# NEXT-GENERATION ELECTROCHEMICAL ENERGY STORAGE DEVICES

EDITED BY : Yuxin Tang, Oleksandr Malyi, Wenping Sun and Jianqing Zhao PUBLISHED IN : Frontiers in Chemistry

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ISSN 1664-8714 ISBN 978-2-88963-569-6 DOI 10.3389/978-2-88963-569-6

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## NEXT-GENERATION ELECTROCHEMICAL ENERGY STORAGE DEVICES

Topic Editors: Yuxin Tang, University of Macau, China Oleksandr Malyi, University of Oslo, Norway Wenping Sun, Zhejiang University, China Jianqing Zhao, Soochow University, China

Citation: Tang, Y., Malyi, O., Sun, W., Zhao, J., eds. (2020). Next-generation Electrochemical Energy Storage Devices. Lausanne: Frontiers Media SA. doi: 10.3389/978-2-88963-569-6

# Table of Contents

Xiaoxia Liu, Yufei Ren, Lan Zhang and Suojiang Zhang

Xiaosheng Tang, Jie Yang, Shiqi Li, Weiwei Chen, Zhiping Hu and Jing Qiu

*69* In-situ *Functionalization of Metal Electrodes for Advanced Asymmetric Supercapacitors*

Leimeng Sun, Xinghui Wang, Yurong Wang, Dongyang Xiao, Weifan Cai, Yuan Jing, Yanrong Wang, Fangjing Hu and Qing Zhang

Liying Deng, Wangyang Li, Hongnan Li, Weifan Cai, Jingyuan Wang, Hong Zhang, Hongjie Jia, Xinghui Wang and Shuying Cheng

*103 First-Principle Insights Into Molecular Design for High-Voltage Organic Electrode Materials for Mg Based Batteries*

Johann Lüder and Sergei Manzhos

*121 Building High Power Density of Sodium-Ion Batteries: Importance of Multidimensional Diffusion Pathways in Cathode Materials* Mingzhe Chen, Yanyan Zhang, Guichuan Xing and Yuxin Tang

# Enhanced Cycling Stability of Cation Disordered Rock-Salt Li1.2Ti0.4Mn0.4O2 Material by Surface Modification With Al2O3

Baojun Huang, Rui Wang\*, Yansheng Gong, Beibei He\* and Huanwen Wang

Faculty of Materials Science and Chemistry, China University of Geosciences, Wuhan, China

Cation disordered rock-salt lithium-excess oxides are promising candidate cathode materials for next-generation electric vehicles due to their extra high capacities. However, one major issue for these materials is the distinct decline of discharge capacities during charge/discharge cycles. In this study, Al2O3 layers were coated on cation disordered Li1.2Ti0.4Mn0.4O2 (LTMO) using atomic layer deposition (ALD) method to optimize its electrochemical performance. The discharge capacity after 15 cycles increased from 228.1 to 266.7 mAh g−1 for LTMO after coated with Al2O3 for 24 ALD cycles, and the corresponding capacity retention enhanced from 79.7 to 90.9%. The improved cycling stability of the coated sample was ascribed to the alleviation of oxygen release and the inhibition on the undesirable side reactions. Our work has provided a new possible solution to address some of the capacity fading issues related to the cation disordered rock-salt cathode materials.

Keywords: lithium-ion batteries, cathode, cation disorder, rock-salt, Li-excess

### INTRODUCTION

Cathode materials with high energy densities are crucial for next generation of lithium-ion batteries, especially when used in hybrid and electric vehicles (Zu and Li, 2011; Goodenough and Park, 2013; Rui et al., 2015; Wang et al., 2017; Lv et al., 2018; Tan et al., 2018; Zhang et al., 2019). For this reason, Layer-structured Li-excess materials, which could deliver capacities as high as 300 mAh g−1 , have been researched for more than 10 years (Lu et al., 2002; Wang et al., 2013; Yu et al., 2014; Hu et al., 2018). But these materials undergo irreversible O loss in the first cycle, which may cause structure densification and voltage degradation in subsequent cycles (Xu et al., 2011). Recently, cation disordered rock-salt Li-excess materials, sharing similar chemical compositions with layered Li-excess materials, have also attracted lots of attentions because of their high capacities (Yabuuchi et al., 2011, 2015, 2016a,b; Lee et al., 2014, 2015, 2017, 2018; Twu et al., 2015; Freire et al., 2016, 2017; Cambaz et al., 2018; Kitchaev et al., 2018; Zhao et al., 2019). Besides, it is reported that the O redox may undergo reversible reactions in the first cycle in this material (Wang et al., 2015), which provides the possibility to achieve reversible changes during the lengthy electrochemical cycles. However, many of these materials have problems in their cycling performances. Wang et al. prepared a cation disordered Li1.23Ni0.155Ru0.615O2 material, and the capacity drops from 295.3 to 250 mAh g−1 only after 5 charge-discharge cycles (Wang et al., 2017). Okada's group prepared another cation disordered Li1.2Mn0.4Ti0.4O2 material, but its capacity drops from 226 to about 200 mAh g−1 just after 6 cycles (Kitajou et al., 2016).

#### Edited by:

Yuxin Tang, Nanyang Technological University, Singapore

#### Reviewed by:

Xianhong Rui, Guangdong University of Technology, China Jianqing Zhao, Soochow University, China Renheng Wang, Shenzhen University, China

#### \*Correspondence:

Rui Wang<EMAIL_ADDRESS>Beibei He<EMAIL_ADDRESS>
#### Specialty section:

This article was submitted to Electrochemistry, a section of the journal Frontiers in Chemistry

Received: 25 December 2018 Accepted: 11 February 2019 Published: 04 March 2019

#### Citation:

Huang B, Wang R, Gong Y, He B and Wang H (2019) Enhanced Cycling Stability of Cation Disordered Rock-Salt Li1.2Ti0.4Mn0.4O2 Material by Surface Modification With Al2O3. Front. Chem. 7:107. doi: 10.3389/fchem.2019.00107

Similar problems happened to layered Li-excess materials, and some literatures have reported that atomic layer deposition (ALD) method may be an effective method to alleviate the problem. ALD is a powerful technique to precisely render a uniform and conformal layer at Å level on arbitrary substrate surfaces due to its pulsing and controllable reaction. Belharouak's team coated Li1.2Ni0.13Mn0.54Co0.13O2 porous powder with ultrathin Al2O3 film using ALD technique, and the coated material shows higher first cycle coulombic efficiency and improved cycling performance (Zhang et al., 2013). Xiao et al. reported that the AlPO4 coating layer by ALD can effectively protect the Li1.2Mn0.54Co0.13Ni0.13O2 against the attack from the electrolyte, and can significantly improve its initial coulombic efficiency and thermal stability (Xiao et al., 2017a). Meanwhile, Oxide-based coatings at the surface of different cathode materials via ALD method have been demonstrated to be very conformal and uniform as reported in the literatures (Zhao and Wang, 2012; Zhao et al., 2012; Zhao and Wang, 2013a,b), thus they can effectively prevent from the electrolyte attack for enhanced cycling stability of the coated cathode.

In this study, we coated a cation disordered rock-salt Li1.2Ti0.4Mn0.4O2 material with different thicknesses of Al2O3 layers by ALD technique and studied their effects on the cycling stabilities. Li1.2Ti0.4Mn0.4O2 is a typical cation disordered rocksalt material but with poor cycling performances (Kitajou et al., 2016; Yabuuchi et al., 2016a). In our previous study, we found that the valence of Mn in disordered materials may be lower than 3+ after the first cycle (Wang et al., 2015), and this low valence may cause Mn dissolution into the electrolyte (Nicolau et al., 2018). In this case, we infer that Al2O3 coating may be an effective method to increase the reversibility and cycling performances of this cation disordered material.

#### EXPERIMENTAL

Li1.2Ti0.4Mn0.4O2 was synthesized by the traditional solid-state reaction using the precursors of Li2CO3 (99%, Alfa Aesar), Mn2O3 (98%, Alfa Aesar), and TiO2 (99%, Sigma-Aldrich). Stoichiometric amounts of precursors were ball-milled for 4 h and then pressed into a pellet. The pellet was calcinated at 900◦C for 12 h in Ar, then pulverized and ball-milled at 300 rpm for 4 h. The obtained material was denoted as LTMO.

Coating was achieved by the Atomic Layer deposition (ALD) method using an Ensure Nanotech ALD system (LabNano-9100). During the ALD process, as-prepared LTMO sample was first placed in a home-made sample holder and heated to 200◦C in the reaction chamber under ca. 1.0 mbar. Nitrogen with a flow rate of 20 sccm was used as the carrying and purge gas. The Al(CH3)3 and water were employed as aluminum and oxygen precursors, respectively. In each cycle, pulse time of Al(CH3)3 was controlled at 0.02 s, and exposure time water precursor was 0.02 s. 8 s nitrogen purge steps were used after Al(CH3)3 and water exposures. The LTMO sample was coated with Al2O3 layers for 16, 24, 40 ALD cycles. The coated samples were denoted as LTMO/nAl2O3, where n stands for the number of ALD cycles.

Morphologies and crystal structures of LTMO with or without Al2O3 layers were analyzed by scanning electron microscope (SEM, SU 8010) and powder X-ray diffractometer (XRD, Bruker D8 Advance), respectively. High Resolution Transmission Electron Microscope (HRTEM, TF 20) was used to record detailed crystalline structures of the samples. Xray photoelectron spectroscopy (XPS, ESCALab 250Xi) with monochromatic Al K-α radiation was carried out to investigate the valences of the species in the samples.

To fabricate the cathode electrode, different Al2O3-coated LTMO samples were firstly ball-milled with Ketjen black at 150 rpm for 4 h, then manually mixed with polytetrafluoroethylene (PTFE) binder. The mixture was rolled into a thin film, in which the weight ratios of LTMO/nAl2O3, Ketjen black, and PTFE are 70: 20: 10. The surface mass density of each electrode film is about 4.4 mg cm−2 . Cells were assembled according to our previous study (Wang et al., 2015). Typically, a two electrode swagelok cell was fabricated using the Al2O3-coated LTMO thin film and lithium foil as the working electrode and the counter electrode, respectively. Borosilicate glass fiber membrane (Whatman) was used as the separator, and 1 M solution of LiPF6 dissolved in ethylene carbonate/dimethyl carbonate (1:1 by volume) was used as the electrolyte. All cells were cycled on an LANHE CT2001A system (Wuhan LAND Electronics Co.) between 4.8 and 1.5 V with a current density of 10 mA g−1 at room temperature. The electrochemical impedance spectroscopy (EIS) data were recorded using a Gamry Reference 3,000 equipment. Cyclic voltammetry (CV) was measured on a CHI760E electrochemical workstation with the scanning rate of 0.5 mV s−1 and the potential range of 4.8–1.5 V.

#### RESULTS AND DISCUSSION

X-ray diffraction (XRD) patterns of the samples are shown in **Figure 1a**, and all the results correspond well to the disordered rock-salt structure. According to previous studies, cations are randomly placed on the 4b position (1/2, 1/2, 1/2), and the O atom is placed on the 4a position (0, 0, 0) (Wang et al., 2015). In cation disordered rock-salt materials, Li excess could build the 0-TM channels, through which Li ions could move out and back into the cathode material (Lee et al., 2015). According to the XRD patterns, no signal of Al2O3 is detected, and the reason may be that the amount of coated Al2O3 is too small or the coated Al2O3 is amorphous. The morphologies of LTMO before and after Al2O3 coating are shown in **Figures 1b,c**. No obvious differences are observed, indicating the coated Al2O3 particle may be too small to be surveyed in SEM. Energy Dispersive Spectrometers of the two samples are shown in **Figure 2a**. The characteristic peak of Al element could clearly be observed in the LTMO/24Al2O3 sample, while didn't show up in the LTMO sample. The LTMO/24Al2O3 sample was further examined by the elemental mapping experiments. and the results are shown in **Figures 2b-f**. Results show that Ti, Mn, O and Al elements distribute uniformly in the sample, which infer the Al2O3 layer exists uniformly in the sample.

In order to directly observe the coating Al2O3 layers on the surface of LTMO particles, we investigate bare LTMO and LTMO/24Al2O3 samples using high resolution transmission

electron microscopy (HR-TEM). **Figure 3a** is the HR-TEM image of bare LTMO particles, showing well-defined lattice fringes in the surface region as well as those in the bulk region. In contrast, an obvious coating film in the surface region is clearly observed for LTMO/24Al2O3 sample (**Figure 3b**). It is different from the inner areas, and we infer it may be the Al2O3 coating layer. The thickness of the coating layer is around 3–5 nm.

In order to confirm Al2O3 layer exists on the sample surfaces, XPS was used to probe the surface compositions of the LTMO and LTMO/24Al2O3. All spectra were calibrated with the C 1s peak at 284.6 eV. **Figure 4A** presents the elemental XPS spectra of LTMO and LTMO/24Al2O3 samples, showing both samples contain the characteristic peaks of Ti, Mn, O, and C. From the partial enlarged figure (inset of the **Figure 4A**), the elemental peak of Al is only observed in the LTMO/24Al2O3 spectrum. The core level binding energies of Ti are aligned at 457.9 eV (2p3/2) and 463.6 eV (2p1/2) (**Figure 4B**), which is in good agreement with Ti4+ ions in other titanium-based compounds (Kim et al., 2017). The Mn 2p XPS spectra of LTMO and LTMO/24Al2O3 reveal peaks at 641.6 eV and 653.2 eV (**Figure 4C**), which are characteristic peaks of Mn3+ 2p3/2 and 2p1/2 (Das et al., 2011, 2013). This result confirms the valences of Mn in both samples are 3+. The Al 2p3/2 core level XPS spectrum displays a binding energy of 74.7 eV for the LTMO/24Al2O3 (**Figure 4D**), affirming that the chemical composition of the surface coating is Al2O3 for the LTMO/24Al2O3 sample.

