Vacuum induced self-assembling nanoporous LiMn2O4 for lithium ion batteries with superior high rate capability

Accepted Manuscript Title: Vacuum induced self-assembling nanoporous LiMn2 O4 for lithium ion batteries with superior high rate capability Author: Wei-Bo Hua Su-Ning Wang Xiao-Dong Guo Shu-Lei Chou Kui Yin Ben-He Zhong Shi-Xue Dou PII: DOI: Reference:

S0013-4686(15)30674-5 http://dx.doi.org/doi:10.1016/j.electacta.2015.10.093 EA 25896

To appear in:

Electrochimica Acta

Received date: Revised date: Accepted date:

1-9-2015 29-9-2015 18-10-2015

Please cite this article as: Wei-Bo Hua, Su-Ning Wang, Xiao-Dong Guo, Shu-Lei Chou, Kui Yin, Ben-He Zhong, Shi-Xue Dou, Vacuum induced self-assembling nanoporous LiMn2O4 for lithium ion batteries with superior high rate capability, Electrochimica Acta http://dx.doi.org/10.1016/j.electacta.2015.10.093 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Vacuum

induced

self-assembling

nanoporous

LiMn2O4 for lithium ion batteries with superior high rate capability Wei-Bo Huaa, Su-Ning Wanga, Xiao-Dong Guoa,b,*, Shu-Lei Choub,*, Kui Yina,c, Ben-He Zhonga, Shi-Xue Doub

a

College of Chemical Engineering, Sichuan University, No.24 South Section 1, Yihuan

Road, Chengdu, 610065, China. b

Institute for Superconducting and Electronic Materials, University of Wollongong,

Wollongong NSW 2522, Australia. c

Southwest Company, China Petroleum Engineering Co., Ltd., No. 6, Shenghua Road,

Chengdu, 610041, China.

---------------------------*Corresponding author: Xiao-Dong Guo, College of Chemical Engineering, Sichuan University, No.24 South Section 1, Yihuan Road, Chengdu, 610065, China. E-mail: [email protected]. Shu-Lei Chou, Institute for Superconducting and Electronic Materials, University of Wollongong, Wollongong NSW 2522, Australia. E-mail: [email protected].

Graphical abstract

A nanoporous spinel LiMn2O4 cathode material with excellent electrochemical property has been successfully prepared via a simple, low-cost and scalable vacuum induced self-assembling process.

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Highlights:   

Nanoporous LiMn2O4 is prepared by vacuum induced self-assembly reaction. Ammonia molecules play a key role in the formation of the crystals. Nanoporous structure of LMO-A improves its electrochemical properties.

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Abstract Spinel LiMn2O4 is an inexpensive, eco-friendly and highly abundant cathode material for lithium ion batteries. Here, we report a synthesis of nanoporous LiMn2O4 cathode material using a simple vacuum induced self-assembly reaction. Ammonia molecules play a key role in the formation of the nanoporous structure in our method. The galvanostatic charge/discharge results show that the nanoporous LiMn2O4 delivers a high specific capacity at high power rates. About 95.9 % of its initial capacity (94.5 mAh g-1) is retained after 100 cycles at 10 C. The enhanced kinetics of nanoporous LiMn2O4 with low apparent activation energies indicates that the nanoporous structure provides short Li-ion diffusion paths and a continuous three-dimensional network of pathways for the transport of Li-ions and electrons. These results reveal that the nanoporous spinel LiMn2O4 material is a promising cathode candidate for next generation of high-power lithium ion battery.

