NMP, CoP-1, CoP-2, CoP-3, and CoP-4 is about 1:0:2, 1:0.5:1.5, 1:1:1, 1:1.5:0.5, 1:2:0, respectively, which is in accordance with our expectation. Therefore, it could be concluded that the stoichiometrical precursors are successfully synthesized via hydrothermal process, and that Co substitution is conducive to the good morphological consistency and distribution of the NMP particles.

The SEM images of all samples are illustrated in **Figures 2k–o** and Figure S1. Particles possessing porous character are observed for all the samples, which could be ascribed to the decomposition of carbonate precursor during heating process. It is also noted that the particles morphology for all the samples are inherited from those of the precursors. For example, particles from the NM sample show irregular morphology and random distribution. However, for the Co-1 and Co-2, the particles exhibit hollow spherical morphology and good distribution. As the increase of Co concentration, particles from the Co-3 exhibit spindlelike morphology and hollow multi-porous architecture. EDS mapping of Co-3 (Figure S1) shows a uniform distribution of Ni, Co, Mn, and O throughout the particle. These results confirm that Co incorporates into the Co-3, and Co substitution changes the morphology of both primary and aggregate particles, which might be benefit for lithium and electron diffusion during lithiation/delithiation reaction.

Schematic diagram for the synthesis process of the porous hollow superlattice NiMn2O4/NiCo2O4 mesocrystals are illustrated in **Figure 2p**. The acetate molecules and hexamethylenetetramine dissolve in the solvent then gather together to form small droplets. During solvothermal treatment, the hexamethylenetetramine decomposes, and creates CO2 bubbles. By capturing the metal ions on bubble surface, an oval-like microsphere precursor is obtained successfully. During calcination process, metal ions move into crystal structure of the superlattice NiMn2O4/NiCo2O4 mesocrystals, meanwhile, CO2 release in the carbonate precursors, leaving pores on the obtained hollow complex oxides (Luo D. et al., 2014; Luo et al., 2015; Ma Z. et al., 2015; Shen et al., 2015).

To understand the oxidation/reduction and phase transformation processes in electrode reactions of the synthesized samples, the CV curves for the initial three cycles of NM and Co-3 samples are presented in **Figures 3A,B**, which are tested at a scan rate of 0.1 mV s−1 in the voltage range of 0.01–3.00 V vs. Li/Li+ at room temperature, respectively. The CV curves of both samples for the first cycle are significantly different from those for the following cycles, because of the irreversible electrochemical reaction during the first discharge cycle. However, no obvious alteration is observed between the second cycle and the third cycle, suggesting the good reversibility of lithium insertion and extraction reactions. In the first cycle, as can be seen in **Figure 3A**, there is a sharp reduction peak at 0.9 V and a broad reduction peak between 0.3 and 0.8 V. The sharp one associates with the initial reduction of Mn3+–Mn2+, and the broad one can be attributed to the reduction of NiMn2O4 to metallic Mn and Ni and the formation of Li2O (Li J. F. et al., 2013b; Kang et al., 2015; Ma Y. et al., 2015; Ma Z. et al., 2015). With the increase of Co concentration (Figure S2), it is clear that the sharp reduction peak at 0.9 V is disappeared, and the reduction peaks which is attributed to the reduction of the oxides to metal and the formation of Li2O are narrower than that of NiMn2O4. In comparison, as shown in **Figure 3B**, a new reduction peak appear at 1.0 V, which can be corresponded to the reduction of Co3+–Co2+, and the broad reduction peaks

of Co-3 are completely separated. Yuan et al. (Ma Z. et al., 2015) and Lee et al. (Kang et al., 2015) reported that "the sharp peak at 0.5 V should correspond to the reduction of NiMn2O4 to metallic Mn and Ni," while Yang et al. (Zhang et al., 2016) and Wang et al. (Leng et al., 2015) reported that "the peak at 0.8 V should correspond to the reduction of NiCo2O4 to metallic Co and Ni," and both peaks "relate to the formation of Li2O and solid-electrolyte interface (SEI)" (Li J. F. et al., 2013c; Zhang et al., 2014, 2018; Gao et al., 2015). These results further confirm that Co-3 is a kind of mixed-phase NiMn2O4/NiCo2O4 composite material, which is in accordance with the XRD results and TEM results as aforesaid. For the Co-3 sample, in the anodic sweep, a wide oxidation peak (∼2.0 V) can be attributed to the oxidation of metals to oxides accompanying the decomposition of Li2O. Based on the above CV analysis, in conjunction with the previously reported storage mechanism for NiMn2O4 and NiCo2O4 (Li J. F. et al., 2013c; Zhang et al., 2014, 2016, 2018; Gao et al., 2015; Kang et al., 2015; Leng et al., 2015; Ma Y. et al., 2015;

Schematic diagram (p) for the synthesis process of the porous hollow superlattice NiMn2O4/NiCo2O4 mesocrystals.

