the availability of lithium resources block their large-scale applications [6]. On the basis of these factors, multivalent-ion batteries (such as Mg2þ, Al3þ and Zn2þ) and aqueous electrolytes are receiving significant attention due to their high safety, low cost and environmental friendliness [7,8]. Furthermore, in theory, the multiple electron transfer of multivalent-ion batteries can be used to obtain high capacities that further lead to high energy densities [9]. The advantages of aqueous multivalent-ion batteries are attracting notable research interest. In particular, Zn-ion batteries have a number of unique advantages, as follows. First, the Zn metal anode

* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (Z. Xu), [email protected] (R. Hu). https://doi.org/10.1016/j.jallcom.2020.154273 0925-8388/© 2020 Elsevier B.V. All rights reserved.

possesses a high specific capacity of 819 mAh g
1. Second, rechargeable Zn-ion batteries have higher water compatibility and stability than alkaline metals. Thirdly, Zn-ion batteries can be produced and recycled with a mature industrial process based on commercial alkaline Zn batteries [10,11]. Therefore, for small electronic applications, Li-ion batteries could be replaced by mild aqueous Zn-ion batteries. Although Zn-ion batteries show significant potential, a variety of challenges remain for the cathodes in these devices [12]. Various cathode materials have been evaluated in rechargeable Zn-ion batteries, in order to find electrode materials with a high energy density as the host of Zn2þ ions and reliable cycling stability for long-term application [13e15]. For example, Prussian blue and its analogues possess an average operation voltage up to 1.7 V vs. Zn and an open skeleton structure that allows for rapid diffusion of Zn2þ ions but delivers a low specific capacity [16,17]. Vanadiumbased materials, such as Na2V6O16$1.63H2O, V2O5$nH2O and Zn0.25V2O5, are considered as potential candidates, owing to their high specific capacity and good structural stability [18e20]. However, the low operating voltage of 0.2e1.6 V vs. Zn is a limitation of their widespread application.

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In contrast, manganese dioxide (MnO2), fabricated by different methods, has attracted attention as a result of its advantages in terms of considerable theoretical capacity and higher operation voltage (with an average up to 1.7 V vs. Zn). Pan et al. demonstrated a highly reversible zinc/manganese oxide system, in which a-MnO2 nanofibers were used as the cathode material and an optimal mild aqueous ZnSO4-based solution was used as the electrolyte. The system exhibited a specific capacity of 285 mAh g
1 at a current rate of 60 mA g
1 [10]. Zhang et al. reported a high performance rechargeable zinc-manganese dioxide battery with a mild aqueous zinc triflate electrolyte, exhibiting an excellent specific capacity of 225 mAh g
1 [6]. However, it has been found that the capacity of MnO2 cathodes fades rapidly due to the gradual dissolution of MnO2 via disproportion in the electrolyte. A promising method to combat this is to induce a Mn2þ ion source into the electrolyte or use an aqueous Zn(CF3SO3)2 electrolyte, which has been shown to improve the cycle performance in some reports [7,10,21]. Carbon coatings have also been used to improve the cycling stability of MnO2 cathodes [22,23]. Wu and co-workers successfully assembled a cathode with a-MnO2 nanowires coated by graphene scrolls, achieving significant electronic conductivity and alleviating the Mn2þ dissolution at high current densities [24]. In addition, different from V2O5-based cathodes, the insertion of Zn2þ ions in MnO2 leads to a structural change from the initial phase to a layered structure [25]. Significantly more Zn2þ ions may insert into the MnO2 cathode under low current density accompanied by a large volume expansion, which usually leads to structural damage and thus inferior reaction reversibility and fast capacity fading. The porous structure and nanoscale size of MnO2 composites, which are promising solutions, can be served to effectively eliminate phase changes and facilitate charge storage [26,27]. Currently, plasma assisted milling (P-milling), which is a combination of heating and high energy electron bombardment effects produced by plasma, and milling mechanical effects have attracted extensive research interest in many fields by virtue of their unique advantages in the high efficiency promotion of powder refinement, activation and insitu chemical reaction [28e30]. However, there have been no reports on P-milling strategies to elaborate MnO2-based composites at the nanoscale for aqueous zinc-ion batteries. Herein, we report a practical P-milling strategy to create bMnO2@C nanocomposites, derived directly from pristine MnO2 and expanded graphite powders, as cathodes for aqueous Zn-ion batteries. After a short period of P-milling, the b-MnO2@C nanocomposites consisted of MnO2 nanocrystals, which were combined and wrapped with thin carbon layers. In an aqueous electrolyte solution containing 3 M Zn(CF3SO3)2 and 0.1 M MnSO4, the P-milled b-MnO2@C cathode delivered a high capacity of more than 100 mAh g
1 at current rate of 300 mA g
1, with a capacity retention of 100% after 400 cycles. Furthermore, ex-situ X-ray diffraction (XRD) and soft X-ray absorption spectroscopy (XAS) investigations revealed the reversible formation of Zn4SO4(OH)6$4H2O, which contributed to the stable cycling performance for the P-milled bMnO2@C cathode. In addition, the plasma assisted preparation of bMnO2@C was simple and facile, thereby representing a large-scale method that may open new opportunities for the design of high performance and low-cost rechargeable aqueous ZneMnO2-ion batteries.