Keywords: lithium ion batteries; cathode material; nanoporous; vacuum induced self-assembly; rate capability

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1. Introduction Rechargeable lithium ion batteries (LIBs) have played a very important role in today’s portable electronics, and they are considered as exciting power sources for hybrid electric vehicles (HEVs) and full electric vehicles (EVs) [1-3]. These batteries are based on electrode reactions via a typical intercalation reaction for which Li-ions are inserted/extracted into/from an open host structure with an accompanying addition/removal of electrons. Although such batteries can provide higher energy density (Wh kg-1) than other secondary electrochemical energy storage systems (i.e. double layer capacitors and pseudocapacitors), the poor high power capability is still one of the great challenge for LIBs for use in EV applications [4, 5]. Spinel LiMn2O4 is a promising candidate to replace layered Ni or Co oxides as cathode material in LIBs because of its lower cost, more abundance, better environmental friendliness, and better safety [6, 7]. The high rate capability of LiMn2O4 is hindered by several issues. First, the intrinsic low lithium-ion diffusion coefficient in the solid phase (10-8-10-11 cm2 s-1) prevents the improvement of high rate capability, because only a fraction of the bulk can be available for Li-ions insertion/extraction at high charge–discharge rates [8]. Second, the instability of the deep delithiated phase of LiMn2O4 in organic electrolyte results in serious capacity fading due to the slow dissolution of manganese ions through a disproportionate reaction ( 2 M n

3

  Mn

2

 Mn

4

)

and the phase transformation from cubic to tetragonal phase [9, 10]. Coating or doping is an efficient approach to alleviate the dissolution of manganese [6, 11]. To address the poor rate capability of bulk LiMn2O4, preparing nanosized material is usually suggested to enhance the power capability of intercalation electrode materials because of the resultant short Li-ion and electron diffusion lengths [12]. However, nanosized cathode material with a high surface area could cause many undesired side reactions [13].

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To obtain highly conductive LiMn2O4 and superior electrochemical performance, considerable efforts have been put into the research on synthesis methods, because the synthesis route was found to play a great impact on the structure, morphology, particle size, crystallinity and porosity of LiMn2O4, which have a pronounced effect on its Li-ion conductivity and electrochemical performance [14-18]. For example, Lee et al. [19] synthesized ultrathin LiMn2O4 nanowires using ɑ-MnO2 nanowires as a self-sacrificial template, and the resultant nanowires presented excellent high rate capabilities. Jiang et al. [20] reported a synthesis of hollow LiMn2O4 nanocones by a template-based reaction which delivered the enhanced electrochemical properties. Zhu et al. [21] also reported the nanoporous LiMn2O4 hollow structure based on the CaCO3-template synthesis method, which showed superior rate capability and good cycling stability. These methods were found to be effective for preparation of LiMn2O4 cathode materials with small particle size. Despite these successes, their preparation technologies are complicated and time-consuming. Therefore, it is really a challenge to develop a facile and scalable method for preparing high power capable LiMn2O4 cathode material. Herein, a simple, environmentally benign and surfactant-free chemical approach, namely, vacuum induced self-assembling process, is proposed for synthesizing LiMn2O4 cathode material for the first time. In this method, ammonia controls both the crystal nucleation and crystal growth. According to the references [22-24], ammonia molecules could passivate the surface and thereby affect the growth and aggregation behaviors of the crystals. As a consequence, when the reaction starts with high ammonia concentration, during the ammonia evaporation process, milky Mn(NH3)n2+ complexes and a few Mn(OH)2 nuclei are formed in the solution. When the NH3 molecules of Mn(NH3)n2+ are continuously liberated, the primary particles tend to become attached to each other, producing the nanoporous Mn(OH)2 precursor. Subsequently, lithiation

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reaction was done to obtain the final nanoporous LiMn2O4 material at high temperature. The asprepared LiMn2O4 material exhibit outstanding electrochemical properties and especially superior rate capability. In the present investigation, we also synthesized LiMn2O4 by a traditional solid-state route for comparison.