Ma Z. et al., 2015; Su et al., 2015), the electrochemical reactions for the electrodes can be summarized as follows:

$$\text{NiMnO}\_{4} + 8\text{Li}^{+} + 8\text{e}^{-} \rightarrow \text{Ni} + 2\text{Mn} + 4\text{Li}\_{2}\text{O} \quad \text{(1)}$$

$$\text{Ni} + 2\text{Mn} + 3\text{Li}\_{2}\text{O} \leftrightarrow \text{NiO} + 2\text{MnO} + 6\text{Li}^{+} + 6\text{e}^{-} \text{(2)}$$

$$\text{NiCo}\_{2}\text{O}\_{4} + 8\text{Li}^{+} + 8\text{e}^{-} \rightarrow \text{Ni} + 2\text{Co} + 4\text{Li}\_{2}\text{O} \quad \text{(3)}$$

$$\text{Ni} + 2\text{Co} + 3\text{Li}\_{2}\text{O} \leftrightarrow \text{NiO} + 2\text{CoO} + 6\text{Li}^{+} + 6\text{e}^{-} \text{(4)}$$

$$\text{NiMn}\_{2-x}\text{Co}\_{x}\text{O}\_{4} + 8\text{Li}^{+} + 8\text{e}^{-} \rightarrow \text{Ni}$$

$$\qquad + (2-x)\text{Mn} + x\text{Co} + 4\text{Li}\_{2}\text{O} \qquad \text{(5)}$$

$$\text{Ni} + (2-x)\text{Mn} + x\text{Co} + 3\text{Li}\_{2}\text{O} \leftrightarrow \text{NiO}$$

$$\qquad + (2-x)\text{MnO} + x\text{CoO} + 6\text{Li}^{+} + 6\text{e}^{-} \qquad \text{(6)}$$

**Figure 3C** depicts the initial charge-discharge curves of the NM, Co-1, Co-2, Co-3, and Co-4 samples. The electrochemical tests are carried out at a current density of 0.05 A g−1 with the voltage of 0.01–3.00 V vs. Li/Li+ at room temperature. As can be seen in **Figure 3C**, there is a representative plateaus of the profile for

NM around 0.7 V in the first discharge process. By comparison, a representative profile of Co-3 which is composed of two plateaus around 0.9 V and 0.7 V in the first discharge process can be observed clearly, which relates to the insertion of lithium into the crystal structure of NiCo2O4 and NiMn2O4, respectively, accompanied by the formation of Li2O (Wang J. X. et al., 2014; He et al., 2015, 2017; Mo et al., 2015; Li et al., 2016). The initial charge capacity for the NM, Co-1, Co-2, Co-3, and Co-4 is 513.5, 580.8, 636.9, 665.4, and 595.7 mAh g−1 , respectively. It is obvious that the initial charge capacities of the Co-substituted samples are superior than that of the NM, particularly, the Co-3 sample exhibit the optimal initial charge capacities of all samples.

The rate capability of all samples are given in **Figure 3D**. The cells are charged/discharged at a current density of 0.05 A−1.6 A g−1 , and then charged/discharged at a current density of 0.1 A g−1 with the voltage of 0.01–3.00 V vs. Li/Li+ at room temperature. It is noted that the Co-substituted samples show improved rate capability compared to the pristine counterparts at high rates. Specifically, the Co-3 demonstrates the best rate performance and delivers the highest reversible capacities of all samples. Remarkably, it can maintain the charge capacities of 611.4, 591.1, 554.6, 499.1, 467.6, 435.6, and 606.8 mAh g−1 after 5 cycles at the current density of 0.1, 0.2, 0.5, 1, 1.2, 1.6, and 0.1 A g−1 , respectively. Even at a current density as high as 1.6 A g −1 , the Co-3 can still deliver a stable capacity of about 435.6 mAh g−1 after 5 cycles, which is still higher than the theoretical capacity of the graphite (372 mAh g−1 ), indicating the excellent rate capability.