by annealing expandable graphite that was heated at 900  C for 2 min in a muffle furnace according to high-temperature puffing mechanism. The powder mixtures with a mass ratio of 9:1 of bMnO2 and expanded graphite were treated by P-milling for 10 h in an Ar atmosphere [31]. The weight ratio of ball to powder was set at 50:1, according to our previous work [32]. In addition, to prevent samples from overheating, the milling machine was stopped for 15 min every 2 h during milling. Finally, the samples yielded were named as b-MnO2@C (9e1). For comparison, different mass ratios of b-MnO2@C composites (8e2, 7e3 and 5-5) were also produced. To research the effect of milling time on performance, the bMnO2@C (9e1) composite milled for 5 h was also investigated. 2.2. Microstructural characterization The phase structures of the b-MnO2@C samples were investigated by XRD with Cu-Ka radiation. The microstructures of the samples were characterized by field-emission scanning electron microscopy (SEM) and transmission electron microscopy (TEM) at 200 kV. The contents of Mn, C and O were measured by electron probe X-ray microanalysis (EPMA, EPMA-1600). The degree of graphitization of the expanded graphite after milling was measured using a Raman spectrometer at a 632.8 nm excitation wavelength. The X-ray photoelectron spectroscopy (XPS) analysis was performed using an Escalab 250 photoelectron spectrometer (Thermo Fisher Scientific, USA). Soft XAS at the Mn L-edge was collected in total electron yield mode on beamline BL08U at the Shanghai Synchrotron Radiation Facility. Cycled electrodes were washed several times with deionized water and absolute ethanol to remove the electrolyte and then dried in a vacuum oven before the microstructural analysis. 2.3. Electrochemical measurements The electrochemical properties of these composites were measured using coin-type half cells (CR2025). The b-MnO2@C cathodes were prepared by a slurry coating process on Ti foils with Super P and a polyvinylidene difluoride binder in a weight ratio of 8:1:1 in N-methyl-2-pyrrolidone and then dried at 80  C in a vacuum for 12 h. Metallic Zn foils were used as the anode and a glass fiber membrane was used as the separator. An aqueous solution containing 3 M Zn(CF3SO3)2 and 0.1 M MnSO4 was used as the electrolyte. For comparison, electrolytes of 2 M ZnSO4 with 0.1 or 0.2 M MnSO4 added were also prepared. Cyclic voltammetry (CV) profiles were measured on an electrochemical workstation (Interface 1000, Gamry, USA) at a voltage region between 1.0 and 1.8 V vs. Zn2þ/Zn with scan rates from 0.1 to 2 mV s
1. The galvanostatic charge-discharge tests and rate performance analysis were carried out on a battery test system (CT2001A, LAND, China). Electrochemical impedance spectroscopy (EIS) was performed on a Gamry Interface 1000 workstation in a frequency range from 100 kHz to 0.1 Hz at a 10 mV amplitude signal without applied voltage. The capacities were calculated based on the total mass of the bMnO2@C composites in this work. 3. Results and discussion 3.1. Morphological structures

2. Experimental 2.1. Preparation of b-MnO2@C composites

b-MnO2 (99.9% purity, Aladdin Industrial Inc., China) and expanded graphite powders were used as the raw materials without further purification. The expanded graphite was obtained