2. Experimental 2.1 Preparation of LiMn2O4 All chemical reagents were of analytical purity and used without further purification. The procedure used for the synthesis of the nanoporous LiMn2O4 (LMO-A) is shown in Fig. 1. Vacuum evaporation technology was introduced into the vacuum induced self-assembly reaction to prevent oxidation of Mn(OH)2 and evaporate the volatile species. The reactor pressure was typically controlled to -0.09 MPa throughout the reaction (see Fig. S1 in the supporting information). 80 ml of ammonia solution (10 M) was firstly pumped into a reactor at 45 oC. Then, 0.2 mol of Mn(CH3COO)2·4H2O dissolved in 100 ml H2O, was directly added into the reactor under a suitable stirring rate of about 200 rpm. 4.45 g of LiOH·H2O (s) was fed into the reactor until the mixture became a brown sludge. The resultant brown sludge was dried and calcined for 8 h at 700 oC in air to obtain the nanoporous LMO-A. The reactions involved are in accordance with the following equations [22, 25]: Mn

2+

2

+nN H 3 2

M n (N H 3 ) n + 2 O H

M n (N H 3 ) n 

(n = 1, 2, 3, 4)

o

45 C ，  0 .0 9 M P a

      M n (O H ) 2  n N H 3 o

8 M n (O H ) 2 + 4 L iO H  3 O 2    4 L iM n 2 O 4  1 0 H 2 O 700 C

(1) (2) (3)

For comparison, we also used a traditional solid-state method [26] to prepare LiMn2O4 (LMO-B) as a reference sample. Briefly, the MnO2 powder was thoroughly mixed with a

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stoichiometric amount of Li2CO3 by mechanical milling. The obtained mixtures were calcined at 800 oC for 10 h in air. 2.2 Characterization The crystalline structures of the samples were identified by powder X-ray diffraction (XRD, Philip PW 1730 diffractometer using a Cu Kα radiation, λ = 1.5406 Å). The morphology and particle size of the materials were observed by a field-emission scanning electron microscopy (SEM, Hitachi S4800). The microstructure of the samples was also examined by transmission electron microscopy (TEM, JEM-2100) and high resolution transmission electron microscopy (HRTEM, JEOL 2100F). Brunauer-Emmett-Teller (BET, QUADRASORB SI) N2 sorption isotherms were used to measure the specific surface area and pore size of the materials. 2.3 Electrochemical measurements The electrochemical properties of the samples were examined using two-electrode coin cells (type CR2032). The positive electrodes were assembled by coating a mixing composed of 80 wt% active material (LMO-A or LMO-B), 13 wt% electronic conductor (carbon black) and 7 wt% binder (polyvinylidence difluoride, PVDF), which was dissolved in N-methylpyrrolidinone (NMP), on an aluminium foil current collector with a diameter of 12 mm. After dried at 100 oC in a vacuum for 10 h, the electrodes were compressed at 20 MPa for 10 s between two stainless steel plates. The assembly work was performed in a high-purity argon-filled glove box. Lithium metal was used as the negative electrode. The electrolyte was 1 M lithium hexafluorophosphate (LiPF6) dissolved in a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) (1:1 v/v). A porous polypropylene (PP) film (Celgard 2300) was applied as the separator. The assembled cells were cycled at different current densities in the voltage window of 3.1–4.3 V (vs. Li/Li+) using a battery tester (BTS-3000, Neware Co., Ltd., Shenzhen, China). Electrochemical

8

impedance spectroscopy (EIS) was conducted using an electrochemical workstation (Zennium, IM6) in the frequency range of 100 kHz–10 mHz with an alternating-current (AC) amplitude of 5 mV.

3. Results and discussions The ammonia is a key factor in the fabrication of small particles and porosities for LiMn2O4 in the vacuum induced self-assembling process. In our case, there is a large amount of ammonium hydroxide in the reaction, which chelates with Mn 2+ to produce Mn(NH3)n2+ complex in the initiation of the complexing reaction [27]. The thus-formed Mn(NH3)n2+ complex prevents direct nucleation [24]. Subsequently, the evaporation of NH3 in Mn(NH3)n2+complex ions results in the nucleation of layer-structured Mn(OH)2 [22]. With the continuous loss of the NH3 molecules of Mn(NH3)n2+ during the vacuum distillation process, the continuous dissociation of the Mn(NH3)n2+ complexing compounds supplies the indispensable nutrients for the crystallization and growth of Mn(OH)2 nanoparticles (see Fig S2). The pores are also left in the intermediate Mn(OH)2 crystal in this stage [28, 29]. The XRD pattern of the typical intermediate product for LMO-A is shown in Fig S3. Main peaks of the intermediate product can be indexed to the layered Mn(OH)2 (JCPDS, No. 73-1604). In the following calcination, the Li-ions insert themselves into the intermediate structure to form the final nanoporous spinel LiMn2O4 product. The structures and phases of two samples were investigated by XRD. Fig. S4 shows the XRD patterns of the as-prepared materials. All the Bragg peaks are consistent with the spinel cubic LiMn2O4 phase (JCPDS, No. 35-0782). No other impurity peak is detected in