The charge and discharge profiles of uncoated and coated LTMO samples for the first cycle are shown in **Figure 5A**, showing all the samples have the identical profiles except for the voltage platform and capacity. It is obviously seen that all the coated LTMO samples have higher first discharge plateau and coulombic efficiency, and we think this may be the effect of Al2O3 layer. Though the O reaction in LTMO may happen in the form of O redox, there might still be little amount of oxygen release in the first cycle. We infer the Al2O3 layer may alleviate this oxygen release and improve the O reaction reversibility, thus improve the discharge plateau and coulombic efficiency. The charge and discharge profiles for the 10th cycle are shown in **Figure 5B**, and the LTMO/24Al2O3 sample presents the largest discharge capacity. **Figure 5C** shows the cycling stability of the pristine and ALD coated samples. Bare LTMO shows an initial discharge capacity of 286.3 mAh g−1 (818.6 Wh kg−1 ) and the capacity drops rapidly to 228.1 mAh g−1 (587.7 Wh kg−1 ) after 15 cycles, which is only 79.7% for its initial discharge capacity. After coated with Al2O3 for 16 ALD cycles, an initial discharge of 289.1 mAh g−1 (856.6 Wh kg−1 ) is observed and the capacity retention reaches 85.2%. It is worth noting that the LTMO/24Al2O3 sample could deliver a capacity of 293.4 mAh g−1 and energy density of 884.5 Wh kg−1 in the first cycle. After 15 cycles, the capacity and energy density could still maintain 266.7 mAh g−1 and 709.5 Wh kg−1 , respectively. And the capacity retention after 15 cycles is 90.9%, which is 11.2% higher than that of the pristine LTMO. Besides, the coulombic efficiencies also increased from 91.5 to 93.0% after the 24 cycles ALD coating. However, further increasing the ALD cycle numbers of coated Al2O3, the inferior performance for LTMO/40Al2O3 is obtained, and the capacity retention decreases to 84.6%. The reason should

be the thick insulated Al2O3 layer may inhibit the electric and ionic transportation (Jung et al., 2010). Based on above results, it can be concluded that proper thickness of Al2O3 layer can greatly improve the cycling stability of LTMO electrodes, while the over-thick Al2O3 layer leads to the reverse effect on the electrode.

**Figure 6** show the cyclic voltammetric (CV) curves of LTMO and LTMO/24Al2O3 samples in the first three cycles at a

scanning rate of 0.5 mV s−1 . For the LTMO electrode (shown in the **Figure 6A**), the anodic peaks located at about 3.95 V corresponds to the oxidation of Mn3+ to Mn4+ (Xiao et al., 2017b) during the initial charge process. Another anodic peaks at around 4.63 V may be attributed to the oxygen loss from the crystal structure and formation of O2− 2 or O2 from O2−, which is usually observed on the Mn-based Li-excess cathode materials (Ma et al., 2016). While this peak disappears in the subsequent cycles, showing the irreversibility of the oxidation of O2− in the LTMO electrode. In the following cathodic process, the peak at around 2.89 V corresponds to the reduction of Mn4+ to Mn3+ (Kong et al., 2017). It can be clearly seen that this peak moves to lower voltage range in the following cycles, indicating the decreasing of the valence of Mn. As to LTMO/24Al2O3 (shown in **Figure 6B**), two differences can be observed. First, the anodic peak at around 4.63 V still can be observed in the second and third cycles. Second, the cathodic peak located at around 3.17 V doesn't move during the cycles. These results infer the oxygen loss from the crystal structure in LTMO/24Al2O3 is milder than in LTMO. As a result, the Mn redox reaction in LTMO/24Al2O3 happens between Mn4+ and Mn3+. While Mn in pristine LTMO reduces to lower than 3+ because of the O loss and related densification in the first cycle.

In order to investigate the intrinsic factor of the improvement in the electrochemical performance of the LTMO/24Al2O3

sample, electrochemical impedance spectra (EIS) were collected on the bare LTMO and the LTMO/24Al2O3 after charging to 4.8 V and resting for 4 h at the different cycles, shown in **Figures 7A,B**, respectively. All the Nyquist plots comprised a depressed semicircle from high to middle frequencies and an inclined line at low frequency. The simulated equivalent circuit is presented as an inset. The R1 represents the Ohmic resistance coming from the separator, electrolyte and other components. The semicircle shows the charge transfer reaction composed of a charge transfer resistor (R2) and a constant phase element (CPE1), the inclined line stands for the Warburg diffusion impedance (ZW). The LTMO/24Al2O3 electrode shows the smaller R2 of 101.6 before cycling and remains at 228.4 , 433.8 and 851.3 after 1st, 5, and 10th cycling, respectively. Nevertheless, the much larger R2 values of the LTMO electrode are seen (i.e., 255.8 before cycling and changed to 557.7 after 1st cycling, even to 10,207 for the 5th cycling, respectively), which infer complex side reactions may happen on the electrode surface and these reaction products may have blocked the ionic transfer process and result in worse cycling performance. Thus, it can be concluded that an appropriate thickness of Al2O3 layer guarantees the stable charge transfer and structural integrity of the cathode electrodes, which led to good cycling performance.

### CONCLUSION

ALD technique was successfully used to deposit ultrathin Al2O3 coating layer onto the surface of the LTMO particles. Comparing with the uncoated LTMO, ALD process can reduce the polarization, restrain the undesirable side reactions, and suppress the increasing charge transfer resistance during cycling, which results in the significantly improved electrochemical performance of Al2O3-coated LTMO (LTMO/24Al2O3). The controllable ALD technology provides a promising guideline for the surface modification of disordered rock-salt cathode materials with high electrochemical performance.

### AUTHOR CONTRIBUTIONS

RW conceived and designed the research. BaH carried out the experiments. YG and HW contributed to the discussion. RW wrote the manuscript with the help of BeH. All authors reviewed the manuscript.

#### ACKNOWLEDGMENTS

This work was supported by the National Natural Science Foundation of China (51702294 and 51702295).

#### REFERENCES

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Huang, Wang, Gong, He and Wang. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

## The Electrochemical Performances of n-Type Extended Lattice Spaced Si Negative Electrodes for Lithium-Ion Batteries

Moonsang Lee1†, Dockyoung Yoon2,3†, Uk Jae Lee4 , Nurzhan Umirov 3 , Aliya Mukanova5 , Zhumabay Bakenov 5 and Sung-Soo Kim3 \*

*1 Korea Basic Science Institute, Daejeon, South Korea, 2 SK Innovation, Daejeon, South Korea, 3 Graduate School of Energy Science and Technology, Chungnam National University, Daejeon, South Korea, 4 School of Integrative Engineering, Chung-Ang University, Seoul, South Korea, 5 National Laboratory Astana, School of Engineering, Nazarbayev University, Institute of Batteries, Astana, Kazakhstan*

#### Edited by:

*Jianqing Zhao, Soochow University, China*

#### Reviewed by:

*Yong Min Lee, Daegu Gyeongbuk Institute of Science and Technology (DGIST), South Korea Zhiqiang Xie, University of Tennessee Space Institute (UTSI), United States*

> \*Correspondence: *Sung-Soo Kim<EMAIL_ADDRESS>
*†These authors have contributed equally to this work*

#### Specialty section:

*This article was submitted to Electrochemistry, a section of the journal Frontiers in Chemistry*

Received: *18 February 2019* Accepted: *14 May 2019* Published: *31 May 2019*

#### Citation:

*Lee M, Yoon D, Lee UJ, Umirov N, Mukanova A, Bakenov Z and Kim S-S (2019) The Electrochemical Performances of n-Type Extended Lattice Spaced Si Negative Electrodes for Lithium-Ion Batteries. Front. Chem. 7:389. doi: 10.3389/fchem.2019.00389* The electrochemical performances of lithium-ion batteries with different lattice-spacing Si negative electrodes were investigated. To achieve a homogeneous distribution of impurities in the Si anodes, single crystalline Si wafers with As-dopant were ball-milled to form irregular and agglomerated micro-flakes with an average size of ∼10µm. The structural analysis proved that the As-doped Si negative materials retain the increased lattice constant, thus, keep the existence of the residual tensile stress of around 1.7 GPa compared with undoped Si anode. Electrochemical characterization showed that the As-doped Si anodes have lower discharge capacity, but Coulombic efficiency and capacity retention were improved in contrast with those of the undoped one. This improvement of electrochemical characteristics was attributed to the increased potential barrier on the side of Si anodes, inherited from the electronic and mechanical nature of Si materials doped with As. We believe that this study will guide us the way to optimize the electrochemical performances of LIBs with Si-based anodes.

Keywords: lithium-ion batteries, silicon, dopant, arsenic, discharge capacity, retention

## INTRODUCTION

Lithium-ion batteries (LIBs) have been considered as an important energy storage device for modern portable electronics and electric vehicles (Goodenough and Park, 2013; Nitta et al., 2015; Kim et al., 2016). Until now, the graphite is typically used as anode material in commercial LIBs. However, its theoretical capacity (372 mAh g−1 ) is not sufficient for the next-generation mobile applications requiring high energy density (Peng et al., 2010; Ma et al., 2014; Ko et al., 2015). To overcome the energy density problem of current LIBs, the use of Si as the anode in LIBs can be an alternative way due to its unique physical properties such as high theoretical capacity of 3,580 mAh g −1 (Li15Si4 state) and low discharge potential (<0.5 V vs. Li/Li+) (Wang and Dahn, 2006; Liu et al., 2012; Shi et al., 2016). Even if the Si material has such good advantages for LIBs, the commercial success of LIBs with Si anodes has been impeded by fast capacity fading of Si anodes. This fading results from large volume change more than 300% of Si anodes during lithiation and delithiation, leading to their disintegration (Ryu et al., 2004; Kim et al., 2011). Furthermore, semiconductor characteristics of Si make it even more difficult to be utilized in the practical application of LIBs as

**13**

a negative electrode owing to the electrical losses due to large ohmic resistance (Wen and Tian, 2013; Wang et al., 2015).

Recently, various research groups have reported that the electrochemical performances of LIBs with Si anodes can be improved by using Si/Metal alloy electrode (Umirov et al., 2019a,b), the formation of nanostructured-anodes (Mukanova et al., 2018a,b), hollow and yolk-shell structured-composites (Yao et al., 2011; Wang et al., 2013), and doping of Si with impurities such as copper (Cu) (Wen and Tian, 2013), arsenic (As) (McSweeney et al., 2014), aluminum (Al) (Legrain and Manzhos, 2015), silver (Ag) (Talla et al., 2015), boron (B) (Long et al., 2011; Rousselot et al., 2012; Yi et al., 2013), nitrogen (N) (Han et al., 2016), and phosphorus (P) (Domi et al., 2016). Among other approaches, incorporating dopants into Si architectures has gained considerable attention as a suitable route for improving the rate capability, electrical conductivity, and capacity retention of the Si anodes (Mukanova et al., 2018b).

To date, many methods have been proposed for the Si powder preparation, including solution and solid-state based techniques, etc. Taking into account that the electronic/silicon-based semiconductor technologies have widely entered into society and industry, it is easy to imagine how much semiconductor class Si containing waste is produced every day. These silicon materials can be recycled for use in LIBs.

In this paper, we used the Si wafer in order to demonstrate the possibility of the Si electronic wastes to be reused in LIBs. As-modulated Si micro-flake anodes were investigated as an anode. To achieve a homogeneous distribution of As-impurity in the Si anodes, As-doped Si wafer was prepared by ballmilling and employed for the anode in a lithium cell. Moreover, this work covers the comparative analysis in the structural and electrochemical responses of As-doped and undoped Si microflake anodes to investigate the effects of the lattice spacing in Si anodes.

We believe that this paper will shed lights on the trends in the electrochemical performance [cycle retention, Coulombic efficiency (CE), and rate capability] of LIBs with the Si microflake anode with different lattice spacing.

#### EXPERIMENTAL

The commercial As-doped single crystalline (100) Si wafers with a doping concentration of 1 × 1019 cm−3 were provided by SK Hynix and used as the anode material for LIBs. The doping concentrations of Si wafer were confirmed by secondary ion mass spectrometry (SIMS) (not shown in this paper). Furthermore, the intrinsic (undoped) Si wafers were used as the control samples to compare the effects of As doping. To remove native oxide on Si surface, Si wafers were immersed in diluted 10% hydrofluoric acid solution for 3 min and followed by rinsing in deionized (DI) water of ultra-purity (resistance 18.2 Mcm) for 5 min. Si wafers were mechanically grounded to achieve Si microflakes using zirconia balls for 2 h in paint shaker (KM-2000T, Nara Sci).

The undoped and As-doped Si based electrodes were fabricated via direct casting of slurry onto Cu foil using the Dr. Blade method. The slurry was prepared by thoroughly

FIGURE 1 | Plan-view SEM image of (A) undoped and (B) As-doped Si powders after ball-milling.

mixing of active material and PAI (Polyamide-Imide, Solvay Torlon 4000T) binder with the wt.% ratio of 9:1, respectively, in NMP (N-Methyl-2-pyrrolidone, Merck) solvent. After casting the electrodes were dried in convection oven at 110◦C to remove NMP. The loading level of electrodes were adjusted to 5 mg/cm2 .

Electrochemical performances were studied using coin-type half configuration cells with Li metal as the counter electrodes, which were assembled in an Ar-filled glove box with <1 ppm of both oxygen and moisture. The electrolyte solution used was 1.3 M lithium hexafluoro phosphate (LiPF6) dissolved in a solvent mixture of ethylene carbonate (EC) and diethyl carbonate (DEC) with a 3:7 volume ratio. Galvanostatic measurements were carried out to determine charge-discharge capacity and cycling retention in the potential range between 0.005 and 2.000 V vs. Li/Li+ under a current density of 1.0 A g−1 (0.28 C) using a computer-controlled battery measurement system (TOSCAT 3000 U).