Fig. 1(a) illustrates the prepared procedures of the b-MnO2@C composites by a practical P-milling method. The mixtures of MnO2 and expanded graphite were treated by discharge plasma energy and mechanical milling energy synchronously. After P-milling, the coarse MnO2 particles were refined to micro-nanoparticles wrapped with carbon. Fig. 1(b) presents the XRD patterns of the as-

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Fig. 1. (a) Preparation processes of b-MnO2@C composites: (b) XRD patterns; (c) Raman spectrum of b-MnO2@C with mass ratio of 9e1 after P-milling for 10 h and un-milling expanded graphite; (d) and (e) SEM images of the as-milled b-MnO2@C (9e1) hybrid powders at different resolutions.

milled b-MnO2@C (9e1) sample in comparison with that of the unmilled mixture. Even after milling for 10 h, the diffraction peaks of b-MnO2 (JCPD00-024-0735) at (110), (101), (210) and so on can be clarified. It is noteworthy that the diffraction peaks of the expanded graphite in XRD could not be clearly found comparing to that the un-milling sample exists the diffraction peaks of C (JCPDS 00-0261079). This suggested that the absence of expanded graphite was ascribed to amorphous carbon after milling. In Fig. S1(a), different mass ratios of b-MnO2@C composites (8e2, 7e3 and 5-5) are presented, which were similar to the XRD patterns with the mass ratio of b-MnO2@C (9e1). To further confirm the existence of amorphous carbon, the bMnO2@C sample was investigated by Raman spectroscopy. As shown in Fig. 1(c) and S1(b), the D peak was ~1344 cm
1, corresponding to the defects or disorder in the carbon materials. The G peak was at 1595 cm
1 with the relationship of the E2g mode of the carbon sp2 atoms. The intensity of I D/I G (0.63) changed into 1.35 after P-milling indicating that the existence of expanded graphite derived from the expanded graphite during milling [33,34]. Fig. 1(d) shows a typical SEM image of the near spherical morphology of the as-milled b-MnO2@C (9e1) sample. The size of the as-milled spherical particles was in the range of 5e10 mm. The

magnified image in Fig. 1(e) further reveals that the b-MnO2@C composites were actually an aggregation of many irregular grains, with sizes ranging between 100 and 500 nm, even for the mass ratios of b-MnO2@C (8e2, 7e3 and 5-5) composites shown as a similar morphology with b-MnO2@C (9e1) in Fig. S1(c)e(e). However, among these composites, there were numerous multiscaled pores and the MnO2 and carbon could be not identified, which should be due to the homogenous mixing of these two phases after high energy P-milling for 10 h. This could be confirmed by the elemental mappings of the b-MnO2@C sample at different resolutions. As shown in Table S1 and Fig. S2, at different resolutions, the b-MnO2@C (9e1) composites have almost the same mass ratio of Mn, O and C elements, consolidating the highly dispersed state of b-MnO2 and carbon in the as-milled composite. Fig. 2(a) shows a typical TEM image of the P-milled b-MnO2@C (9e1) composite, showing that the MnO2 nanoparticles were well veiled in the carbon matrix. The existence of MnO2 was confirmed by the selected-area electron diffraction patterns, as shown in Fig. 2(b). Furthermore, the elemental mappings of Mn and C were obtained for the particles given in Fig. 2(a), and as shown in Fig. 2(c), the Mn was distributed uniformly on the carbon matrix in the P-milled composite. The high-resolution TEM image in Fig. 2(d)

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Fig. 2. (a) TEM image of b-MnO2@C. (b) Selected-area electron diffraction patterns. (c) Elemental mapping of Mn, O and C and (def) high-resolution TEM images.