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the diffractogram. Sharp diffraction peaks can be observed for both samples, suggesting that both the materials formed a pure phase with high crystallinity [30]. This result also confirms that vacuum-induced self-assembly method can be used to produce high crystalline LiMn2O4 using lower temperature calcination and shorter time than traditional solid state reaction [19, 26]. The crystalline sizes for LMO-A and LMO-B are about 51 and 65 nm, respectively, which are calculated using the Scherrer equation for the (111), (311) and (400) diffraction peaks. The Rietveld refinement was successfully accomplished by assuming a cubic spinel type structure with space group Fd3m. The Rietveld refinement results on the XRD patterns for the samples are shown in Fig. 2. Good fitting with low weighted profile R-factor (Rwp) indicates the formation of wellcrystallized pure phase LiMn2O4 by the different synthesis routes. Atomic sites and coordinates x, y, z, with the number of atoms (N) for the samples are listed in Tables S1 and S2 in the supporting information. The lattice constants of LMO-A and LMO-B are 8.219 and 8.234 Å, respectively, which is in good agreement with previously reported data in the literature [18, 31]. Fig. 3 contains SEM images of the LMO-A and LMO-B powders. It can be seen from Fig. 3(a) and (b) that the secondary particles of LMO-A are composed of many interconnected primary particles with an average size about 151 nm [see Fig. S5(a)]. The primary particles have a spherical or polyhedral morphology, and these fine particles tend to aggregate together to form many void channels. In the case of LMO-B, more compact aggregation and a larger primary particles size (about 223 nm) are observed as shown in Fig. 3(c) and (d). Based on the particle size distribution of the two materials in Fig. S5, the as-synthesized LMO-A sample retains a more homogeneous distribution than the

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LMO-B sample. The N2 adsorption-desorption isotherms for both samples [Fig. 4(a)] exhibit the typical irreversible type Ⅱ shape according to the IUPAC definition. Almost no hysteresis loop in the relative pressure (P/P0) range of 0.1-0.9 is observed in the two curves, indicating few mesopores are present in the samples. The results from the BET analysis reveal that the specific surface area of LMO-A and LMO-B are 4.21 and 1.93 m2 g-1, respectively. Their corresponding pore size distribution were calculated from the nitrogen isotherms using the Barret–Joyner–Halender (BJH) method. From Fig. 4(b), we can see that LMO-A shows a larger nanopore volume of than LMO-B, especially for pore radius of about 18 nm, indicating that LMO-A really possesses a larger specific surface area than LMO-B. According to the references [31, 32], the increased specific surface area will promote undesirable reactions arising from the addition of supplementary interfaces (i.e. active material/electrolyte solution), although the specific surface area of LMO-A is lower than the values reported in the literature [19, 21, 31]. In comparison, the relatively low surface area of LMO-A with submicron particles could decrease the risk of manganese ions dissolution and contribute to the enhanced cycling performance of cathode. In brief, these results tell us that the vacuum-induced self-assembly method may be an effective way to prepare nanoarchitectured cathode material with moderate specific surface area and nano-sized pores. The microstructure of LMO-A is further investigated by TEM and HRTEM. Fig. 5(a) clearly reveals the nanoporous structure of LMO-A. As shown in Fig. 5(b), the primary particles of LMO-A have spherical or polyhedral shape and a particle size of about 150 nm, which is consistent with the analysis in SEM images. HRTEM images and the corresponding fast Fourier transform (FFT) patterns of LMO-A are presented in Fig. 5(c) and (d). The HRTEM image of