The surface morphologies of Si microflake anodes were evaluated by field emission scanning electron microscope (FE-SEM). The stress evolution, phase transition, and crystallinity of Si microflake anodes were investigated by micro Raman spectroscopy and X-ray diffraction (XRD).

#### RESULTS AND DISCUSSION

**Figures 1A,B** shows the plan-view SEM images of undoped and As-doped Si negative electrodes, respectively, after ball-milling for 2 h. The irregular and agglomerated particles were made up with the average size of 10µm for both undoped and As-doped Si anodes. The particle size and distributions were measured by particle size analyzer (not shown in this paper). The similar morphology and size of micro particles of undoped and Asdoped Si anodes were observed, indicating that As-dopants do not influence the morphology and particle size distribution of negative anodes.

**Figure 2** exhibits XRD patterns and micro-Raman spectra of undoped and As-doped Si negative electrodes. The phase and crystallinity of Si anodes used in this work were investigated by XRD as shown in **Figure 2A**. It should be noted that all XRD peaks are sharp. This feature could be ascribed to the property of

single crystalline Si. It is found that all the peak positions of XRD patterns can be indexed to those of cubic Si. It is interesting to note that the peak positions of As-doped Si anodes were shifted to the lower angles, which is attributed to the increased lattice

constant by As-doping into Si structure as depicted in the inset of **Figure 2A**. Typical stress profiles of undoped and As-doped Si negative electrodes were studied by micro-Raman spectroscopy measurements at room temperature, as shown in **Figure 2B**. It is well-known that F2g-phonon modes in crystalline Si can be identified with peaks at around 520 cm−1 (Shimizu et al., 2015). We can easily notice that the peak position for the As-doped Si anode has been red-shifted, indicating the residual tensile stress of approximately ∼1.7 GPa compared to the undoped one. The residual stress can be calculated using the following equation (Wu et al., 2007):

$$\sigma(\text{MPa}) = 434 \times \Delta\omega \tag{1}$$

where σ is the biaxial stress, 1ω is the F2g phonon peak shift. These data are in good agreement with the XRD results in **Figure 2A**.

In order to evaluate the electrochemical characteristics of Si negative electrodes with and without dopants, the various electrochemical examinations were performed. **Figure 3** illustrates the first charge and discharge behavior of undoped and As-doped Si thick film anodes. It is clearly seen that there are plateaus in the curves, which is attributed to lithiation and delithiation by the reactions of Si atoms with Li ions (Chevrier et al., 2010). The initial charge potential plateau is located near 0.1 V, which is in a good agreement with the formation of LixSi via lithiation of crystallized Si (Han et al., 2016). The first charge and discharge capacity for undoped and As-doped Si negative electrodes are 3,733, 2,782, 3,578, and 2,722 mAh g−1 , respectively. These values correspond to the initial CE of 74.5 and 76.1%, respectively.

**Figure 4A** illustrates the cycle performance of the cells with the undoped and As-doped Si negative electrodes. The discharge capacities of the LIBs with both types of the anodes were evaluated in a cut-off potential range of 0.005–2.000 V as seen in **Figure 3**. The discharge capacity of LIB with undoped Si negative electrode was higher than that of As-doped Si in all cycle ranges. One can easily observe that the discharge capacity fading of undoped Si anode is comparable with that of As-doped Si. Notably, a fading rate of discharge capacity with the undoped Si negative anode changed more sharply near 10 cycles. The

initial discharge capacities of both anodes were decreased by approximately 50% after 30 cycles and their discharge capacities arrived at similar extent. **Figure 4B** shows the CE of both investigated materials. The subsequent cycle CE of cells with undoped and As-doped Si anodes start from 98.5 and 99.6% and further reaches over 95.0 and 95.8% after 30 cycles, respectively. In addition, the capacity retentions of both cells were measured as shown in **Figure 4C**. Despite the discharge capacity of undoped Si anode was higher than that with As-doped Si one, its capacity retention displays similar trends, corresponding to about 50% of retention. This phenomena can be explained by the fact that the As-doped Si anodes consumed more Li-ions for the generation of a solid-electrolyte interface (SEI) compared with the undoped Si anode and high energy barrier in energy band diagram, which will be mentioned later.

**Figure 5** depicts 2theta X-ray rocking curves of the undoped and As-doped Si negative electrodes with various cut-off voltages (OCV) in charging/discharging states after the 1st cycle. After the formation process, there is no sign of the Li15Si4 phase presence in XRD curves for undoped and As-doped Si anodes. Besides, XRD patterns of both undoped and As-doped Si anodes discharged up to a potential of 0.005 V clearly revealed the presence of the peaks at 39◦ and 41◦ , indicating the formation of Li15Si4 silicate (Li and Dahn, 2007). It is essential to note that the intensity of XRD patterns of undoped Si negative electrodes was much higher than that of As-doped Si one. This implies that the phase transition of undoped Si anodes into Li15Si4 phase is easier to occur compared to the doped electrode. In the discharging states, the peak intensity of Li15Si4 reduced steadily with increasing the OCV, and finally disappeared at OCV of 2 V for undoped and As-doped Si anodes.

In order to estimate the phase transition of undoped and As-doped Si anodes, we performed XRD analysis as a function of cycle number as illustrated in **Figure 6**. No Li silicide peaks were found in the XRD spectra of the samples before cycling (not shown in this paper). The Li15Si4 phase peaks in XRD

curves after 10 and 30 cycles were easily confirmed as seen in **Figures 6A,B**. Indeed, the XRD intensity of Li15Si4 phase in undoped Si anode is much higher than that in As-doped Si, which is consistent with the results obtained in **Figure 5**. In addition, the gradual increase of Li15Si4 phase upon 30 cycles is assigned to the activation process of Li-ion pathway in the electrode-electrolyte system, despite the fact that activation was accompanied by the fast capacity fading of both electrodes as observed in galvanostatic cycling.

Furthermore, we investigated the morphology of the electrodes at their cross-section and surface before and after 30 cycles using SEM imaging. Cross-sectional SEM images of undoped and As-doped Si (**Figures 7A,B**) showed that both electrodes have the thickness of approximately 20µm before cycling, while after 30 cycles, thickness increased drastically (>200%) that can be observed in **Figures 7C,D**. Besides, both samples exhibited similar changes in volume expansion. From the top-view images of undoped Si vs. Asdoped Si (**Figures 7E–H**), one can see that the morphology of electrodes looked the same before and after cycling. The extensive formation of cracks and partial pulverization of active electrode material after 30 cycles was observed in both cases shown in **Figures 7G,H**. Accordingly, SEM results supported the capacity fading phenomena revealed by the galvanostatic tests earlier for the undoped and As-doped Si anodes.

**Figure 8** compares the discharge capacity retentions of undoped and As-doped Si electrodes at various current densities from 0.1 to 4.0 C (1C = 1.0 A/g), respectively. Here, undoped Si exhibits considerable fading in discharge capacity at higher current densities compare with As-doped Si. More specifically, at the highest current density of 4.0 C, the undoped Sibased electrode shows 69% in capacity retention vs. 0.1 C. By contrast, the As-doped Si demonstrates better rate performance with ∼80% capacity retention. Obvious improvement in rate capability can be ascribed to the As-doping effect in Si structure and will be further discussed with the electronic energy band diagram.

To understand the charge/discharge characteristics of undoped and As-doped Si negative electrodes, we took the electronic energy band diagram of both Si anode-electrolyte interfaces into consideration as illustrated in **Figure 9**. It is well-known that when the semiconductor-electrolyte interface is formed, electrolyte and semiconductor reach electrochemical equilibrium by matching their electrochemical potentials. This occurs by charge transfer between both media (Prados et al., 2014). Considering this, each barrier heights of undoped and Asdoped Si anodes are represented in **Figures 9A,B**, respectively. It is noteworthy to mention that the energy barrier height from As-doped Si anode side, qVAs−doped, was found to be higher than qVundoped of undoped Si anode after equilibrium state. Hence, the flux of electrons across the electrolyte in As-doped Si with the increased potential barrier height can be impeded, thus hindering charge flow in the cell. In previously published works, the addition of dopants such as B and P decreased the lattice spacing of Si anodes and led to the increase of the insertion energy of Li ions, thus suppressing the discharging and charging capacity in LIBs (Long et al., 2011; Domi et al., 2016). However, the charging/discharging capacities of As-doped Si anodes have a different aspect, apart from volume change. To understand the electrochemical behaviors of cells with As-doped Si anodes, the energy band diagrams were illustrated, as shownv in **Figure 9**. It was found that the potential barrier of cells with As-dopant is much higher than that of undoped Si anode. We can infer that the higher potential barrier on the anode side by highly n-type doping results in the reduced charging/discharging capacity of As-doped Si anodes, compared to that of undoped Si ones. This is in a fair agreement with the results of **Figures 5**, **6**. Moreover, we consider that slightly higher CE of As-doped Si

#### REFERENCES

negative electrode is related to the potential barrier. A lower charge/discharge capacity leads to the suppression of the phase transition from c-Si to a-Si, which may result in the improved CE, corresponding to the XRD results.

#### CONCLUSION

In summary, the electrochemical performance of the cells with As-doped Si negative electrodes was investigated. The ballmilling of single-crystalline Si wafers was used to fabricate the anodes, which particles were irregular and agglomerated microflakes. The XRD and Raman analysis revealed that As-doped Si anodes have larger lattice spacing compared to undoped Si samples, indicating the residual tensile stress of approximately ∼1.7 GPa. Despite the increased lattice spacing in As-doped Si anode, its charge/discharge capacity was lower compared with that of undoped Si anodes. The Coulombic efficiency and cyclic retention of As-doped Si negative electrodes, however, were comparable to those of the undoped Si anode. This effect was attributed to the increased potential barrier height on the side of As-doped Si, affecting the electrochemical performance of LIBs with the As-doped anode. We believe that this study will provide an effective approach to achieve the enhanced electrochemical performances of LIBs with Si-based anode derived from electronic waste materials.

#### DATA AVAILABILITY

The datasets generated for this study are available on request to the corresponding author.

#### AUTHOR CONTRIBUTIONS

ML and DY conceived design of the study and carried out the experiment; UL, ZB, and S-SK contributed to the interpretation of the results; NU, AM, ZB, and S-SK wrote sections of the manuscript. All authors contributed to manuscript revision, read and approved the submitted version.

#### ACKNOWLEDGMENTS

The authors acknowledge the financial support of the KETEP (Grant No.20164010201070) and the Nazarbayev University by the research grant No.110119FD4504 Development of 3D all solid-state thin film Li-ion microbatteries.

Goodenough, J. B., and Park, K. -S. (2013). The Li-ion rechargeable battery: a perspective. J. Am. Chem. Soc. 135, 1167–1176. doi: 10.1021/ja3091438

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Lee, Yoon, Lee, Umirov, Mukanova, Bakenov and Kim. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

Frontiers in Chemistry | www.frontiersin.org

# High-Performance Solid Composite Polymer Electrolyte for all Solid-State Lithium Battery Through Facile Microstructure Regulation

Jingjing Yang1 \* † , Xun Wang1 , Gai Zhang1 , Aijie Ma1 , Weixing Chen1 , Le Shao2 \* † , Chao Shen3 and Keyu Xie3 \*

*1 School of Materials and Chemical Engineering, Xi'an Technological University, Xi'an, China, 2 Shaanxi Coal Chemical Industry Technology Research Institute Co. Ltd., Xi'an, China, 3 State Key Laboratory of Solidification Processing, School of Materials Science and Engineering, Center for Nano Energy Materials, Northwestern Polytechnical University and Shaanxi Joint Laboratory of Graphene (NPU), Xi'an, China*

#### Edited by:

*Jianqing Zhao, Soochow University, China*

#### Reviewed by:

*Wangwang Xu, Louisiana State University, United States Xiaowei Wang, National University of Singapore, Singapore*

#### \*Correspondence:

*Jingjing Yang<EMAIL_ADDRESS>Le Shao<EMAIL_ADDRESS>Keyu Xie<EMAIL_ADDRESS>
*†These authors have contributed equally to this work*

#### Specialty section:

*This article was submitted to Electrochemistry, a section of the journal Frontiers in Chemistry*

Received: *27 March 2019* Accepted: *14 May 2019* Published: *31 May 2019*

#### Citation:

*Yang J, Wang X, Zhang G, Ma A, Chen W, Shao L, Shen C and Xie K (2019) High-Performance Solid Composite Polymer Electrolyte for all Solid-State Lithium Battery Through Facile Microstructure Regulation. Front. Chem. 7:388. doi: 10.3389/fchem.2019.00388* Solid composite polymer electrolytes are the optimal candidate for all solid-state lithium batteries, because of their enhanced ionic conductivities, long-life cycle ability and compatibility to lithium anode. Herein, we reported a kind of solid composite polymer electrolyte comprised of poly(ethylene oxide), graphitic-like carbon nitride and lithium perchlorate, which was prepared by a facile solution blending method. Microstructure of the solid composite polymer electrolyte was regulated by thermal annealing and interaction among components and was characterized by XRD, DSC, FTIR-ATR, and ROM. The obtained solid composite polymer electrolyte achieved an ionic conductivity as high as 1.76 ×10−5 S cm−1 at 25◦C. And the electrochemical stable window and the lithium ion transference number, t+, were also obviously enhanced. LiFePO4/Li solid-state batteries with the annealed PEO-LiClO4-g-C3N4 solid polymer electrolyte presented a high initial discharge capacity of 161.2 mAh g−1 and superior cycle stability with a capacity retention ratio of 81% after 200 cycles at 1C at 80◦C. The above results indicates that the thermal annealing treatment and g-C3N4 as a novel structure modifier is crucial for obtaining the high-performance solid composite polymer electrolytes used in the all solid-state lithium battery.