The comparison of the cycling stability between the NM and the Co-substituted samples are presented in **Figure 3E**. The cells were charged/discharged 2 times at a current density of 0.05 A g−1 , and then charged/discharged 100 times at a current density of 1 A g−1 with the voltage of 0.01–3.00 V vs. Li/Li+ at room temperature. The charge capacity retention reaches 61.3, 79.7, 73.8, 90.4, and 70.2% after 100 cycles for NM, Co-1, Co-2, Co-3, and Co-4, respectively. In particular, as can be seen in **Figure 3F**, the Co-3 maintains 532.2 mAh g−1 with 67.8% capacity retention after 500 cycles at a current density of 1 A g−1 with the voltage of 0.01–3.00 V vs. Li/Li+ at room temperature. It means that the Co-3 can still deliver a stable capacity of about 360.8 mAh g−1 after 500 cycles, which is still near the theoretical capacity of the graphite (372 mAh g−1 ), demonstrating the outstanding cycling stability. In conclusion, the porous Co-substituted samples in this work exhibit a superior electrochemical performance compared with the NM, specially, the Co-3 exhibits the optimal electrochemical performance of all samples.

To further understand the effects of the porous architecture on the electrochemical properties of Co-substituted samples, the Co-1 is selected as a reference, which has the similar structure but different morphological architecture, compared to the NM. The SEM images of NM and Co-1 electrodes before and after 100 cycles are shown in **Figure 4**. The cells are charged to 3.00 V at high delithiation state. As can be seen in **Figure 4**, after 100 cycles, both samples (**Figures 4b,d**) show fuzzy surface, which may be due to the formation of Li2O and the electrolyte decomposition (Wang J. X. et al., 2014; Mo et al., 2015; Li et al., 2016; He et al., 2017). Compared to the fresh electrode (**Figure 4a**), the irregular morphological NiMn2O4 particles suffer serious damage and expand into aggregate particles after 100 cycles (**Figure 4b**). The aggregation can be ascribed to the volume expansion on

FIGURE 4 | SEM images of (a,b) NM electrodes, (c,d) Co-1 electrodes, (a,c) before cycling, (b,d) after 100 cycles at a current density of 1A g−1 with the voltage of 0.01–3.00 V vs. Li/Li+.

FIGURE 5 | N2 adsorption and desorption isotherm: (A) NM, (B) Co-1, and (C) Co-3. Nyquist plots of as-prepared electrodes, (D) after the 1st cycle and (E) after the 100th cycle, (F) the relationships between Z're and ω −1/2 after the 1st cycle for the NM, Co-1 and Co-3 electrodes.

account of the irreversible electrochemical process from the high lithiation state (Pan et al., 2014; Zhang et al., 2015). In contrast, for Co-1, as shown in **Figures 4c,d**, there is no obvious alteration in the hollow multi-porous particle morphology after 100 cycles. These results demonstrate that the hollow multiporous architecture can provide enough void spaces to alleviate the architectural change, and offer a stable structure for the intercalation and de-intercalation cycling (Zhang et al., 2015). In addition, Nitrogen adsorption/desorption isotherms are shown in **Figures 5A–C**, the sorption isotherms are of type III isotherms, they have a distinct H3-type hysteresis loop in the range of P/P0 = 0.2–0.99, which does not close until the saturation pressure is reached. Obviously, the BET specific surface area (SSA) of the Co-1 is 6.909 m2 g −1 , which is larger than that of the NM sample (5.199 m2 g −1 ). The larger SSA can be derived from the hollow multi-porous architecture, leading to a shorter Li ions diffusion distance during the charge-discharge processes.

Furthermore, nyquist plots of the as-prepared electrodes and corresponding EIS and lithium ion diffusion coefficient (DLi+) results are exhibited in **Tables 1**, **2** and **Figures 5D–F**. The semicircle at high and intermediate frequency region relates the SEI film impedance (Rsei) and the charge transfer impedance (Rct), respectively. The inclined straight line at the low frequency region is attributed to Warburg impedance (Rw), which is associated with Li+ diffusion into the bulk (Qiu et al., 2011; Li et al., 2012; Chen et al., 2014a,b; Sun et al., 2014; Wang L. Y. et al., 2014; Xu et al., 2014; Choi et al., 2015; Wu F. X. et al., 2016; Wu et al., 2018). The lithium ion diffusion coefficient was calculated via a widely accepted method (Reddy et al., 2013; Li et al., 2015). With the increase of cycle number, it is obvious that the values of R for all samples rise. After 100 cycles, the resistance of all samples increase, and both samples (**Figures 4b,d**) show fuzzy surface. These results indicate that the SEI formed on the electrode. In addition, as shown in **Figures 4c,d**, there is no obvious alteration in the hollow multiporous particle morphology after 100 cycles, which suggests that the uniform hollow multi-porous architecture can alleviate the pore size/function change originating from the volume expansion and the SEI formation during cycling. Specially, the