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LMO-A displays two sets of clear lattice fringes with a spacing of about 4.75 and 2.91 Å at an angle of 90o, corresponding to the (111) and ( 2 2 0 ) planes, respectively, of typical spinel-type LiMn2O4. In the FFT pattern of the LMO-A, the interplanar distances of the diffraction dots are indexed to the crystal planes ( 2 2 0 ) and ( 2 2 0 ), respectively, and zone axis [ 1 12 ]. So LMO-A exhibits a well-ordered spinel-type structure. It is well known that the electrochemical properties of LiMn2O4 cathodes strongly depend on the morphology, particle size distribution and the porosity of the architecture [9, 10]. The homogeneous micro/nano-scaled particles of LMO-A provide a short lithium ion and electron diffusion lengths, while the nanoporosities between the primary particles forming the three-dimensional (3D) framework allow for the penetration of the liquid electrolyte. Hence, together with the above results, the enhanced electrochemical performance of LMO-A cathode used in LIBs is expected. To evaluate the electrochemical performance of the as-obtained samples as cathodes in LIBs, galvanostatic charge-discharge tests were carried out in the voltage range of 3.1– 4.3 V at room temperature in this work. Fig. 6 (a) shows the first galvanostatic chargedischarge curves of the LMO-A and LMO-B cathodes at 0.1 C (1 C = 148 mA g -1). The charge and discharge curves of the two electrodes show two plateaus at about 3.9 and 4.1 V vs. Li/Li+, the well-known characteristic of intercalation reactions of LiMn 2O4 [19, 33]. The lithium extraction-insertion mechanism in LiMn2O4 electrode could proceed through a solid-state solution insertion reaction and thus cause its two slope-like charge/discharge profile [34]. The insertion/extraction behaviour of Li-ions can be expressed as: L iM n 2 O 4



L i 1 -x M n 2 O 4  x L i  x e



(4)

During the charge process, the first plateau means an extraction of half of the Li-ions from the tetrahedral sites in LiMn2O4 with Li-Li interaction [10, 20]; the second plateau

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reflects the further extraction of the remaining Li-ions from the tetrahedral sites without the aid of any Li-Li interaction [10, 33]. The initial charge/discharge capacity of LMO-A cathode is 132.2 and 131.0 mAh g-1 at 0.1 C, respectively, higher than those of the LMOB cathode (119.9 and 117.8 mAh g-1). This may be attributed to the surface nanostructure, which can improve lithium-ion utilization. The differential capacity vs. voltage (dQ/dV) curves of the LMO-A and LMO-B electrodes [Fig. 6 (b)] show two pairs of separated redox peaks at around 4.05/3.98 V and 4.15/4.12 V that correspond to the two-step transformation of Li0.5Mn2O4/LiMn2O4 and λ-MnO2/Li0.5Mn2O4 [33]. The LMO-A cathode exhibits a sharper characteristic peaks than the LMO-B cathode revealing the superior kinetic for Li-ions transports in LMO-A. The rate capability of the LMO-A and LMO-B electrodes was investigated on the Li half cells at different current rates in the range of 0.2–20 C. As shown in Fig. 7(a), the discharge capacities of LMO-A cathode are much higher than those of LMO-B cathode. The specific capacities of LMO-A electrode at the rates of 0.2, 5, 10 and 20 C are 120.7, 101.2, 96.9 and 92.6 mAh g-1, respectively, which are much higher than those of LMO-B electrode (107.3, 58.8, 44.0 and 24.1 mAh g-1). This indicates that the discharge capacity at ultra-high rate of 20 C is about 76.7 % of the capacity at 0.2 C, which is superior to the reported values for LiMn2O4 cathode material in the literature [8, 9, 19, 35-37]. The specific energy and the specific power of the electrodes were calculated using a previously reported method [2, 3, 38]. From the ragone plots in Fig. 7(b), it is clearly found that the high power capability of LMO-A electrode is much better than that of LMO-B electrode. At 20 C, the specific energy of the LMO-A cathode maintains about 72.1 % of its energy obtained at 0.2 C, which is much higher than for the LMO-B cathode