Keywords: all solid-state lithium battery, solid composite polymer electrolyte, microstructure, graphite-like carbonitride, electrochemical property

## INTRODUCTION

During recent years, high-energy-density storage batteries are urgently needed to satisfy the increasing demand in electric vehicles, consumer electronics and grid energy storage (Manthiram et al., 2017; Fan et al., 2018). Lithium metal secondary batteries have been considered as the potential candidate for high-energy-density storage batteries, since lithium metal anode possessed the highest theoretical specific capacity (3,862 mAh g−1 ) and the lowest reduction potential (−3.04V VS. standard hydrogen electrode) (Shen et al., 2018; Wu et al., 2018). Moreover, lithium metal used asthe anode can even act as the lithium source in the battery with non-lithiated materials such as sulfur or oxygen as the cathode to achieve higher storage capacity compared to the current

**21**

commercial lithium-ion batteries (Pan et al., 2019). However, practical and large-scale development of lithium metal batteries with liquid electrolytes is limited by safety concerns and the rapid capacity attenuation and short cycling capability (Fan et al., 2018). During the charge-discharge process of lithium metal secondary batteries with organic liquid electrolytes, the propensity of faster lithium dendrites formation was resulted from the plating and stripping of Li ions on the anode, which would result in the capacity attenuation and poor cycling capability and short-circuit of cells.

In order to further develop the practical use of lithium metal as anode, scientists made great efforts to solve the lithium dendrites problems and achieved much progress on suppressing dendrite formation and growth (Li D. et al., 2018; Nie et al., 2018; Shen et al., 2018; Wu et al., 2018; Pan et al., 2019). Strategies such as electrolyte modification, multifunctional barriers, composite metallic lithium electrode, and 3D current collectors were proposed to suppress the formation of Lithium dendrites (Wu et al., 2018; Xiao et al., 2018). Replacement of current organic liquid electrolytes with the solid-state electrolytes is considered as the most promising way to realize excellent performance of lithium metal secondary batteries (Goodenough, 2018; Nie et al., 2018). Because the organic liquid electrolyte is flammable, toxic and environmental contamination, solidstate electrolytes are therefore of crucial importance because of their excellent safety and mechanical ability for dendritic growth inhibition (Ban et al., 2018). And all solid-state lithium metal batteries are widely regarded as promising candidates for next generation of energy storage devices with improved energy density and superior safety performances.

Solid-state electrolytes are generally divided into inorganic superionic conductors (or the solid ceramic electrolytes) and solid polymer electrolytes. Compared with the inorganic superionic conductors, solid polymer electrolytes, benefiting from shape versatility, flexibility, light weight and low processing costs, are being investigated as promising candidates to replace currently available organic liquid electrolytes in lithium metal batteries (Bae et al., 2018b), (Zhang X. et al., 2018). However, due to the inherent low ionic conductivity at RT and poor mechanical property of solid polymer electrolytes, polymer-based solid composite electrolytes taking advantage of the merits of both inorganic and polymeric materials have been attracting more and more attention in all solid-state lithium metal battery applications (Tan et al., 2018). And Poly(ethylene oxide) (PEO) based solid composite electrolytes were widely considered as promising candidates solid electrolytes for high energy density lithium metal batteries (Xue et al., 2015; Tan et al., 2018). Numerous fillers were added into the PEO-based composite solid electrolyte to improve the ionic conductivity and the interfacial properties in contact with the electrodes (Tan et al., 2018). And it was the simplest method to achieve the composite polymer electrolyte, whose physical properties can be easily controlled by compositional change (Tan et al., 2018; Xiao et al., 2018; Zhang J. et al., 2018). However, the particle agglomeration especially at high concentration of fillers would inhibit the further improvement of properties of the solid polymer electrolyte. New kind of effective filler or percolated nanofiller structure needs to be developed by researchers.

Graphitic-like carbon nitride (g-C3N4) is a stacked 2D structure, which was mainly composed of carbon, nitrogen and a few residual –NH2 or –NH groups, and was metal-free and lightweight (Wang et al., 2012; Shi et al., 2014). And its intrinsic polarity and semiconductor property may be helpful to modulate the spatial and interface distribution of ionic charge carriers (Hu et al., 2017). These properties make g-C3N4 a potential candidate as polymer electrolyte filler for all-in-solid lithium batteries. Shi et al. proved that the introduction of 6.0 wt% g-C3N4 nanosheets in the sodium alginate (SA) nanocomposite films could greatly enhance the thermal stability and mechanical properties of SA biopolymer electrolyte nanocomposite films (Shi et al., 2014). Hu et al. reported about a kind of composite electrolyte with bis(triffuoromethanesulfonimide) lithium salt (LiTFSI), di(ethylene glycol) dimethyl ether (DGM) and g-C3N4, which was used to suppress lithium dendrite growth during cycling and improve the long-term cycling property of the composite electrolyte (Hu et al., 2017). But quantitative and even qualitative descriptions of the interactions between separate components of the state electrolyte system are still lacking. Understanding of the different interactions in the system and its influence on the microstructure and final electrochemical properties of the solid composite polymer electrolyte is crucial for achieving high performance solid polymer electrolyte.

In this article, we present for the very first time, to the best of our knowledge, the addition of g-C3N4 as fillers in PEO and its influence on the lithium ionic dissociation in lithium perchlorate (LiClO4) and transport and the crystallization of PEO in solid composite polymer electrolytes. Thermal annealing treatment and the interactions regulation among components by addition of g-C3N4 were introduced to tailor the microstructure of the solid composite polymer electrolyte. And through controlling the microstructure of solid composite polymer electrolytes, excellent improvement of the desired physical and chemical properties and electrochemical properties was achieved. When the weight percent of g-C3N4 was 10%, the annealed solid composite polymer electrolyte gave the best performance in the all solidstate lithium battery with lithium metal as the anode and LiFePO4 as the cathode.

### EXPERIMENTAL SECTION

### Materials

Poly(ethylene oxide) (Mw = 6 × 105 , Alfa Aesar), lithium perchlorate (LiClO4, Aladdin, battery grade), poly(vinylidene fluoride) (PVDF, Solef R 5130, Solvay), and LiFePO4 (DY-3, Shenzhen Dynanonic Co., Ltd) were commercially obtained and were dried before usage. Acetonitrile (ACN, Aladdin, AR), Nmethly-2-pyrrolidone (NMP, AR), and urea (Aladdin, AR) were used as obtained. The Li metal used in our experiments was commercially obtained with a diameter of 15.8 and 0.6 mm of thickness.

## Preparation of g-C3N4

The g-C3N4 filler was synthesized by thermal polymerization as we reported in the previous article (Wang et al., 2019). Twenty gram of urea was put into a crucible with a cover in the muffle furnace. And then the crucible was heated to 580◦C to stay for 4 h, with a heating rate of 3◦C/min. And the powder product was in light yellow color. The detailed characterization of pure g-C3N4 was displayed in **Figures S1**, **S3**.

### Solid Composite Polymer Electrolyte Films Preparation

Firstly, the prepared g-C3N4 filler was ultrasonic dispersed in ACN. Then PEO and LiClO4 (molar ratio of ethylene oxide (EO) to LiClO4 is 15:1) were added into the solution under stirring, in order to obtain uniform mixture solution. The weight percentage of g-C3N4 compared to PEO was controlled as 0, 5, 10, 15, and 20%, respectively. The mixture solution was casted into a homemade Teflon mold and dried in the vacuum oven at 40◦C for 48 h. And the above mentioned state polymer electrolytes films were firstly annealed at 120◦C for 2 h and then quenched to RT before use. In contrast, the state polymer electrolyte films were also prepared without thermal annealing process. The thickness of obtained solid state polymer electrolyte films was regulated to about 100µm. And the state polymer electrolytes films were cut into rounds with diameter of 19 mm before loaded into the glovebox.

### Characterization

Bruker Vertex 70 FTIR spectrometer with Access ATRTM attachment was used to characterize the composition of samples. The melting temperature of PEO in samples was collected by a differential scanning calorimeter (DSC, Mettler DSC823e) under nitrogen atmosphere. Samples with a mass of 5∼10 mg were sealed into aluminum crucible, and then heated at a rate of 10◦C/min. X-ray powder diffraction (XRD) patterns were collected by a XRD-6000 (Shimazu, Japan) X-Ray diffractometer equipped with Cu Kα radiation in the range of 5∼50◦ at a scanning rate of 4◦ /min. The reflection optical microscopy (ROM) observations were performed with a Leica DM2500P optical microscope equipped with a C-5050ZOOM camera.

The ionic conductivity of SPEs was measured by electrochemical impedance spectroscopy (EIS) by using a Princeton PARSTAT 4000A (AMETEK) electrochemical workstation over frequency range from 1 to 105 Hz with an applied voltage of 10 mV. The measured SPEs were sandwiched between two stainless steel electrodes, which were assembled into CR2025 coin cells. And the temperature was controlled by an oven in the range of 25∼120◦C. The electrochemical stability window was determined by linear sweep voltammetry (LSV) by a Princeton PARSTAT 4000A (AMETEK) electrochemical workstation. The cell was assembled by sandwiching SPEs between a stainless steel and lithium metal (SS/SPE/Li) as the counter and reference electrode, respectively. And the measurement was cycled from 2.0 to 6.0 V (VS. Li/Li+) at a scan rate of 10 mV/S. The lithium transference number (t+) was tested in a symmetric cell (Li/SPE/Li) by chronoamperometry combined with electrochemical impedance spectroscopy according to the method provided by Bruce et al. (Evans et al., 1987). In order to characterize the battery performances, the prepared state polymer electrolytes were fabricated as CR2025 coin-type cells with lithium metal anode and LiFePO4 cathode. And LiFePO4 was blended with carbon black (Super-P) and

poly(vinylidene fluoride) at a weight ratio of 8:1:1 with NMP as the solvent. The weight of cathode used in the test was controlled to about 2.40 mg. As a result, the weight of active material was about 1.92 mg. The battery performances of coin cells were examined at 80◦C using a Battery Testing System (NEWARE CT-4008, Shenzhen, China), including cycle property and C-rate capability (0.1, 0.2, 0.5, 1, 2, and 5C). The cut-off voltages were chosen as 4.2 V (charge) and 2.5 V (discharge) during the charge-discharge cycle performance.

### RESULTS AND DISCUSSION

### Microstructure Regulation by Thermal Annealing and Interaction Among Components in Solid Composite Polymer Electrolytes

In order to avoid the effect of thermal history on the crystallization of PEO during the solid composite polymer electrolyte preparation process to an extreme, 120◦C annealing and then quenching method was introduced, as shown in **Figure 1**. And we have reported on this preparation method in detail in a previous work (Wang et al., 2019). During the solid polymer electrolyte preparation process, samples were firstly dried at 40◦C for 48 h, which would provide an isothermal crystallization condition for PEO. As a result, the crystallinity of PEO in the samples was high. However, when the sample was quenching from the amorphous state, the non-isothermal crystallization might happened, which gave a lower crystallinity of PEO in the sample. And this deduction was further proved by the following XRD and DSC results.

In order to clarify the influence of thermal annealing and the interaction among components of the solid composite polymer electrolyte on the microstructure of solid composite polymer electrolytes, XRD was used to characterize the crystallization of samples as shown in **Figure 2**. As shown in **Figure 2A**a, pure g-C3N4 gave a typical diffraction peak at 27.5◦ , which was the typical peak of (002) crystal plane of g-C3N4. Moreover, the interlamellar spacing of g-C3N4 was calculated as about 0.32 nm by Bragg equation. On the other hand, XRD patterns of the solid composite polymer electrolyte samples shown in **Figures 2A**b,c displayed three typical peaks at 19.3◦ , 23.4◦ and 27.5◦ , respectively. The two characteristic peaks at 19.3◦ and 23.4◦ were corresponded to (120) and (132)/(032)/(212)/(112) crystal planes of PEO, respectively (Wang et al., 2019). It was obvious that the intensity of typical diffraction peak of PEO at

19.3◦ and 23.4◦ greatly decreased after the sample annealing at 120◦C and then quenching to RT, as shown in **Figures 2A**b,c.

Furthermore, the diffraction peaks of crystalline phase and amorphous phase in XRD patterns were fitted by the Jade 6.0 software. And the crystallization degree of PEO in samples was calculated by divided the summation of peak areas of crystallization phase and amorphous phase by the summation of peak areas of crystalline phase. Therefore, the crystallization degree of PEO was decreased from about 15.49% to about 8%, when the solid composite polymer electrolyte was thermal treated. This result demonstrated that the thermal annealing process could greatly suppress the crystallization of PEO.

As shown in **Figure 2B**, the intensity of diffraction peak observed at 27.5◦ increased with the content of g-C3N4 increasing. **Figures 2B**a–e showed that intensity of typical peaks of PEO crystals was firstly decreased and then increased due to the addition of g-C3N4, with a minimum value appeared in the solid composite polymer electrolyte with 10% g-C3N4. Moreover, when the content of g-C3N4 was chosen as 0, 5, 10, 15, and 20%, the crystallization degree of PEO was calculated as about 28, 25, 8, 19, and 20%, respectively. It was obvious that the introduction of g-C3N4 could significantly decrease the crystallinity of PEO. But the crystallization degree of PEO was firstly decreased and then increased as the content of g-C3N4 increased, with a minimum value in the solid composite polymer electrolyte with 10% g-C3N4. In the state polymer electrolyte, the porous microstructure of g-C3N4 sheets (as shown in **Figure S1**) might serve as the hard limitation and inhibit the ordered chain arrangement of PEO, while on the other hand also serve as the nucleation sites for PEO chains. As a result, the heterogeneous nucleation could be enhanced, when the content of g-C3N4 sheets was over 10%, which would improve the crystallization of PEO. Furthermore, the g-C3N4 nanosheets might change the priority growth of PEO crystal planes, as shown by the relative intensity of fitting peaks at 19.3◦ and 23.4◦ in **Figures 2B**a,e. It implied that the interaction between g-C3N4 and PEO could greatly influence the crystallization preference of crystal planes in PEO.