TABLE 2 | The values of specific surface area and DLi+ for the NM and Co-3 samples.

value of R for NM after 1 cycle and 100 cycles is 218.5 and 719.3 , respectively, the dramatical increase of the resistance could be ascribed to the formation and expansion of aggregate particles (as seen in **Figure 4**), which is certainly harmful to electrolyte penetration and the lithium ions and electron transport. In comparison, the Co-substituted samples do not have significant changes, which should be attributed to that the hollow multi-porous architecture can provide enough void spaces to prevent particle aggregation, resulting in enhanced the charge transport kinetics during discharging/charging processes (Zhang et al., 2015). It is clear from the above discussions that the excellent rate capability and cycle stability of Cosubstituted samples are attributed to the porous architecture by following factors: (1) provides enough void spaces to suppress the architectural collapse during cycle, (2) reduces the lithium ions diffusion and electron-transportation distances; (3) facilitates the electrolyte penetration resulting in enhanced the charge transport kinetics.

To further evaluate the effects of the superlattice structure on electrochemical properties of the NiMn2O4/NiCo2O4, the electrochemical kinetics of the Co-3 and Co-1 samples have been studied. The reason why choose Co-3 and Co-1 is partly due to the Co-3 exhibits superior electrochemical properties and superlattice structure, and partly due to Co-1 owns similar porous morphology with the Co-3, while is made up with a single phase. The BET results reveal that the SSA of the Co-1 and Co-3 samples is 6.909 and 7.050 m2 g −1 , respectively. The similar SSA means the similar Li ions diffusion distance during the charge-discharge process. Therefore, it is supposed that the electrochemical kinetic difference between the Co-1 and Co-3 mainly lies on structure. As shown in **Table 1**, the Co-3 exhibits the lowest initial R and the most stable resistance during cycling among all samples. Furthermore, it can be seen from **Table 2**, the DLi+ of the Co-3 sample is 1.19 × 10−11 cm2 S −1 , which is much higher than that of the Co-1 sample (3.49 × 10−12 cm2 S −1 ). Since the morphology and the SSA of the Co-1 is similar to those of Co-3, the superior electron and Li ion transport diffusive capacity of the Co-3 should be related to the superlattice NiMn2O4/NiCo2O4 structure. A similar phenomenon has also been reported by Chen et al (Suo et al., 2013). They pointed out that the composite phases can supply high ionic conductivity. These results demonstrate that the superlattice NiMn2O4/NiCo2O4 structure can facilitate the ionic conductivity to enhance charge transport kinetics of the bulk material.

The schematic illustrations of the porous hollow superlattice NiMn2O4/NiCo2O4 compound and the mechanism of electrochemical process are exhibited in **Scheme 1**. Many reports (Kang et al., 2015; Ma Y. et al., 2015; Ma Z. et al., 2015; Zhang et al., 2015, 2016) indicate that the aggregation from the irreversible volume expansion and the low charge transport kinetics, including the poor intrinsic electronic conductivity and the slow ion-diffusion rates, is responsible for the fading capacity of NiMn2O4. In this work, the porous hollow architecture can provide enough void spaces to suppress architectural collapse originating from the volume expansion during cycling, and offer a short stable path for the lithium

ions diffusion and electron-transportation. Moreover, the superlattice NiMn2O4/NiCo2O4 structure can facilitate the electronic and ionic conductivity to enhance charge transport kinetics of the bulk material. The synergistic effects of the superlattice NiMn2O4/NiCo2O4 structure and the uniform hollow multi-porous architecture are guaranteed for the excellent electrochemical performance.