reveals that a crystalline MnO2 particle was well enclosed by the carbon veil with a thickness of less than 10 nm. According to the lattice fringes in the high-resolution images shown in Fig. 2(e) and (f), the interlayer spacing of b-MnO2 was distinguished and marked. The lattice fringe distances of 0.884, 1.62 and 2.11 Å were consistent with the (332), (211), and (110) lattice spaces of tetragonal MnO2, respectively. The TEM image of Fig. 2(f) also clearly reveals the ultrafine grains of MnO2. Amongst these small grains with a size of ~5 nm, there was a large number of interfaces, which could serve as transport channels for the electrolyte during electrochemical reactions. 3.2. Electrochemical performance To evaluate the electrochemical properties of the resulting composite, Fig. 3(a) shows the CV curves of the b-MnO2@C (9e1) composite electrode carried out in an aqueous electrolyte of 3 M Zn(CF3SO3)2 and 0.1 M MnSO4. In the first cathode scan, a sharp peak at ~1.14 V was observed. Meanwhile, a reversible peak at ~1.56 V appeared in the first anode scan. In addition, in the subsequent scan, cathode peaks at 1.17 and 1.4 V and anode peaks at 1.56 and 1.63 V were clearly observed and repeated, respectively, indicating the existence of the activation process and the phase transition in the second cycles [6,35]. The each discharge peak at around 1.40 V indicates that Zn2þ intercalation occurs during the first plateau and Hþ intercalation may occur within the first plateau simultaneously with Zn2þ intercalation. And the second discharge peak almost at 1.20 V is the formation of Zn4SO4(OH)6. In the charge potential of 1.56 V, most of the Zn4SO4(OH)6 outcomes have dissolved at the first charge point, and at the charge plateau of 1.63 V, the Zn4SO4(OH)6 outcomes have completely dissolved. Meanwhile, the main amount of Zn2þ is extracted [36]. Fig. 3(b) shows the galvanostatic capacity vs. voltage profiles of the b-MnO2@C (9e1) cathode at a current rate of 50 mA g
1 between 1.0 and 1.8 V vs. Zn/ Zn2þ, with platforms at 1.4 and 1.2 V in the discharge process and

1.54 V in the charge process, which was in accordance with the CV profiles. Moreover, as shown in Fig. 3(b), the specific capacities of the discharge/charge were gradually increased, which may be attributed to the following reasons. First, it was related to the gradual activation of MnO2 in the electrodes. For instance, the electrolyte gradually soaked into the inner part of the cathode materials, which was a universal phenomenon in aqueous batteries and capacitors [22,24,37]. As the EIS results show in Fig. 3(c), the gradual decrease in the ohmic resistance and charge transfer resistances of the b-MnO2@C electrode with cycling could also reflect the gradual activation of MnO2 in the electrode. The apparent Znion diffusion coefficient (D2þ Zn ) also exhibited a noticeable similar effect. The values of D2þ Zn were calculated from the inclined lines in the Warburg region [38,39]: 2 2 2 4 4 2 2 D2þ Zn ¼ R T /2A n F C s

(1)

2where A is the surface area of the electrode, n is the number of the electrons per molecule attending the electronic transfer reaction, F is the Faraday constant (96,500 C mol
1), C is the concentration of Zn ion in electrode, R is the gas constant (8.314 J K
1 mol
1), T is the room temperature in our experiment, s is the Warburg factor associated with
Z” (
Z”fsu
1/2), respectively [40]. After the linear fitting, the relation plot between
Z00 and the reciprocal square root of the angular frequency u shown in the inside image of Fig. 3(c) is used to assess the Warburg factor s and the apparent Zn2þ ion diffusion coefficient [41]. The D2þ Zn values of the b-MnO2@C cathode at the 1st, 2nd, 5th, 10th and 30th cycles are calculated to be 3.56  10
16, 6.75  10
16, 1.26  10
15 and 1.80  10
15 cm2 s
1, respectively. Second, the increasing specific capacity may also be due to the tunnelstructured MnO2 undergoing a phase transition allowing for subsequent intercalation of Zn2þ ions into the latter structure during the initial cycles [24]. Eventually, with an electrolyte of 3 M Zn(CF3SO3)2 and 0.1 M MnSO4, the b-MnO2@C electrode could reach a high capacity of 250 mAh g
1 at a low current rate of 50 mA g
1.

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Fig. 3. Electrochemical performance of b-MnO2@C (9e1) between 1.0 and 1.8 V. (a) CV curves for the initial three cycles with a scanning rate of 0.1 mV s
1. (b) Galvanostatic voltage profiles for the 1st, 5th, 10th, 20th, 30th, 40th and 50th cycles at a current rate of 50 mA g
1. (c) EIS test for the 1st to 5th, 10th and 30th cycles, with the inset showing the relationship between -Z00 and the reciprocal square root of the frequency in the low-frequency region. (d) Cycle performance and Coulombic efficiency at a current rate of 200 mA g
1. (e) Rate performance at current rates from 100 mA g
1 to 2 A g
1.