The effect of thermal annealing and the interaction among components of the solid composite polymer electrolyte on the crystallization of PEO was further proved by DSC. As shown in **Figures 3A**a,b the melting temperature of PEO was decreased from 52.4 to 48.7◦C, when the state polymer electrolyte was annealed at 120◦C and then quenched to RT. Also shoulder peak appeared at 46.0◦C in **Figure 3A**a. It meant that PEO crystals in the sample undergone thermal annealing formed lamellae with thinner thickness and the thickness of lamellae gave a wider distribution. It implied that the thermal annealing process could regulate the crystallization of PEO.

Furthermore, the melting behavior of PEO in PEO homopolymer, PEO-LiClO4 electrolyte and the solid composite polymer electrolytes with different content of g-C3N4 (5, 10, 15, and 20%), which were treated by the thermal annealing process, was displayed in **Figure 3B**. Due to the interaction between PEO chains and LiClO4, the melting temperature of PEO crystals decreased from 66.7 to 53.3◦C, as shown in **Figures 3B**a,b. Moreover, the addition of g-C3N4 further restrained the crystallization of PEO, as shown in **Figures 3B**c–f. The melting temperature of PEO crystals in the solid composite polymer electrolyte was 51.5, 48.7, 52.3, and 53.1◦C, respectively, with the content of g-C3N4 increasing from 5 to 20%. The melting temperature of PEO firstly decreased and then increased, with a minimum value of 48.7◦C in the sample with 10% g-C3N4. Also the crystallinity of PEO was calculated by dividing the melting enthalpy obtained from DSC curves in **Figure 3B** by the equilibrium enthalpy of PEO fusion (197 J/g). And the crystallinity of PEO in different samples shown in **Figure 3B** was about 54, 25, 22, 7, 18, and 20%, respectively. With the addition of LiClO4, the coordination effect of EO unit and Li+ would decrease the crystallization of PEO. Moreover, the introduction of g-C3N4 further decreased the crystallization of PEO. The residual –NH2 or –NH groups presenting in g-C3N4 can act as active sites for hydrogen bonding with ether oxygen in PEO.

same [EO]: [Li+] ratio of 15:1.

FIGURE 4 | ROM images of solid composite polymer electrolytes prepared with thermal annealing at 120◦C [(a) PEO-LiClO4; (b) PEO-LiClO4-5% g-C3N4; (c) PEO-LiClO4-10% g-C3N4; (d) PEO-LiClO4-15% g-C3N4; (e) PEO-LiClO4-20% g-C3N4] and without thermal annealing [(f) PEO-LiClO4-10% g-C3N4].

And the hydrogen bonds formation restrained the chain folding of PEO, and then reduced the crystallinity of PEO in the solid composite polymer electrolytes. On the other hand, the existence of g-C3N4 also served as the heterogeneous nucleation sites for PEO crystallization. As a result, the crystallinity degree of PEO firstly decreased and then increased slightly in the solid composite electrolytes, with 10% as the optimum content. This result was in good accordance with the foregoing XRD results.

The influence of thermal annealing and the interaction among components of the solid composite polymer electrolyte on the microstructure morphology of solid composite polymer electrolyte was observed by ROM. Perfect spherulites with diameter of about 100µm were formed in the solid PEO-LiClO4 polymer electrolyte, as observed in **Figure 4a**. **Figures 4b–e** showed that the morphology perfection and the crystal's size and number of PEO obviously changed with the content of g-C3N4. PEO spherulites were firstly disappeared and then re-formed with the weight percent of g-C3N4 increasing, with 10% as a critical content. It was ascribed to the hydrogen bonding interaction between g-C3N4 and PEO that could inhibit the crystallization of PEO, while g-C3N4 could also serve as the heterogeneous nucleation sites for PEO. Therefore, when the content of g-C3N4 was over 10%, more and more spherulites were formed in the sample. It implied that the interaction between PEO and g-C3N4 could greatly affect the morphology of the solid composite polymer electrolyte.

Moreover, compared **Figure 4c** with **Figure 4f**, the thermal annealing treatment might greatly change the morphology of PEO crystals from spherulites to lamellae stack. It demonstrated that the thermal treatment also influenced the morphology of the solid composite polymer electrolyte. The morphology of the solid composite polymer electrolyte was tailored by the thermal treatment and the interactions among components in the sample. And the morphology would play an important role in the compatibility between solid electrolytes and electrodes. Uniformly distributed PEO crystals might control the nucleation of lithium, which would help to inhibit the lithium dendrite formation.

### Effects of Interaction Among Components in Solid Composite Polymer Electrolytes on the Lithium Salt Dissociation

FTIR spectra of PEO, pure g-C3N4, the PEO-LiClO4 solid polymer electrolyte and the PEO-LiClO4-g-C3N4 solid composite polymer electrolyte were shown in **Figure S4**. And the typical band of LiClO4 was identified at 624 cm−1 . In order to study the effect of g-C3N4 on the dissociation of LiClO4 in the solid composite polymer electrolytes, FTIR-ATR spectra in the range of 610–650 cm−1 were carefully analyzed. It was reported that the peaks at ∼624 and ∼635 cm−1 were assigned to the dissociated free anion and the bonded ion pair in LiClO4, respectively, Lin et al. (2016). As shown in **Figure 5**, the absorbance peak of dissociated free anion shifted from 624 to 623 cm−1 , as the weight percent of g-C3N4 increased to 15%. It demonstrated that the vibration energy level transition of free ClO− 4 was easier, due to the interaction of g-C3N4 and LiClO4. It is because the residual –NH2 and –NH could coordinate with the lithium ion, which would accelerate the dissociation of LiClO4. Moreover, the relative absorbance of peaks at ∼624 and ∼635 cm−1 (labeled as A624/A635) was obtained from curves in **Figure 5**. And A624/A635 was 2.24, 3.01, 4.79, 4.34, and 4.41, respectively, with the weight percent of g-C3N4 increasing from 0 to 20%. It demonstrated that the introduction of g-C3N4 could enhance the dissociation of LiClO4, with 10% as the optimum

content. Also the residual –NH2 and –NH groups in g-C3N4 could form hydrogen bonds with PEO, which would limit the ordered folding of PEO chains and helped to enhance the chain segment motion. As a result, the dissociation of LiClO4 was accelerated, due to the coordination between Li+ and the ether oxygen. The above factors contributed to the separation of Li+ and ClO− 4 in the solid composite polymer electrolyte. And the dissociation of LiClO4 in the solid composite polymer electrolyte is critical for the good ionic conductivity.

### Ionic Conductivity and the Electrochemical Properties of the Solid Composite Polymer Electrolytes

In order to confirm our hypothesis of the increasing conductivity by thermal annealing treatment and the addition of g-C3N4, the ionic conductivity and electrochemistry properties of the solid composite polymer electrolytes were further studied. The ionic conductivity of the solid composite polymer electrolytes with different content of g-C3N4 was investigated by EIS. The Arrhenius plot of ionic conductivity with temperature from 25 to 120◦C was displayed in **Figure 6**. Compared with that of PEO-LiClO4 polymer electrolyte (**Figure 6**e), significant improvement

of ionic conductivity was observed in the annealed solid composite polymer electrolyte. Moreover, the ionic conductivity was firstly increased and then decreased at each temperature with the content of g-C3N4 increasing from 5 to 20%, as shown in **Figures 6**a–d. The solid composite polymer electrolyte with 10% g-C3N4 exhibited the maximum ionic conductivity (1.76 × 10−5 S cm−1 ) at 25◦C, and the conductivity reached to about 1.08 × 10−3 S cm−1 at 100◦C. And this ionic conductivity at room temperature was much higher than that of the recently reported PEO/g-C3N4/LiTFSI solid electrolytes (2.3 × 10−6 S cm−1 , at 30◦C) (Sun et al., 2019). Also this ionic conductivity is nearly close to that (2.2 × 10−5 S cm−1 at 28◦C) of the PEO based solid composite electrolyte, which needed LiTFSI as the lithium salt and 2D MXene as the filler (Pan et al., 2019). DSC and XRD results demonstrated that the crystallinity of PEO in the solid composite polymer electrolytes was greatly affected by the content of g-C3N4. And when the content of g-C3N4 was 10%, PEO gave the minimum crystallinity degree, which was preferred for the high ionic conductivity. Besides effects on the crystallinity of PEO, the residual –NH2 or –NH groups exited on the 2D structure of g-C3N4 could coordinate with Li+. Also the porous structure might provide a potential transport path for Li+. When the content of g-C3N4 was 10%, continuous percolating pathways for Li+ diffusion might be formed. Nonetheless, when the content of g-C3N4 was over 10%, the redundant inactive fillers would hinder lithium ionic conduction according to effective medium theory (EMT) (Li W. et al., 2018). Moreover, the modification of PEO-LiClO4 using inorganic nanoparticles fillers were always enslaved to the aggregation of particles, which was also harmful for the ionic conducting (Bae et al., 2018a), Evans et al. (1987). As a result, it demonstrated that the thermal annealed PEO-LiClO4-g-C3N4 solid composite polymer electrolyte using 10% g-C3N4 would give an easy way to achieve excellent ionic conductivity in the solid polymer electrolyte.

FIGURE 7 | Chronoamperometry of the Li/ PEO-LiClO4/Li (A) and Li/ PEO-LiClO4-10% g-C3N4/Li (B) cells at a potential of 10 mV at 80◦C. Inset: the AC impedance spectra of the same battery before and after the polarization.

Since the PEO-LiClO4-10% g-C3N4 gave the best ionic conductivity, it was chosen for the following electrochemical studies. The lithium ion transportation in the solid composite polymer electrolytes was studied as well. And the lithium ion transference number, t+, was analyzed according to the method provided by Bruce and coworkers (Evans et al., 1987). **Figure 7** showed the relation between time and current crossing a symmetric Li/PEO-LiClO4/Li battery (A) and Li/PEO-LiClO4- 10% g-C3N4/Li (B) battery polarized by a small voltage of 10 mV at 80◦C. The inset was the AC impedance spectra of the same battery before and after the polarization. And the t+ can be calculated by the equation: t+ = I s (1V−I 0R 0 ) I 0 (1V−I sRs ) , where the I 0 and I s are the initial current and the steady-state current, respectively; 1V is the potential applied across the cell, and R 0 and R s are the initial and steady-state interfacial resistances, respectively. Therefore, the t+ of the PEO-LiClO4 electrolyte was 0.241, while the t+ of the annealed PEO-LiClO4-g-C3N4 electrolyte was 0.37. It indicated that the thermal annealing treatment coupled with the addition of g-C3N4 would greatly enhance the mobility of Li+. And this result was in good accordance with the FTIR-ATR results.

Effects of g-C3N4 on the electrochemical stability of the solid composite polymer electrolytes were characterized by the electrochemical window analysis. The electrochemical stability window of PEO-LiClO4 electrolyte and annealed PEO-LiClO4-g-C3N4 electrolyte at 80◦C were measured by LSV, as shown in **Figure 8**. Detected current of PEO-LiClO4 electrolyte increased sharply as the applying voltage exceeded 3.5 V, as seen from **Figure 8**a. But **Figure 8**b showed that the detected current of PEO-LiClO4-g-C3N4 electrolyte was more stable and displayed a decomposition voltage at about 4.8 V. It implied that the annealed solid composite polymer electrolyte might guarantee a higher working voltage of electrode. And due to the interaction between g-C3N4 and PEO, the stability of the solid polymer electrolyte was greatly enhanced. It was reasonable to modify the properties of PEO electrolyte through physically thermal annealing treatment and using g-C3N4 as the filler.

The battery performance of all-solid-state lithium battery (LiFePO4/Li) based on the solid polymer electrolyte was also performed. **Figure 9** showed the long-term cycling performance and coulombic efficiencies of LiFePO4/Li batteries assembled using the annealed state polymer electrolyte with 10% g-C3N4 and PEO/LiClO4 with a current density of 1 C at 80◦C. As shown in **Figure 9A**, the cell based on PEO-LiClO4 electrolyte gave an initial discharge capacity of only 133.3 mAh g−1 , and displayed a capacity retention ratio of 80% after 115 cycles. On the other hand, **Figure 9B** demonstrated that the cell based on the solid composite polymer electrolyte with a capacity retention ratio

of 80% after 200 cycles possessed an initial discharge capacity of 161.2 mAh g−1 , which was very close to the theoretical capacity of LiFePO4 (170 mAh g−1 ). Moreover, this cycling performance was better than the reported LiFePO4/Li solid-state batteries based on PEO-LiTFSI-g-C3N4 polymer electrolytes, which gave an initial discharge specific capacity of 161.3 mAh g −1 but cycled for 120 cycles at 60◦C (Sun et al., 2019). The coulombic efficiency of batteries based on the solid composite polymer electrolyte during charge-discharge cycling test was as high as 99.7% in **Figure 9B**, which was higher than that of the battery based on the PEO-LiClO4 electrolyte shown in **Figure 9A**. Moreover, the prevailing experiments using LiFePO4 cathode in SPE based solid state lithium batteries were cycled within 3.8 V, in order to guarantee the cycle stability (Nie et al., 2018). In this article, the LiFePO4/Li solid battery could operate in the range of 4.2∼2.5 V and gave excellent cycle stability. Furthermore, the charge-discharge voltage profiles of the corresponding LiFePO4/Li solid battery during the cycling

performance were displayed as the insert in **Figures 9A,B**. The voltage plateau at around 3.48 V was steady in **Figure 9B**, while a sloping voltage plateau was observed in **Figure 9A**. And the attenuation of specific capacity during charge-discharge cycling process was more clearly in the LiFePO4/Li batteries with PEO/LiClO4 electrolyte shown in **Figure 9A**, compared with that in the LiFePO4/Li batteries with annealed PEO-LiClO4-g-C3N4 electrolyte shown in **Figure 9B**. It implied that the addition of g-C3N4 enhanced the electrochemical stability of PEO-based solid polymer electrolytes.