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(~ 20.0 %). These findings demonstrate that the vacuum-induced self-assembly process is a quite powerful method to prepare nanoporous LiMn 2O4 cathode materials with excellent rate performance. The discharge profiles of the two electrodes at various rates are shown in Fig. 7(c) and (d). The discharge curves at low rates exhibit two distinct plateaus, indicating two-step Li-ion extraction behaviour. With increasing discharge current, the separation of the two plateaus gradually becomes indistinct, and the discharge plateau potentials shift toward lower voltages due to the increased ohmic voltage drop and cell polarization at high current density [20, 39]. It can also be seen from Fig. 7(c) and (d) that the polarization of LMO-A electrode is smaller than that of LMO-B electrode, reflecting the higher Li-ions diffusion in a small particle or a faster charger reaction rate on the particle surface. This result is attributed to the unique nanoporous structure, which could promote the entry of electrolyte into the interior of the LMO-A electrode, and thus forms a continuous 3D network for Li-ion and electron pathways. The interpenetrating 3D network can provide short Li-ion and electron diffusion distances and some electrolyte-filled pores between particles. To investigate the cycleability of LMO-A electrode, the fabricated coin cells were cycled at 1 and 10 C at room temperature. The potential versus charge-discharge capacity traces are presented in Fig. 8(a). It can be seen that both coin cells exhibit similar curve shape, showing the signature of intercalation reactions of LiMn2O4 cathodes. Furthermore, it is obvious that the cell cycled at 10 C shows a larger IR drop between charge and discharge over the total voltage range than the cell cycled at 1 C because of a polarization effect caused by the limited Li-ion diffusion. Nevertheless, the LMO-A electrode exhibits

14

excellent cycling performance, as shown in Fig. 8(b). At 1 C, the discharge capacity of the LMO-A is 97.5 mAh g-1 after 100 cycles, maintaining 90.9 % of the initial capacity. At the higher rate of 10 C, after 100 cycles, almost 95.9 % of the first capacity (94.5 mAh g-1) and near 100 % of its coulombic efficiency are achieved for LMO-A cathode. This confirms that nanoporous LMO-A electrode can deliver better cycling stability than nanoparticles [18], micro-sized spheres [40], one-dimensional nanorods [12] and nanotubes [41]. The good cycling performance of LMO-A electrode is clearly attributable to the nanoporous structure, which can not only accommodate the strains and volume expansion/contraction during the repeated Li-ion insertion/extraction processes, but also provide a relatively low specific surface area. The superior cycleability and excellent ultra-high rate capability of the nanoporous LMO-A are crucial for the application of LiMn2O4 cathodes for high-power LIBs. In order to further investigate the electrode kinetics, the apparent activation energies of the two electrodes were calculated from electrochemical impedance spectroscopy (EIS) experiments. Fig. 9(a) and (b) provide the Nyquist impedance plots and the equivalent circuits of two cathodes. The impedance data were collected at different temperatures after completing 0.2-20 C galvanostatic charged/discharged tests. Z' and -Z'' are respectively the magnitude of the real and the imaginary impedance. The impedance spectra for two cathodes are composed of a semicircle in the high frequency region and an oblique line at approximately 45o in the low frequency region. In general, the intercept in Z′-axis at the Nyquist curve reflects the ohmic resistance (Rs); the semicircle at the high frequency represents the charge transfer resistance and capacitance (Rct||CPE1); and the oblique line reflects the solid-state diffusion of Li-ions, namely, the Warburg impedance