Fig. 3(d) shows the cycle performance of the b-MnO2@C (9e1) cathode at a current rate of 200 mA g
1. After the initial activation cycles, the specific capacity reached almost 150 mAh g
1 after 250 cycles. The Coulombic efficiency was also nearly 100% after the initial cycles. It is worth noting that the stable cycle performance of

the b-MnO2@C cathode should be attributed to the ultrafine grain size of the MnO2 and the high dispersion state of MnO2 in the amorphous carbon, which helped to make an effective conductive network and inhibited MnO2 dissolution to some extent. Fig. 3(e) exhibits the rate performance of the b-MnO2@C cathode at

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increasing current rates from 100 mA g
1 to 2 A g
1. At high current rate of 2A g
1, the cathode delivered a reversible capacity of 40 mh g
1 after 50 cycles, comparable to the reversible capacity of the Prussian blue cathode cycled at very low current rates [16,17]. After that, as the cathode was continuously cycled at a low current rate of 300 mA g
1, a stable capacity of 100 mAh g
1 could be maintained even after 300 cycles. These results consolidated the observation that the as-milled b-MnO2@C had a strong tolerance for Zn2þ intercalation/deintercalation and was endowed with good rate performance and excellent cycling stability. In Fig. S3(a)e(c), the charge and discharge curves of the different mass ratios of b-MnO2@C (9e1, 8e2 and 5-5) electrodes are shown, which exhibited a small rise during cycling, as mentioned above (such as Fig. 3). Fig. S3(d) further shows the b-MnO2@C (9e1, 8e2, 7e3 and 5-5) electrode performances, which illustrated that the electrode performances were relative to the MnO2 content in composites and the appropriate carbon coated like many reported papers [24,42]. Therefore, the b-MnO2@C (9e1) composite possessed the highest performance among them. The b-MnO2@C (9e1 and 8e2) electrodes exhibited increasing specific capacity values, reaching 120 and 90 mAh g
1, respectively. However, the bMnO2@C (7e3 and 5-5) electrodes obtained specific capacities of 82 and 56 mAh g
1, which exhibited a slight reduction. Furthermore, for comparison, the effect of milling time on performance for the bMnO2@C (9e1) composite with milling for 5 h is presented in Fig. S3(e), which also showed long-life cycling, with 79 mAh g
1 obtained at around 165 cycles at a current rate of 300 mA g
1 but

less than that one milling with 10 h. Evidently, under the same composition, the milling time had an influence on electrode performance, owing to the small grain size of the b-MnO2@C composite. This suggested that P-milling was conducive to stabilizing cycling for Zn-ion batteries in our work. Remarkably, the aqueous electrolytes may also have influences on the stability of the P-milling b-MnO2@C cathode materials. As the CV curves show in Fig. 4(a) and (b), the cathodes of b-MnO2@C (9e1) tested with electrolytes of 2 M ZnSO4 with 0.1 or 0.2 M MnSO4 had similar oxide and reduction peaks with those tested in the electrolyte of 3 M Zn(CF3SO3)2 and 0.1 M MnSO4. Moreover, as shown in Fig. S4(a) and (b), the cathodes tested in these three kinds of electrolytes presented very similar capacity-voltage profiles, suggesting that the electrolyte selection did not change the reaction mechanism of b-MnO2@C toward Zn2þ ion storage. In Fig. 4(c), the cathodes with Zn4SO4eMnSO4 had higher specific capacities than the cathode with Zn(CF3SO3)2eMnSO4. Nevertheless, the latter one enabled the b-MnO2@C cathode having superior cycling stability, which achieves almost 130 mAh g
1 after 400 cycles at a current rate of 300 mA g
1. This should be due to the fact that the MnO2 dissolution is alleviated in the Zn(CF3SO3)2eMnSO4 electrolyte [43]. Fig. 4(d) and Table S2 show the comparison of capacity retention of the simple P-milled b-MnO2@C cathode with various manganese oxide cathodes reported in recent studies. Table S3 gives a specific capacity comparison of various manganese oxide cathodes, displaying that the P-milled b-MnO2@C cathode possesses superior cycle stability and excellent cycle performance even

Fig. 4. (a) and (b) CV curves of b-MnO2@C (9e1) cathodes with 2 M ZnSO4-0.1 and 0.2 M MnSO4. (c) Cycle performance of b-MnO2@C (9e1) with 2 M ZnSO4-0.1 M MnSO4, 2 M ZnSO4-0.2 M MnSO4 and 3 M Zn(CF3SO3)2-0.1 M MnSO4. (d) Capacity retention of P-milled b-MnO2@C (9e1) in comparison with different manganese oxide cathodes in recent studies.