Moreover, in order to ascertain the usefulness of the solid composite polymer electrolyte, the C-rate cycling property of PEO-LiClO4 electrolyte and the annealed PEO-LiClO4-g-C3N4 electrolyte at 80◦C were characterized and results were shown in **Figure 10**. As shown in **Figure 10A**, the discharge specific capacity of the all solid-state battery based on the PEO-LiClO4 electrolyte was 148.0, 150.6, 139.5, 123.1, 80.4, and 22.2 mAh g −1 , respectively, with the C-rate increasing from 0.1, 0.2, 0.5C, 1, 2 to 5C. And it was obvious that the discharge specific capacity was gradually decreased at higher C-rates, which was reported as a typical characteristic of LiFePO4 material (Brutti et al., 2012). The low discharge specific capacity at 5C-rate was attributed to the electrode polarization and solid electrolyte interface caused by the low electronic conductivity, limited diffusion of Li+ into the LiFePO4 structure (Kumar et al., 2014). On the other hand, **Figure 10B** showed the performance of the battery with PEO-LiClO4-g-C3N4 electrolyte. It delivered an initial discharge specific capacity of 161.0 mAh g−1 at 0.1C-rate. And the discharge specific capacity was 159.5, 159.7, 158.8, and 155.6 mAh g−1 , respectively, at 0.2C-rate, 0.5C-rate, 1C-rate, and 2C-rate without any obvious attenuation. The battery was even able to deliver a specific capacity of 59.7 mAh g−1 even at 5C-rate, which was 168% higher than that of the battery based on the PEO-LiClO4 solid polymer electrolyte. Moreover, when the rate returned to 0.1C, the average discharge capacity of the battery was still as high as 157.2 mAh g−1 . Also the coulombic efficiencies were close to 100%. This result implied that the PEO-LiClO4-g-C3N4 solid polymer electrolyte could facilitate the cycling performance of LiFePO4 cathode, due to its good compatibility with the cathode and good lithium ionic transport capability. Moreover, the solid state battery based on the solid composite polymer electrolyte was able to light up several LED lamps under a normal condition as shown in **Figure S5**. It demonstrated that the annealed PEO-LiClO4-g-C3N4 electrolyte could be an outstanding candidate used in all solid-state lithium batteries.

### CONCLUSIONS

To conclude, we reported a novel method to prepare the solid polymer electrolyte with high performance using facile thermal annealing treatment and by addition g-C3N4 as the modification filler. XRD and DSC results demonstrated that the crystallinity degree of PEO was greatly decreased by the thermal annealing treatment and by interaction between PEO chains and the g-C3N4 filler. Moreover, the dissociation of LiClO4 was obviously enhanced as proved by the FTIR-ATR results. The final morphology of the obtained solid composite polymer electrolyte was even changed with the content of g-C3N4 increasing and changed by the thermal annealing treatment. When the weight percent of g-C3N4 was 10% in the annealed solid composite polymer electrolyte, spherulites morphology of PEO crystals was totally destroyed, which would facilitate the compatibility between the solid electrolyte and electrodes. And the crystallinity of PEO was only about 8%, which was helpful for the lithium ion transport in the solid composite polymer electrolyte.

As a result, the obtained solid composite polymer electrolyte with 10% g-C3N4 achieved an ionic conductivity as high as 1.76 × 10−5 S cm−1 at 25◦C, due to the above microstructure regulation.

### REFERENCES

Bae, J., Li, Y., Zhao, F., Zhou, X., Ding, Y., and Yu, G. (2018b). Designing 3D nanostructured garnet frameworks for enhancing ionic conductivity and The electrochemical stable window was increased to about 4.8 V. The lithium ion transference number, t+, was changed from 0.24 to 0.37, and increased about 54.8%. Finally, the annealed PEO-LiClO4-10% g-C3N4 electrolyte was also assembled into an all solid-state battery using LiFePO4 as the cathode and lithium metal as the anode. And the battery demonstrated excellent performance, with a high initial discharge capacity of 161.2 mAh·g −1 and superior cycle stability with a capacity retention ratio of 81% after 200 cycles at 1C at 80◦C. Furthermore, the C-rate cycling results also demonstrated that the stability of the assembled solid state battery was markedly improved by using PEO-LiClO4-10% g-C3N4 electrolyte. The above results indicates that the thermal annealing treatment and g-C3N4 as a novel structure modifier provide an useful and facile way to prepare the state polymer electrolytes used in the all solid-state lithium battery.

### DATA AVAILABILITY

The raw data supporting the conclusions of this manuscript will be made available by the authors, without undue reservation, to any qualified researcher.

### AUTHOR CONTRIBUTIONS

LS and JY developed the concept and designed the experiment. They contributed equally to this work. XW and JY conducted the experiments. KX, GZ, and AM co-supervised the experiments. JY and LS wrote the manuscript. KX, CS, and WC helped to revise the work critically. All authors listed have made a substantial, crucial and direct contribution to the work, and approved it for publication.

### FUNDING

This work was supported by the project sponsored by the Key R&D Program of Shaanxi (2017ZDCXL-GY-08-03), the National Natural Science Foundation of China (61604120, 51804259 and 21501139), Scientific Research Program Funded by Shaanxi Provincial Education Department (17JK0380 and 18JK0390), Natural Science Foundation of the province of Shaanxi (No.2018JQ2037), Innovation and entrepreneurship training program for college students<PHONE_NUMBER>) and the Presidential Research Fund of Xi'an Technological University (No. XAGDXJJ18007).

### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2019.00388/full#supplementary-material

flexibility in composite polymer electrolytes for lithium batteries. Energy Storage Mater. 15, 46–52. doi: 10.1016/j.ensm.2018.03.016

Bae, J., Li, Y., Zhang, J., Zhou, X., Zhao, F., Shi, Y., et al. (2018a). A 3D nanostructured hydrogel framework-derived high-performance composite polymer lithium-ion electrolyte. Angew. Chem. Int. Ed. 57, 1–6. doi: 10.1002/anie.201710841/

electrolyte nanocomposite films, a novel investigation. ACS Appl. Mater. Interf. 6:429. doi: 10.1021/am4044932

**Conflict of Interest Statement:** LS was employed by the company Shaanxi Coal Chemical Industry Technology Research Institute Co. Ltd.

The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Yang, Wang, Zhang, Ma, Chen, Shao, Shen and Xie. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Advances in Microwave-Assisted Production of Reduced Graphene Oxide

Xinxin Xie1,2, Yanping Zhou1,2 \* and Kama Huang1,2

*1 College of Electronics and Information Engineering, Sichuan University, Chengdu, China, 2 Key Laboratory of Wireless Power Transmission of Ministry of Education, Sichuan University, Chengdu, China*

Efficient reduction of graphene oxide to obtain high-quality graphene nanosheets is desirable for energy storage, catalysis, electronics and environmental remediation. In this brief review, we mainly focus on the microwave-assisted production of reduced graphene oxide in three categories: (1) microwave-assisted chemical reduction of graphene oxide; (2) microwave-assisted thermal reduction of graphene oxide; (3) microwave-assisted simultaneous thermal exfoliation & thermal reduction of graphite oxide. We also summarize common techniques for characterizing reduction efficiency and quality of as-obtained rGO.

Keywords: microwave-assisted synthesis, reduced graphene oxide, chemical reduction, thermal reduction, exfoliation of graphite oxide

#### Edited by:

*Wenping Sun, University of Wollongong, Australia*

#### Reviewed by:

*Jiantie Xu, South China University of Technology, China Chao Wu, University of Wollongong, Australia*

> \*Correspondence: *Yanping Zhou<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Electrochemistry, a section of the journal Frontiers in Chemistry*

Received: *21 February 2019* Accepted: *30 April 2019* Published: *04 June 2019*

#### Citation:

*Xie X, Zhou Y and Huang K (2019) Advances in Microwave-Assisted Production of Reduced Graphene Oxide. Front. Chem. 7:355. doi: 10.3389/fchem.2019.00355* INTRODUCTION

Graphene is a two-dimensional sheet of sp2-hybridized carbon.  Clearly, there will be an increasing demand for high quality graphene.

Currently, reduction of graphene oxide or simultaneous exfoliation and reduction of graphite oxide are regarded to be most promising for large scale production of chemically derived graphene, which can be defined as reduced graphene oxide (rGO). In general, there are mainly two strategies used for graphene oxide reduction: chemical reduction and thermal reduction, both of which are facing certain challenges. In chemical reduction methods, the reduction efficiency is unsatisfactory, requiring long treating time and yielding rGO with oxygen content usually higher than 15% (Stankovich et al., 2007; Shao et al., 2010; Mattevi et al., 2011; Park et al., 2011b; Xiao et al., 2013; Zhang et al., 2013; Wen et al., 2014). In thermal reduction process, long-time high-temperature treatment is not only energy consuming but also results in fragmentation of graphene and lots of structural defects (Jones et al., 2002; Chen and Yan, 2010; Chen et al., 2010; Gao et al., 2010; Lin et al., 2010; Zhu et al., 2010b; Zhang et al., 2011).

Microwave heats materials directly through dielectric loss rather than heat convection as in the conventional heating method, and hence can enable fast heating and selective heating. It is regarded to be promising in shortening the reaction time and yielding hot spots with extraordinary high temperature (Thostenson and Chou, 1999; Kappe, 2004; Tompsett et al., 2006; Schwenke et al., 2015). As such, material scientists have devoted much effort to explore the potential application of microwave in overcoming the current problems and facilitating the fabrication of high-quality rGO in large-scale. To this end, we present an up-to-date critical survey of literature relevant to MW-assisted production of rGO. This brief review is organized as follows. First, techniques for characterizing quality of asobtained rGO are discussed. Next, we provide a comprehensive summary of microwave-assisted chemical reduction of graphene oxide, microwave-assisted thermal reduction of graphene oxide, and microwave-assisted simultaneous thermal exfoliation and reduction of graphite oxide**,** respectively.

### TECHNIQUES FOR CHARACTERIZING QUALITY OF RGO

Generally, X-ray diffraction, X-ray photoelectron spectroscopy, Raman spectroscopy, electrical conductivity, thermal gravity analysis are common techniques used for characterizing reducing efficiency and quality of rGO.

### X-Ray Diffraction

The X-ray diffraction (XRD) pattern of graphene/graphite oxide usually shows a characteristic diffraction peak at 2θ = 10.9◦ with the corresponding d-spacing of 0.81 nm (Stobinski et al., 2014; Strankowski et al., 2016), due to insertion of hydroxyl and epoxy groups between the carbon sheets. After reduction, two types of XRD spectra may be obtained: (1) The XRD pattern with no peaks, suggesting the complete reduction and the formation of single-layered rGO (Hassan et al., 2009). (2) The XRD pattern showing a broad peak around 26.5◦ (Zedan et al., 2010; Liu et al., 2011; Park et al., 2011b; Wang et al., 2011; Pokharel et al., 2014). Since the oxygen-containing functional groups could lead to a reduction in crystallinity, the higher the broad peak's intensity is, the higher the crystallinity degree is, and the better graphene/graphite oxide is reduced.

### X-Ray Photoelectron Spectroscopy

X-ray photoelectron spectroscopy (XPS) could help to obtain qualitative and quantitative analysis results of the content of oxygenated groups. For example, the carbon/oxygen atomic ratio (C/O ratio) could be evaluated from the areas of C1s and O1s peaks and the atomic sensitivity factor. Usually, the C/O ratio of graphene oxide is ∼2.2–2.7 (Stankovich et al., 2007; Chen et al., 2010; Park et al., 2011a; Li et al., 2013; Wen et al., 2014; Han et al., 2015), attributed to the abundant oxygen-containing functional groups introduced during oxidation of graphite. Further, the C1s peak can be decomposed into three peaks: C=C (284.7 eV), C-O (286.9 eV), C=O (287.77 eV) (Zhao et al., 2014), through the areas of which existing statuses of oxygen could be analyzed in a quantitative manner.

#### Raman Spectroscopy

The D band around 1,350 cm−1 in Raman spectra could reflect the disorder degree of the crystal structure of carbon while the G band around 1,580 cm−1 represents a first-order scattering E2g vibration mode for characterizing the sp2 bond structure of carbon. ID/IG is the ratio of D band intensity to G band intensity, which could be used to evaluate the quality of the graphene structure. The higher ID/IG is, the more defects of C atom crystal there are. In the Raman spectrum of graphene, the D-band peak could not be observed. However, the ID/IG of graphene oxide could be as high as ∼0.8–1 (Chen et al., 2010; Lin et al., 2010; Liu et al., 2011; Park et al., 2011a; Zhao et al., 2014). Besides, the number of rGO layers can be estimated by observing the shape and position of the 2D peak in Raman spectra which is the second-order two-phonon process. For example, single-layer graphene exhibits a single, sharp 2D band located below 2,700 cm−1 , while bilayer sheets have a broader 2D peak around 2,700 cm−1 , and sheets more than five layers have a broad 2D peak above 2,700 cm−1 (Hassan et al., 2009).