### CONCLUSIONS

In summary, the Co substituted NiMn2O4 mesocrystals with controlled phase structure and morphology are successfully synthesized via a simple solvothermal method. The optimal sample with well-distributed hollow multi-porous superlattice NiMn2O4/NiCo2O4 mesocrystals exhibits excellent rate ability and enhanced cycling performance, which maintains 532.2 mAh g−1 with 90.4% capacity retention after 100 cycles, and can still deliver a stable capacity of about 360.8 mAh g −1 after 500 cycles, at a current density of 1 A g−1 . The excellent electrochemical performance can be attributed to the synergistic effect of the superlattice NiMn2O4/NiCo2O4 structure and the uniform hollow multi-porous architecture. The superlattice structure can enhance the Li ion diffusion coefficient to enhance the charge transport kinetics. While, the uniform hollow multi-porous architecture can alleviate the architectural change originating from the volume expansion during cycling and facilitate electrolyte penetration. These combined effects

### REFERENCES

Chen, Y. J., Zhu, J., Qu, B. H., Lu, B., and Xu, Z. (2014a). Graphene improving lithium-ion battery performance by construction of NiCo2O4/graphene hybrid nanosheet arrays. Nano Energy 3, 88–94. doi: 10.1016/j.nanoen.2013.10.008

make the NiCo1.5Mn0.5O4 a promising anode material for the practical application in EV/HEV and energy storage systems.

### AUTHOR CONTRIBUTIONS

LL and QY conceived the idea; QY and LL prepared all materials; QY, HY, and JL conducted SEM experiments; KY, BL, and ZL conducted XRD experiments; LL, QY, and BZ analyzed the data; QY wrote the manuscript and BZ, ZC, and JD commented on it; LL supervised the implementation of the project.

### ACKNOWLEDGMENTS

This work was financially supported by the National Natural Science Foundation of China (No. 51774051 and No. 51404156), the Changsha City Fund for Distinguished and Innovative Young Scholars (No. KQ1707014), Scientific Research Fund of Hunan Provincial Education Department (No. 17B002), the Key Laboratory of Efficient and Clean Energy Utilization, The Education Department of Hunan Province (No. 2016 NGQ005).

### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2018.00153/full#supplementary-material

Chen, Y. J., Zhuo, M., Deng, J. W., Xu, Z., Li, Q. H., and Wang T. H., (2014b). Reduced graphene oxide networks as an effective buffer matrix to improve the electrode performance of porous NiCo2O4 nanoplates for lithium-ion batteries. J. Mater. Chem. A 2, 4449–4456. doi: 10.1039/c3ta1 4624c

**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.

The reviewer, BQ, and handling Editor declared their shared affiliation.

Copyright © 2018 Li, Yao, Liu, Ye, Liu, Liu, Yang, Chen, Duan and Zhang. 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 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.

# SiC Nanofibers as Long-Life Lithium-Ion Battery Anode Materials

Xuejiao Sun1†, Changzhen Shao2†, Feng Zhang2 , Yi Li 2 \*, Qi-Hui Wu1 \* and Yonggang Yang2

*1 Department of Materials Chemistry, School of Chemical Engineering and Materials Science, Quanzhou Normal University, Quanzhou, China, 2 Jiangsu Key Laboratory of Advanced Functional Polymer Designand Application, Department of Polymer Science and Engineering, College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, China*

The development of high energy lithium-ion batteries (LIBs) has spurred the designing and production of novel anode materials to substitute currently commercial using graphitic materials. Herein, twisted SiC nanofibers toward LIBs anode materials, containing 92.5 wt% cubic β-SiC and 7.5 wt% amorphous C, were successfully synthesized from resin-silica composites. The electrochemical measurements showed that the SiC-based electrode delivered a stable reversible capacity of 254.5 mAh g−1 after 250 cycles at a current density of 0.1 A g−1 . It is interesting that a high discharge capacity of 540.1 mAh g −1 was achieved after 500 cycles at an even higher current density of 0.3 A g−1 , which is higher than the theoretical capacity of graphite. The results imply that SiC nanomaterials are potential anode candidate for LIBs with high stability due to their high structure stability as supported with the transmission electron microscopy images.