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after 400 cycles. These results fully demonstrated the feasibility of our strategy of using P-milled b-MnO2@C and then using a Zn(CF3SO3)2eMnSO4 electrolyte to significantly enhance the cycling stability of MnO2-based cathodes for rechargeable aqueous Zn-ion batteries. The variation in peak shape of the CV curves with scan rate reflected the kinetics of metal ion intercalation/deintercalation at the electrode/electrolyte interface and/or the rate of ion diffusion inside the electrode. As shown in Fig. 5(a), a series of CV curves were recorded with the b-MnO2@C electrode as a function of scan rate in the range from 0.1 to 2 mV s
1, during which the reduction or oxidation peaks shifted with increased scan rates. The CV profiles were followed by the relationship between the current (i) and the scan rate (v): i ¼ avb

(2)

where a and b are adjustable parameters. The b value is used to provide kinetic information regarding the electrochemical system in batteries [44,45]. A b value of 0.5 means semi-infinite diffusioncontrolled faradaic processes while a b value of 1 represents ideal capacitive-controlled behavior [46]. As shown in Fig. 5(b), the bMnO2@C cathode had b values for oxidation and reduction of 0.52

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and 0.70, respectively. This indicated that both diffusion-controlled faradaic processes and capacitive-controlled behavior had a synergetic effect in the electrochemical performance of the b-MnO2@C cathode. Furthermore, the capacity contribution of diffusioncontrolled (k1v1/2) and capacitive-controlled (k2v) fractions could be quantified according to the following equation: i(V) ¼ k1v1/2 þ k2v

(3)

By transforming Eq. (3) to i(V)/v1/2 ¼ k1 þ k2v1/2

(4)

k1 and k2 can be facilely achieved by plotting i(V)/v1/2 vs. v1/2 and thus the capacity of the capacitive-controlled fraction can be calculated by ic(V) ¼ k2v [47,48]. For instance, at a given scan rate of 0.8 mV s
1, the CV profile for the capacitive contribution compared with that of the total measured current is shown in Fig. 5(c). Based on the same method, the capacitive contribution of other scan rates from 0.5 to 2 mV s
1 were determined. As shown in Fig. 5(d), the capacitive contributions were 46.5%, 53.0%, 57.6% and 67.5% at scan rates of 0.5, 0.8, 1 and 2 mV s
1, respectively, which gradually increased with the scan rate raised. The capacitive contribution in

Fig. 5. (a) CV profiles with scan rates from 0.1 to 2 mV s
1. (b) Log (i) versus log (v) plots of the cathodic current response at two peaks shown in (a). The slopes of this line determine b values. (c) Separation of capacitive and diffusion currents at a scan rate of 0.8 mV s
1. The capacitive contribution to the total current is shown by the shaded region. (d) Contributions of capacitive portion at different scan rates from 0.5 to 2 mV s
1.

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this P-milled b-MnO2@C was somewhat lower than those of nanostructured MnO2-based cathodes prepared by other chemical methods, which was mainly attributed to the larger grain size of the P-milled MnO2@C composite [22]. 3.3. Mechanistic investigation To probe the structural evolution of the b-MnO2@C (9e1) cathodes cycled with the Zn(CF3SO3)2-0.1 M MnSO4 aqueous

electrolyte, the cathodes at different discharge and charge states were investigated by XRD, XPS and soft XAS. As the XRD patterns recorded in Fig. 6(a) show, at the end of discharging to 1.0 V in the second cycle, the diffraction peaks of the MnO2 disappear, instead the obvious characteristics peaks in the XRD patterns were assigned to Zn4SO4(OH)6$4H2O (JCPDS 044e0673). Even at the 10th and 50th cycles, the precipitates of Zn4SO4(OH)6$4H2O were confirmed at the end of discharge (Fig. S5). No crystalline phase with MnO2 could be determined. However, as the cathode