15

(W1). The Rs and Rct values of the two cathodes were obtained by applying ZView 2.0 (shown in Table S3). The apparent activation energy of the interface reactions was investigated to understand the kinetics of the interfacial Li-ion transfer process. The apparent activation energy (Ea) for the lithium intercalated into LiMn 2O4 can be calculated from the following equations [42]: i0  R T n F R c t

(5)

i0  A e x p   E a R T



(6)

Here i0 is the exchange current, R is the gas constant, T (K) is the absolute temperature, A is a temperature-independent coefficient, n is the number of transferred electrons, and F is the Faraday constant. Fig. 9(c) shows the Arrhenius curves of ln (i0/A) as a function of 1/T. The activation energies (Ea = -Rk, where k is the slope of the fitting line) of LMO-A and LMO-B electrodes are calculated to be 18.64 and 42.15 kJ mol -1, respectively. These results demonstrate that LMO-A features a lower apparent activation energy for the intercalation of Li-ions from the LiMn2O4 than LMO-B. The lower activation energy of LMO-A indicates that the surface of LMO-A is highly conductive, thereby accelerating the ion and electron diffusion through the primary particles and thus ensuring ultra-high rate capability. The EIS can also be applied to calculate the “apparent” chemical diffusion coefficient of Li-ion (DLi) using Eqs. (7) and (8) below [5, 43]: D Li  R T 2

2

2A

2

n F C  4

Z   R s  R ct   

4

2

2



(7)

1 2

(8)

16

Where R is the gas constant, T (K) is the absolute temperature, A is the geometric surface area of the cathode, n is the number of electrons transferred in the insertion/extraction reaction, F is the Faraday constant, C is the concentration of lithium ion in the solid, and σ is the Warburg factor, which is relative to the real impedance (Z′). The Warburg factor (σ) can be obtained from the slope of the lines between Z′ and ω -1/2 [see Fig 9(d) and (e)]. Actually, it is impossible to obtain the absolute values of DLi in EIS technique because it has uncertainties such as the active surface area [44-46]. However, this work is intended to be a comparative research on the kinetics of lithium ion intercalation and deintercalation reaction in two half-coin cells. Fig. S6 shows the Li-ion diffusion coefficients of two electrodes. It can be seen in Fig. S6 that the DLi of LMO-A is always higher than that of LMO-B at different temperatures, indicating the improvement in the kinetics of LMO-A. The diffusion apparent activation energy (EaD) of the two electrodes can be also calculated from EIS results using a previously reported method [42, 47]. As shown in Fig. 9(f), the EaD of LMO-A is about half of that of LMO-B. These results indicate that, compared to the LMO-B, the LMO-A can offer smaller energy barrier for Li-ion motion at the surface of the electrode to ultra-high rates.

4. Conclusion In conclusion, nanoporous LiMn2O4 cathode material has been successfully prepared via a simple, low-cost and scalable vacuum-induced self-assembly method. The nanoporous morphology appears to be a significant step towards enhancing the kinetic properties of the electrode material. The nanoporous spinel LiMn 2O4 presents superior

17

rate capability and cycleability when used as cathode material for LIBs. Specifically, the LMO-A electrode delivers a discharge capacity of 92.6 mAh g -1 at the rate of 20 C. After over 100 cycles at 10 C, almost 95.9 % of the initial capacity (94.5 mAh g-1) is retained. These results are superior to those of most LiMn2O4 cathode materials prepared by other reported synthetic approaches. The EIS results indicate that, compared to LMO-B, LMOA has a higher Li-ion diffusion coefficient and lower apparent activation energy because of its unique structural properties. The enhanced electrochemical performance coupled with the facile fabrication method could make the nanoporous LiMn 2O4 a promising cathode material for high power LIBs. Furthermore, we believe that this process is accessible to scale up for synthesizing other nanostructured electrode materials.

AUTHOR INFORMATION Corresponding Author *Email: [email protected] (Dr. Xiao-Dong Guo); *Email: [email protected] (Dr. Shu-Lei Chou). Author contributions The manuscript was written through contributions of all authors. All authors have given the approval to the final version of the manuscript. Notes The authors declare no competing financial interest.