### Electrical Conductivity

Though the conductivity of graphene is as high as 1×108 S m−1 , that of graphene oxide is about 0.02–0.07 S m−1 (Li et al., 2010; Zhu et al., 2010a; Park et al., 2011a) due to the existence of oxygen-containing functional groups which disrupted sp2 bonding networks. During the reduction process, electrical conductivity can be recovered by restoring the p-network as oxygen-containing groups leave, leading to a conductivity improvement by 3–6 orders of magnitude (Chen et al., 2010; Dreyer et al., 2010; Zhao et al., 2014; Han et al., 2015). Therefore, the restoration degree of conductivity can indirectly reflect the reduction degree and quality of rGO.

### Thermal Gravity Analysis

Graphene/Graphite has a very good thermal stability even when being heated up to 900◦C (Bastiurea et al., 2015). For graphene oxide, it starts to have significant mass loss from ∼200◦C due to the decomposition of labile oxygen functional groups leads to poor thermal stability (Hassan et al., 2009; Chen et al., 2010; Zhu et al., 2010a; Gannavarapu et al., 2018). Therefore, the removal of oxygen-containing functional groups can help to restore its thermal stability. The significant mass loss temperature of rGO usually happens around 500–800◦C, depending on the reduction degree of rGO.

### MICROWAVE-ASSISTED PRODUCTION OF REDUCED GRAPHENE OXIDE

Methods for producing rGO could generally be categorized into chemical reduction of graphene oxide, thermal reduction of graphene oxide, and simultaneous thermal exfoliation and reduction of graphite oxide. Here, we would also discuss microwave-assisted production of rGO in these three categories separately.

### Microwave-Assisted Chemical Reduction of Graphene Oxide

Strong reducing reagents have been widely used to reduce graphene oxide. However, the reduction process is very slow. For example, hydrazine hydrate enabled reduction of graphene oxide in the oil bath for 12 h and 24 h resulted in a C/O ratio about 5 (Park et al., 2011b) and 10.3 (Stankovich et al., 2007), respectively. Microwave irradiation was found to be capable of accelerating the reducing rate significantly. For example, Hassan et al. (2009) reported microwave assisted hydrazine hydrate reduction of graphene oxide. After microwave treatment for 2 × 30 s (on for 10 s, off and stirring for 20 s), the ID/IG in Raman spectra was 0.1–0.12, indicating a high reduction degree even after such a short treating time This was confirmed by the improved thermal stability, showing no significant mass loss up to 750◦C. Elazab et al. (2015) obtained Pd/Fe3O4 nanoparticles supported on graphene nanosheets by a one-pot microwave heating with the existence of hydrazine hydrate. After microwave irradiation for 2 min, the C/O of rGO was determined to be 8.1. Kumar et al. (2015) reported microwave-enhanced chemical reduction of graphene oxide using HI/CH3COOH as reducing reagent. They found the effect of 4 h microwave irradiation was comparable to 48 h conventional reaction process.

Since strong reducing regents are very much expensive and hazardous, organic solvents have also been explored to reduce graphene oxide. Zedan et al. (2010) reported MW-assisted reduction of graphene oxide in dimethyl sulfoxide (DMSO), a solvent with high microwave absorbing capability. The reduction degree of the sample prepared from 2 min MW irradiation at 1,200 W is even much higher than that of the sample derived from 7 h conventional heating treatment, as evidenced by XRD results of both samples shown in **Figures 1a,b**. After 2 min of MW irradiation in DMSO, the color of the product changed to black (**Figure 1c**), and the XRD pattern showed no characteristic peak (10.9◦ ) of GO, indicating that most of the GO has been converted into rGO. In contrast, under conventional heating at 180 ◦C, an apparent peak located at 10.9◦ was found even after 7 h treatment. It was not until 12th h that the GO was completely reduced, as can be seen from the disappearance of the peak at 10.9◦ in the XRD pattern and the color change of the sample to black (**Figure 1d**). Chen et al. (2010)reported the microwave-enhanced chemical reduction of graphene oxide to rGO in a mixed solution of N, N-dimethylacetamide (DMAc) and water without any additional reductant under nitrogen purge. The XRD pattern of the graphene oxide showed a wide peak, indicating the damage of the regular crystalline pattern of graphite during the oxidation. After microwave treatment for 2 min × 5 times (2.45 GHz, 800 W), no peak was observed indicating microwave-assisted formation of single-layered graphene. XPS results showed that C/O ratio changed from 2.09 to 5.46, confirming that the removal of oxygen containing groups happened. However, the ID/IG in Raman spectra changed from 0.95 to 0.96, suggesting that the reduction led to an increase of aromatic domains of smaller overall size in graphene. Similarly, Liu et al. (2011) also reported an increase of the product's ID/IG in Raman spectra from 0.80 to 0.94 after being treated by microwave at 750 W for 2 min in N, N-dimethylacetamide (DMF) solution. Zhao et al. (2014) applied DMF to reduce graphite oxide under microwave irradiation at 100 W for 30 min. The XPS results showed that C/O ratio rose from 2.13 to 6.13 while Raman results showed that ID/IG remained almost unchanged.

The characteristics of the as-obtained rGO products derived from different microwave-assisted chemical reduction methods are summarized and listed in **Table 1**. Since microwave heat rapidly, microwave-assisted rate enhancement in chemical reduction of graphene oxide has been widely observed, with most of the reaction time <10 min. However, as in the conventional heating method, the as-obtained rGO is still unsatisfactory, facing the same problem of containing a high oxygen content and a high ID/IG value when organic solvent was used as the reducing reagent. Microwave-enhanced quality of chemically reduced rGO is yet to be explored. Also, since the type and amount of reducing agents were different in different work, it is difficult to draw a conclusion regarding the effect of microwave power and treating time here.

### Microwave-Assisted Thermal Reduction of Graphene Oxide

In conventional thermal reduction process, long-time hightemperature treatment in a protecting atmosphere (e.g., N2/Ar) is usually required, which is not only very energy consuming but also leads to the formation of defects in the graphene basal plane because of the evolution of the oxygen functional groups during reduction, like nanoscopic holes brought by carbon loss as CO or CO2 and Stone-Wales types of defects lead by rearrangement of carbon atoms in the graphene basal plane. Besides, the reducing efficiency is always suppressed by the highly stable ether and carbonyl groups formed between the oxygen functional groups.

In 2010, Li et al. (2010) for the first time reported a microwaveassisted thermal reduction of graphene oxide. Specifically, they made the graphene oxide into a free-standing film via filtration membrane, and then treated it with a microwave irradiation with a frequency of 6.425 ± 1.150 GHz at a power of 500 W without purging any inert gas. As shown in **Figures 2a,b**, after 1 s irradiation, the temperature rose only a little and no distinct difference were found in the electrical properties and the Raman spectra of the sample. However, after 2 s irradiation, the surface temperature of the graphene oxide film ramped up to 400◦C and the conductivity of the graphene oxide film increased from 0.07 to 1 × 104 S m−1 ; the ID/IG in Raman spectra decreased from 1 to 0.3; the mass loss temperature increased from 181 to 668◦C (**Figure 2c**). In addition, they found that in the MW treatment process, fierce degassing and a high gas pressure inside the graphene oxide was generated, which was beneficial for exfoliating the graphene oxide into thinner graphene sheets. Afterwards, they also applied a household MW oven (with a frequency of 2.45 GHz and a power of 1,000 W) to treat the free-standing graphene oxide film. They found that there was an optimal irradiation time of 8 min as evidenced by Raman spectra (**Figure 2d**) and TGA results (**Figure 2e**). Longer treatment time (e.g., 10 min) caused lower mass loss temperature (from 700◦C to around 800◦C), which

could possibly be attributed to the oxidation by air when there was no protecting gas.

Considering the weak microwave absorption of graphene at low frequency, while 2.45 GHz and 915 MHz are most commonly used in industry, Voiry et al. (2016) introduced a conventional mild annealing process before the microwave irradiation process. Firstly, the single-layered graphene oxide solution was slowly injected into an aqueous solution that contained 1 wt% CaCl2, which helps to coagulate GO nanosheets and promote solidification via bridging with the oxygen-containing groups from two GO nanosheets, yielding uniform and continuous gelatinous GO. The coagulated graphene oxide was then washed with DI water and dried, which was subsequently annealed at 300◦C for 1 h under Argon to yield mildly reduced graphene oxide that was capable of absorbing microwaves at 2.45 GHz effectively. After 1–2 s long microwave irradiation in microwave oven at 1,000 W in argon, the efficient absorption of microwave led to rapid heating of the mildly reduced GO and consequently brought large arcing as shown in **Figure 3a**, causing desorption of oxygen functional groups and reordering of the graphene basal plane. As such, the as-obtained rGO exhibited pristine CVD graphene-like features in the Raman spectrum showing sharp G an 2D peaks and a nearly absence of D peak. In contrast, the rGO derived from conventional thermal reduction remains highly disordered, as indicated by the presence of an intense and broad disorder D band and the absence of the 2D band in the Raman spectra (**Figure 3b**). The XPS results were well in agreement with the Raman results. Owning to the large arcing which may possibly resulted in an extremely high temperature, the oxygen content of the MWrGO was about 4%, lower than that theoretically predicted for rGO after annealing at 1,500 K (15–25%). Aberration corrected HR-TEM (**Figure 3c**) also suggested a highly ordered structure, which suggests that there is some reorganization of the carbon bonding during microwave reduction, along with removal of oxygen facilitated by achieving exceptionally high temperatures.

Jiang et al. (2018) further modified Voiry's strategy and developed a triggered microwave-assisted reduction of graphene oxide. In their work, 1 wt% CaCl2 aqueous solution was used to treat one side of the filter paper while graphene oxide was dropped onto the other side of the filter paper. Subsequently, TABLE 1 | Characteristics of rGO products derived from different microwave-assisted chemical reduction.

graphene oxide was treated with CaCl2 solution and then washed with DI water, dried in air and pulled out from the filter paper. Finally, a small piece of rGO "paper" obtained by thermal annealing was placed on the as-obtained large graphene oxide paper to act as a trigger, as shown in **Figure 4a**. After 800 W microwave irradiation for 2 s in air, arc discharge process started. The irradiation was maintained for extra 3– 5 s and high-quality microwave reduced graphene oxide paper was obtained. In XPS spectrum, the characteristic peak for C-O vanished, suggesting the significant removal efficiency of oxygencontaining group. Raman spectra of MW-rGO showed sharp G and 2D bands, and low-intensity D band. In contrast, the Raman spectra of rGO derived from conventional thermal reduction at 800◦C showed a strong D band and none 2D band, indicating microwave heating is very beneficial for recovering the highly ordered graphene-like structures which was also confirmed by the conductivity. The sheet resistance of the MW-rGO was measured to be about 40 cm−2 , far less than that (about 796 cm−2 ) of rGO prepared by 800◦C annealing. They further examined the Raman spectra of graphene-triggered microwave reduction at different positions, and found the part of graphene oxide located nearest to the trigger was not reduced the best, as shown in **Figure 4b.** Hence, they proposed that the microwave reduction can be not only governed by a simple thermal

reduction process but the result of the thermal effect and arc discharge (Jiang et al., 2018).

 Microwave treatment is applied for 2 s one time. The decreasing D band in Raman spectra under different microwave treating times indicate that high crystalline graphene was obtained after microwave treatment.

Low-temperature reduction of graphene oxide is of great importance when graphene oxide film was deposited on glass or plastic substrates. Han et al. (2015) demonstrated microwaveassisted low-temperature thermal reduction of few-layered graphene oxide film. Specifically, they put the graphene oxide film at the center of the highest magnetic field in a single mode microwave reactor which was operated at 2.45 GHz with a power of 42 W to ensure that the temperature was below 250◦C during the whole process. After 5 min treatment in the air, ID/IG in the Raman spectrum of the product decreased from 2.89 to 1.56 and C/O derived from XPS results increased from 7.8 to 17.5. In addition, microwave-enabled exfoliation leading to thinner pieces was also observed in their work.

The characteristics of the as-obtained rGO products derived from different microwave-assisted thermal reduction of graphene oxide methods are summarized and listed in **Table 2**. Arc discharge caused by microwave irradiation is very beneficial for obtaining high-quality rGO, as evidence by the short reaction time, the low ID/IG values, the high C/O values and the conductivity values of most as-derived rGO. It is worth to note that rGO containing only 4% oxygen and exhibiting pristine CVD graphene-like features in the Raman spectrum was successfully fabricated via this reducing strategy. Besides, owning to the ultra-rapid treating process, protecting atmosphere is not always necessary. However, too long treatment time would result in re-oxidation and consequently rGO with worse quality when

the reduction is done in the air. Since the amounts of GO samples and the microwave heating chambers in different work were different, which would affect the power intensity imposed on the GO, it is also difficult to draw a conclusion regarding the effect of microwave power here.

permission from Wan et al. (2018), copyright 2018 John Wiley and Sons].

### Microwave-Assisted Simultaneous Exfoliation and Reduction of Graphite Oxide

Beside of direct reduction of graphene oxide to yield rGO, thermal shock like flash-assisted simultaneous exfoliation and reduction of graphite oxide at temperatures up to 1,050◦C has also been used. Theoretically, microwave heating which is well-known for its rapid heating feature is very promising in enabling simultaneous exfoliation and reduction of graphite oxide. In 2010, Zhu et al. (2010a) treated graphite oxide powder in microwave oven at 700 W for 1 min in ambient atmosphere. Large volume expansion as shown in **Figures 6a**,**b** accompanied with violent fuming followed by sparking was observed. They found that a minimum power of 280 W was required to expand the graphite oxide powder successfully, which was probably due to the insulating feature that resulted in inferior microwave absorption ability. After microwave treatment, the as-obtained rGO had a worm-like morphology, in which many regions are wrinkled and folded (**Figures 6c,d**). XPS results showed that C/O increased from 0.79 to 2.75 (**Figure 6e**). XRD spectrum showed a broad peak centered at 2θ = 25◦ . The conductivity was measured to be about 274 S m−1 .