Fig. 6. (a) XRD patterns of b-MnO2 (9e1) electrode at full charge and discharge in the second cycle with 3 M Zn(CF3SO3)2 and 0.1 M MnSO4 aqueous electrolyte. (b) XRD patterns of b-MnO2@C (9e1) electrode after 2nd, 10th, 50th and 200th charging to 1.8 V. (c) XPS characterization of 2nd and 200th cycles. (d) Soft XAS at Mn L-edge of b-MnO2@C (9e1) electrode after 2nd, 50th and 200th charging to 1.8 V or discharging to 1.0 V, respectively. (e) Schematic showing the reactions during the discharge process for b-MnO2/Zn cell employing the aqueous 3 M Zn(CF3SO3)2-0.1 M MnSO4 electrolyte.

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recharged to 1.8 V in the second cycle, the diffraction peaks of MnO2 appeared, which meant that the consequent formation of Zn4SO4 (OH)6$4H2O on the cathode was reversible. In the second cycle with charging to 1.8 V, the XRD diffraction peaks of MnO2 became wider than the original electrode, suggesting that regenerated MnO2 had a smaller grain size and it is conductive to further embedding Zn2þ ions. Fig. 6(b) shows the XRD patterns of the b-MnO2@C (9e1) cathode fully charged to 1.8 V at the 2nd, 10th, 50th and 200th cycles. At the 2nd and 10th cycles, the peaks of the XRD patterns are assigned to b-MnO2. The peaks of new Zn-insertion phases as typical of ZnMn2O4 (JCPDS 01-071-2499) and MnO2 occurred explicitly at the 50th cycle, while the peaks of b-MnO2 almost disappeared in the 200th cycle, meaning that the Zn2þ ions could insert and extract in and out of the inner MnO2 crystals. As shown in Fig. 6(c), the XPS surveys of the b-MnO2@C cathodes discharging to 1.0 V and charging to 1.8 V showed the dual peaks of Mn 3s. The energy splitting (DE) of the Mn 3s doublet peaks was 4.71, 5.3, and 5.18 eV for the 2nd cycle at 1.0 V (1.8 V) and the 200th cycle at 1.0 V, respectively, indicating that the valence of MnO2 cathode also reduces accompanied by Zn2þ ions inserting into the small grain bMnO2 at the second cycle at 1.0 V and the 200th cycle at 1.8 V. To precisely investigate the variable valence of Mn, as shown in Fig. 6(d), soft XAS at the Mn L-edge of the b-MnO2@C electrode was, respectively, identified after the discharge to 1.0 V and charge to 1.8 V in the 2nd, and charge to 1.8V in the 50th and 200th cycle, which was sensitive in probing the Mn valence near the surface [49e51]. At the second cycle of discharging to 1.0 V, the peak shift to left side, which suggested that the valence of Mn was reduced. Comparing to the second cycle of charging to 1.8 V, the peaks represented mainly Mn4þ consistent with the XRD phase. This meant that Zn ions inserted into the small grains of b-MnO2 at the second cycle of discharging to 1.0 V, consistent with the XRD results in Fig. 6(b) [27,52,53]. In the curves of the 50th and 200th cycles of charging to 1.8 V, the value of Mn mainly stayed at Mn4þ and had a little shift to the left especially at the 200th cycle. This suggested that the b-MnO2@C cathodes could further insert Zn2þ ions and were contributed to remain a stable structure even though having small changes after cycling 200 times. In terms of the above XRD, XPS, soft XAS results, the reactions of the discharge process for the MnO2@C/Zn cell could be illustrated in Fig. 6(e). This was attributed to the synergistic effect of the porous structure of P-milled b-MnO2@C and the electrolyte. A certain amount of Zn2þ ions could enter the small grains of the tunnelstructured MnO2 accompanied by the formation of Zn4SO4(OH)6$4H2O precipitates in holes that alleviated the stress of massive Zn2þ ion insertion, thus protecting the structure integrity of b-MnO2 to some extent [50]. Meanwhile, the Mn2þ dissolution was alleviated by virtue of carbon matrix veils and the electrolyte containing Zn(CF3SO3)2. Therefore, the MnO2@C/Zn cell could achieve stable electrochemical performance. Therefore, this might be a creative strategy for prolonging the cycle life of Zn-ion batteries by constructing enough spaces in the cathode materials to accommodate the reversible Zn4SO4(OH)6$4H2O. Nevertheless, at present, the detailed phase transformations and structural changes during the Zn-insertion/extraction process in MnO2 based cathodes were still complex and in-situ characterization was limited. Therefore, the Zn-insertion/extraction processes still requires further research [27]. 4. Conclusion In summary, we have reported a practical plasma assisted milling strategy to prepare b-MnO2@C nanocomposites, derived directly from pristine MnO2 and expanded graphite powders, as