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ACKNOWLEDGMENT This work was supported by AutoCRC Project 1-111, the National Natural Science Foundation of China (Grant No. 21506133), the Science and Technology Pillar Program of Sichuan Province (2014GZ0077), the Youth Foundation of Sichuan University (No. 2011SCU11081), the Doctoral Program of Higher Education of China (No. 20120181120103), and the Open Found of National Engineering centre for Phosphorus Chemical Industry (2013LF1012). The Analysis and Test Centre of Sichuan University also supported this study. Authors also want to thank for Dr. Silver Tania for her critical reading of the manuscript.

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FIGURES Fig. 1 Schematic illustration of the controlled preparation of nanoporous LiMn2O4. Fig. 2 Rietveld refinement of the XRD patterns for (a) LMO-A and (b) LMO-B. Fig. 3 SEM images at different magnifications of (a, b) LMO-A and (c, d) LMO-B. Fig. 4 (a) Nitrogen absorption and desorption isotherms, and (b) the pore size distributions for the materials. Fig. 5. TEM images of LMO-A at (a) low and (b) high magnification; (c) HRTEM image and (d) Fourier transform (FFT) pattern of the LMO-A. Fig. 6 (a) Initial charge–discharge curves and (b) corresponding dQ/dV profiles of the LMO-A and LMO-B electrodes at 0.1 C between 3.1 and 4.3 V at room temperature. Fig. 7 (a) Cycle life of the two electrodes at different current densities from 0.2 to 20 C at room temperature; (b) Ragone plots for the two electrodes; discharge voltage profiles under different current rates (0.2–20 C) for (c) LMO-A and (d) LMO-B electrodes. Fig. 8 (a) Potential vs. specific capacity profiles and (b) corresponding cycling performance of LMO-A electrode at 1 and 10 C, respectively. Fig. 9 Nyquist plots of (a) LMO-A and (b) LMO-B electrodes at different temperatures; the insets in (a, b) are the equivalent circuit; (c) Arrhenius profiles of ln (i0/A) vs. 1/T for the electrodes; relationship between the real part of the complex impedance and ω-1/2 in the low frequency region for (d) LMO-A and (e) LMO-B electrodes at a series of temperatures; (f) log10 DLi vs. 1/T curves during lithium insertion.

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Fig. 1 Schematic illustration for the controlled preparation of nanoporous LiMn2O4.

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Fig. 2 Rietveld refinement of the XRD patterns for (a) LMO-A and (b) LMO-B.

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Fig. 3 SEM images at different magnifications of (a, b) LMO-A and (c, d) LMO-B.

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Fig. 4 (a) Nitrogen absorption and desorption isotherms, and (b) the pore size distributions for the materials.

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Fig. 5. TEM images of LMO-A at (a) low and (b) high magnification; (c) HRTEM image and (d) Fourier transform (FFT) pattern of the LMO-A.

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Fig. 6 (a) Initial charge–discharge curves and (b) corresponding dQ/dV profiles of the LMO-A and LMO-B electrodes at 0.1 C between 3.1 and 4.3 V at room temperature.

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Fig. 7 (a) Cycle life of the two electrodes at different current densities from 0.2 to 20 C at room temperature; (b) Ragone plots for the two electrodes; discharge voltage profiles under different current rates (0.2–20 C) for (c) LMO-A and (d) LMO-B electrodes.

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Fig. 8 (a) Potential vs. specific capacity profiles and (b) corresponding cycling performance of LMO-A electrode at 1 and 10 C, respectively.

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Fig. 9 Nyquist plots of (a) LMO-A and (b) LMO-B electrodes at different temperatures; the insets in (a, b) are the equivalent circuit; (c) Arrhenius profiles of ln (i0/A) vs. 1/T for the electrodes; relationship between the real part of the complex impedance and ω-1/2 in the low frequency region for (d) LMO-A and (e) LMO-B electrodes at a series of temperatures; (f) log10 DLi vs. 1/T curves during lithium insertion.

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