Under microwave irradiation, graphite oxide itself might not be heated to a high enough temperature to cause its exfoliation and reduction efficiently due to its inferior microwave absorption ability.

Park et al. (2011a) mixed graphite oxide powder with 10% graphene nanosheets to realized efficient exfoliation and reduction under microwave irradiation. The temperature rising rate was estimated to be 2,000◦C/min, which was high enough to build up pressure to overcome the van der Waals forces between the graphene sheets in graphite oxide for exfoliation. Huge volume expansion was observed after 10 s, followed by arc after 40 s. SEM images (**Figure 7**) showed that the prepared rGO also had a worm-like structure composed of

ultra-thin sheets, which is different from randomly agglomerated and wrinkled sheets prepared from solution-based reduction syntheses. HR-TEM images suggested the presence of single or a few layers of graphene nanosheets as shown in **Figure 7**. They further compared the effect of reduction atmosphere on the product quality. The rGO prepared in air, Ar, H2+Ar (1:9) showed a C/O ratio of 3.6, 11.45, 18.5, respectively and ID/IG values of 0.853, 0.842, and 0.785, respectively in Raman spectra, respectively. The rGO prepared in Ar, H2+Ar showed a conductivity of 7.41 × 102 and 1.25 × 103 S m−1 , respectively. Hydrogen containing atmosphere is the most effective because H2O was formed instead of CO2, which prevents the formation of vacancies and defects due to the loss of carbon atoms.

Wen et al. (2014) applied a sequential chemical reduction and microwave exfoliation and reduction of graphite oxide to

TABLE 2 | Characteristics of rGO products obtained by different microwave-assisted thermal reduction.

TABLE 3 | Characteristics of rGO products derived from different microwave-assisted simultaneous thermal exfoliation and reduction.

produce rGO. Specifically, graphite oxide powder was treated by NaBH4 at 95◦C for 2 h, and subsequently made into a film, which was microwave irradiated at 5.8 GHz with a power of 1,400 W for 10 min in N2. After NaBH4 treatment, a C/O value of 3.84, an ID/IG value of 1.05 in Raman, and a conductivity of 1.41 × 102 S m−1 were obtained. During the microwave treatment, the sample expanded to a loose floc and transmitted bright orange light, leading to ultrathin sheets with a C/O ratio of 10.33, an ID/IG ratio of 1.07 in the Raman spectrum and a conductivity of 1.19 × 104 S m−1 . Sole microwave irradiation at 5.8 GHz with a power of 1,400 W for 10 min without a preceding chemical reduction process gave yield to a C/O of 7.24 and an ID/IG of 0.89 in Raman spectrum. This demonstrated that solid-state microwave irradiation not only removed oxygen-containing groups, but also repaired defects on the graphene sheets while NaBH4 though effective for removing oxygen-containing groups may induce more defects. Further, through XPS analysis of content of C-O and C=O groups, they found that microwave irradiation removed C=O more readily than C-O while the chemical reduction method removed C-O more efficiently.

Pokharel et al. (2014) developed a multi-step reduction technique to obtain rGO from graphite oxide. In the first step, graphite oxide with a C/O of 0.83 was exfoliated within 1 min under microwave irradiation at 2.45 GHz with a power of 1,000 W in N2 without becoming red. The expanded graphite

microwave irradiation synthesis. (f) HR-TEM images of GNS. [Reprinted with permission from Park et al. (2011a), copyright 2011 The Royal Society of Chemistry].

oxide had a C/O of 2.42. However, it was so loose that microwave could not concentrate on a definite area for further reduction of graphite oxide. Hence, in the second step, the expanded graphite oxide was collected and then exposed to additional microwave irradiation for 1 and 2 min, resulting in rGO showing a C/O ratio of 5.78 and 10.39, respectively.

The characteristics of the as-obtained rGO products derived from different microwave-assisted simultaneous thermal exfoliation and reduction of graphite oxide methods are summarized and listed in **Table 3**. As can be seen from the C/O values, ID/IG values, mass loss temperatures and conductivities, simultaneous exfoliation and reduction was not able to yield rGO with as high quality as those obtained from thermal reduction of graphene oxides. However, when combined with reducing atmosphere (i.e., H2) or pretreatment with strong reducing reagents (NaBH4), the reduction degree could be greatly enhanced.

### PERSPECTIVES

It is expected that the explored methods for microwave-assisted reduction of graphene/graphite oxide into high-quality rGO could be extended to fabrication of graphene composites. The microwave-induced arc facilitated short reaction time at high temperature is not only beneficial for getting high-quality rGO, but also promising in facilitating the formation of highquality nanomaterials (such as metal nanocrystals, metal oxides nanocrystals, and so on) with advanced performance. In addition, the methods could be extended to fabrication of other conductive materials that could trigger sparks under microwave irradiation, like graphene does.

### AUTHOR CONTRIBUTIONS

All authors listed have made a substantial, direct and intellectual contribution to the work, and approved it for publication.

#### FUNDING

This work was financially supported by the National Natural Science Foundation of China (Grant No.

#### REFERENCES

61801314), National Natural Science Foundation of China (Grant No. 61731013), Fundamental Research Funds for the Central Universities of China (Grant No. YJ201703).

supporting metal nanocrystals in aqueous and organic media. J. Mater. Chem. 19, 3832–3837. doi: 10.1039/B906253J

**Conflict of Interest Statement:** The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Copyright © 2019 Xie, Zhou and Huang. This is an open-access article distributed under the terms of the Creative Commons Attribution License (CC BY). The use, distribution or reproduction in other forums is permitted, provided the original author(s) and the copyright owner(s) are credited and that the original publication in this journal is cited, in accordance with accepted academic practice. No use, distribution or reproduction is permitted which does not comply with these terms.

# Functional Ionic Liquid Modified Core-Shell Structured Fibrous Gel Polymer Electrolyte for Safe and Efficient Fast Charging Lithium-Ion Batteries

Xiaoxia Liu, Yufei Ren, Lan Zhang\* and Suojiang Zhang\*

*Beijing Key Laboratory of Ionic Liquids Clean Process, CAS Key Laboratory of Green Process and Engineering, Institute of Process Engineering, Chinese Academy of Sciences, Beijing, China*

#### Edited by:

*Oleksandr Malyi, University of Oslo, Norway*

#### Reviewed by:

*Xianhong Rui, Guangdong University of Technology, China Weimin Kang, Tianjin Polytechnic University, China Huarong Xia, Nanyang Technological University, Singapore*

#### \*Correspondence:

*Lan Zhang<EMAIL_ADDRESS>Suojiang Zhang<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Electrochemistry, a section of the journal Frontiers in Chemistry*

Received: *16 March 2019* Accepted: *21 May 2019* Published: *12 June 2019*

#### Citation:

*Liu X, Ren Y, Zhang L and Zhang S (2019) Functional Ionic Liquid Modified Core-Shell Structured Fibrous Gel Polymer Electrolyte for Safe and Efficient Fast Charging Lithium-Ion Batteries. Front. Chem. 7:421. doi: 10.3389/fchem.2019.00421* Fast charging is of enormous concerns in the development of power batteries, while the low conductivity and lithium ion transference number in current electrolytes degraded the charge balance, limited the rate performance, and even cause safety issues for dendrite growth. Combine inorganic fillers and ionic liquid plasticizer, here in this paper we prepared a core-shell structured nanofibrous membrane, by incorporating with carbonate based electrolyte, a gel polymer electrolyte (GPE) with high conductivity, outstanding Li+ transference number was obtained. Notably, the Li/electrolyte/LiNi0.6Co0.2Mn0.2O2 (NCM622) half-cell with this composite electrolyte delivers a reversible capacity of 65 mAh/g at 20C, which is 13 times higher than that of with Celgard 2325 membrane. It also shows enhanced long-term cycle stability at both 3C and 5C for the suppression of lithium dendrite. This organic-inorganic co-modified GPE guarantees the fast charging ability and safety of LIBs, thus provides a promising method in high performance electrolyte design.

Keywords: lithium ion battery, fast charging, core-shell structure, nanofibrous separator, safety

### INTRODUCTION

With the increasing demand for energy consumption and the continuous improvement for environmental protection, the development of sustainable and renewable energy is urgently needed currently. Lithium ion batteries (LIBs) have been widely used in the power battery, energy storage applications, portable electronic equipment and other fields, owing to its inherent advantages such as high energy density, high power density, long cycle life, no memory effect, and environmental friendliness (Gao et al., 2009; Goodenough and Kim, 2010; Etacheri et al., 2011; Scrosati et al., 2011; Lu et al., 2014a; Varzi et al., 2016; Tan et al., 2018). Nevertheless, high-capacity and high-power LIBs are facing more and more challenges such as high power output, rapid charge/discharge capability and safety under dynamic conditions (Zhu et al., 2013; Xie et al., 2014; Zhang J. et al., 2014; Li et al., 2018, 2019), and it is of particularly importance to investigate methods to guarantee the safety of LIBs with enhanced fast charge capability (Rui et al., 2016; Tan et al., 2018).

LIBs generally compose of anode, cathode, electrolyte, and packing material. The electrolyte, which including polymer separator and liquid electrolyte in commercial LIBs, mainly plays two key roles: blocking the direct contact between the electrodes to avoid internal electrical short circuit and acting as the Li+ transport passage (Tarascon and Armand, 2001; Xiao et al., 2009; Huang, 2010; Lee et al., 2014). Therefore, the electrolyte not only directly affects the LIB electrochemical performance, such as cycle stability and rate capability, but also, it has significant influence on the safety of LIBs. Most commercial LIBs adopt polyolefin [such as polyethylene (PE) and polypropylene (PP)] membranes (Huang, 2010; Kim et al., 2015; Shen et al., 2018) and LiPF6/carbonates electrolyte, based on which satisfying cycle performances have achieved, while the rate capability and safety, especially during quick charge, is yet to be enhanced.

The major issue in fast charging LIBs is the poor kinetics caused by the sluggish ion transport in the cell, which will probably lead to large polarization and low reversible capacity, furthermore, when the anode potential is lower than that of Li plating, dendrite will be generated and may short the cell thus cause safety concerns.

Electrolyte holds significant influence on both rate capability and safety, on one hand, ionic conductivity and Li+ transference number is the major factor to control the battery kinetics (Diederichsen et al., 2017), on the other hand, the separator, which blocks the direct contact between electrodes, should also work as a physical barrier to suppress the dendrite from shorting the battery (Jana et al., 2015; Shin et al., 2015; Zheng et al., 2018).  Thereinto, the nanofibrous membrane has large specific surface area, three-dimensional (3D) porous structure and high electrolyte uptake ability, on which the rate capability can be enhanced (Xiao et al., 2015; Liang et al., 2016; Park et al., 2017). What's more, functional groups can be easily introduced in the electrospining process, thereby endows the electrolyte with other properties, such as dendrite suppression (Lu et al., 2015a,b; Cheng et al., 2018; Deng et al., 2018), flame retardant (Lu et al., 2017; Jia et al., 2018; Sun et al., 2018), et al. Poly(vinylidene fluoride-co-hexafluoropropene) (PVDF-HFP) is considered as a potential polymer matrix for electrolyte not only as it has a low crystallinity which could promotes rapid ion conduction (Ali et al., 2018; Zhao et al., 2019), but also, it has high dielectric constant and low surface energy that may promote the compact deposition of metal Li (Lopez et al., 2018). It is reported the ionic conductivity of PVDF-HFP-based polymer electrolytes is up to 3.9 mS/cm (Xiao et al., 2012).

Ionic liquid modified gel polymer electrolytes (IL-GPE) have attracted much attention due to their good thermal stability and mechanical properties. Singh et al. studied imidazolyl ionic liquids and found that the EMIMFSI-based GPEs have excellent electrochemical stability, good compatibility and thermal stability(Singh et al., 2018). Guo et al. prepared the PVDF-HFP-LiTFSI/SiO2/EMITFSI GPE with high thermal stability and good electrode compatibility (Guo et al., 2018). However, IL-GPE also has some shortcomings, such as lower ionic conductivity and lithium ion mobility (Zhou et al., 2015). In this work, a piperidine ionic liquid, 1-methyl-1-propylpiperidinium chloride (PPCl), and Li2SiO3 (LSO) nanoparticles were introduced into the PVDF-HFP matrix via coaxial electrospinning technology. The as prepared membrane has high porosity and electrolyte uptake, remarkable ionic conductivity, and outstanding electrochemical performance especially in quick charging. Commercial NCM622 cathode adopting this IL-GPE delivers a high reversible capacity of 65 mAh/g in 20C rate charge/discharging, which is 13 times higher than that of the cell adopting Celgard 2325 membrane.

### EXPERIMENTAL SECTION

### Materials

Poly (vinylidene fluoride-co-hexafluoropropene) (PVDF-HFP, Mw. ∼455,000), 1-methylpiperidine (97%), and (3-chloropropyl) trimethoxysilane (98%) were provided by Sigma-Aldrich. Li2SiO3 (LSO, 99%, Strem Chemicals), N, N-dimethylformamide (DMF, 99.5%, Beijing Chemical Works), LiNi0.6Co0.2Mn0.2O2 (NCM622, Beijing Dangsheng Material Technology Co., Ltd.), Super P (Imerys Graphite & Carbon), Polyvinylidene difluoride (PVDF, Solvay 5130), N-Methyl-2-pyrrolidone (NMP, 99.0%, Sinopharm Chemical Reagent Co., Ltd.), lithium tablet (Li, Tianjin Zhongneng Co., Ltd.), and polypropylene (PP, Japan Ube) were commercially available and used without further purification.

### Preparation of the GPEs