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cathodes for aqueous Zn-ion batteries. The P-milled b-MnO2@C cathode delivered high capacity greater than 100 mAh g
1 at a current rate of 300 mA g
1, retaining a capacity retention of 100% after 400 cycles, which is amongst the best capacity retention values reported so far for MnO2-based cathode materials. This outstanding cycling stability was mainly attributed to the synergistic effects of the porous structure of the P-milled b-MnO2@C with nanoscale size and the Zn(CF3SO3)2-based electrolyte. The electrolyte of 3 M Zn(CF3SO3)2-0.1 M MnSO4 was good for bMnO2@C stability due to inhibiting Mn2þ dissolution. In particular, the b-MnO2@C composite packed with micro/nanoparticles owed enough spaces for Zn2þ ion, which could also facilitate Zn4SO4 (OH)6$4H2O formation thus against the stress of massive of Zn2þ ions insertion to small grains of MnO2. Owing that, it might be a creative strategy for prolonging the cycle life of Zn-ion batteries. Finally, it is worth mentioning that the plasma assisted preparation of b-MnO2@C cathode materials was simple and facile, which has potential for large-scale applications and may pave a practical way for the application of mild aqueous ZneMnO2 ion batteries. Credit author statement We have complied with Journal of alloy and compound’s ethical requirements: We hereby certify that the work has not been submitted previously to the Journal of alloy and compound (in part or in whole), which it is original and has not been submitted and published previously in any journal, and is not under consideration for publication elsewhere. Its publication is approved by all authors and tacitly or explicitly by the responsible authorities where the work was carried out. If accepted, it will not be published elsewhere in the same form, in English or in any other language, without the written consent of the Publisher. Thank you very much for considering our manuscript for potential publication. We are looking forward to hearing from you soon. Acknowledgements This work was supported by the National Natural Science Foundation of China (Nos. 51671088, 51822104, 51831009, and 11575126); the Guangzhou Science and Technology Plan Projects (No. 201904020018); and the Fundamental Research Funds for the Central Universities, SCUT (No. 2019CG24). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.jallcom.2020.154273. References [1] J. Baxter, Z. Bian, G. Chen, D. Danielson, M.S. Dresselhaus, A.G. Fedorov, T.S. Fisher, C.W. Jones, E. Maginn, U. Kortshagen, Energy Environ. Sci. 2 (2009) 559e588. [2] N. Armaroli, V. Balzani, Energy Environ. Sci. 4 (2011) 3193e3222. [3] H. Li, Z. Wang, L. Chen, X. Huang, Adv. Mater. 21 (2009) 4593e4607. [4] R. Hu, Y. Ouyang, T. Liang, X. Tang, B. Yuan, J. Liu, L. Zhang, L. Yang, M. Zhu, Energy Environ. Sci. 10 (2017) 2017e2029. [5] W. Jiang, H. Wang, Z. Xu, N. Li, C. Chen, C. Li, J. Li, H. Lv, L. Kuang, X. Tian, Chem. Eng. J. 335 (2018) 954e969. [6] N. Zhang, F. Cheng, J. Liu, L. Wang, X. Long, X. Liu, F. Li, J. Chen, Nat. Commun. 8 (2017) 405. [7] W. Sun, F. Wang, S. Hou, C. Yang, X. Fan, Z. Ma, T. Gao, F. Han, R. Hu, M. Zhu, J. Am. Chem. Soc. 139 (2017) 9775. [8] Y. Zhang, N. Liu, Chem. Mater. 29 (2017) 9589e9604. [9] K. Zhang, X. Han, Z. Hu, X. Zhang, Z. Tao, J. Chen, Chem. Soc. Rev. 44 (2015) 699. [10] C. Wang, H. Pan, J. Yang, J. Liu, K.T. Mueller, K.S. Han, P. Yan, P. Bhattacharya, X. Li, Y. Cheng, Nat. Energy 1 (2016) 16039.

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