Keywords: nanofibers, silicon carbide, anode, lithium-ion batteries, long-life

### INTRODUCTION

In recent years, due to the wide applications of lithium ion batteries (LIBs) in the portable electronic devices and especially in electric vehicles (EVs), higher criterion on their energy capacity and density as well as cycleability would be established. To further enhance the electrochemical performances of LIBs, it is necessary to find alternative electrode materials to replace the currently commercial ones e.g. graphitic carbons (An et al., 2017a,b; Zhang et al., 2018). SiC, a high bandgap semiconductor, is generally thought to be electrochemically inactive to Li, and often acts as backbone or buffer matrix to enhance the strength of electronic composite materials (Timmons et al., 2007; Jeon and Lee, 2014; Wang C. et al., 2015; Wang W. et al., 2015). But it is interesting that some researchers have studied the cycling performance of SiC-based nanocomposites as anode materials for LIBs. For example, nanoscopic carbon-coated SiC has been synthesized using a simple CVD method (Sri Devi Kumari et al., 2013), which delivered a reversible lithium storage capacity of about 1,200 mAh g−1 after 200 cycles, showing excellent cyclability. Zhang and Xu (2014) prepared a nanocrystalline SiC thin film electrode grown on the stainless steel, which showed a stable discharge capacity of 309 mAh g−1 over 60 cycles. Hu et al. (2016) produced SiC nanowires on the graphite paper, which delivered 397 mAh g−1 after 100 cycles. These results indicated that when a non-conducting SiC material reduces to nano size, it can also serve as anode materials for LIBs. As explained by Li et al. (2016), the capacity of SiC nanomaterials could be activated upon an intimate contact of SiC with graphite, which facilitates the electronic and ionic transport while suppressing the oxidation of SiC. Moreover, during the electrochemical studies of the SiC nanomaterials, the gradual enhancement of capacities was found, which has not been well studied.

Edited by:

*Qiaobao Zhang, Xiamen University, China*

Reviewed by:

*Shaohua Guo, Nanjing University, China Biao Gao, Wuhan University of Science and Technology, China Huiqiao Li, Huazhong University of Science and Technology, China*

\*Correspondence:

*Yi Li<EMAIL_ADDRESS>Qi-Hui Wu<EMAIL_ADDRESS>
*†These authors have contributed equally to this work.*

#### Specialty section:

*This article was submitted to Physical Chemistry and Chemical Physics, a section of the journal Frontiers in Chemistry*

> Received: *12 March 2018* Accepted: *23 April 2018* Published: *14 May 2018*

#### Citation:

*Sun X, Shao C, Zhang F, Li Y, Wu Q-H and Yang Y (2018) SiC Nanofibers as Long-Life Lithium-Ion Battery Anode Materials. Front. Chem. 6:166. doi: 10.3389/fchem.2018.00166* Sun et al. SiC Anode Materials

Up to date, the reports on SiC as anode material for LIBs are few. As a potential material for long-life LIBs, the lithium ions diffusion/reaction mechanism in/with SiC nanostructures needs to be further investigated, which would benefit to later similar studies.

SiC nanostructures can be synthesized by converting well defined SiO2 nanoarchitectures through carbothermal reduction at a high temperature over 1,400◦C, by which the SiC products could preserve the structure regularity of the starting SiO2 nanomaterials (Miller et al., 1979; Martin et al., 1998). Nanosized SiC materials can also be fabricated by pyrolysis of Sicontaining polymers such as polybissilsesquioxanes (Wang C. et al., 2015; Zhang et al., 2015). In our previous work, we reported SiC/C composite nanotubes (69 wt% SiC and 31 wt% C) prepared via pyrolysis of resorcinol-formaldehyde (RF) resinsilica composite nanotubes at 1,400◦C under argon atmosphere. The RF resin-silica composite was synthesized using a sol-gel transcription method taking a chiral gelator as template (Shao et al., 2017b). Herein, SiC nanofibers (with 7.5 wt% residual C) were fabricated using RF resin-silica composite nanofibers as the starting material and then applied as electrode materials for LIBs. The electrochemical data showed that the SiC-based electrode possessed superior lithium ion storage capability and cycling stability. Further, the electrochemical reaction mechanism of lithium and SiC nanofibers was proposed.

### EXPERIMENTAL

### General Methods

Field emission scanning electron microscopy (FE-SEM, 4800 instrument) and transmission electron microscopy (TEM, FEI TecnaiG220) images were taken at 3.0 and 200 kV to observe the samples' morphologies and nanostructures, respectively. A Jobin Yvon Horiba HR 800 LabRAM confocal microprobe Raman system was applied to collect the Raman spectrum with Ar laser excitation at 514.5 nm and a power of 10 mW. Wide angle Xray diffraction (WAXRD) patterns were obtained based on an X'Pert-Pro MPD X-ray diffractometer with Cu Kα radiation (1.542Å). Specific surface area and pore-size distribution were determined according to N2 adsorption isotherm measured from a Micromeritics Tristar II 3020 instrument via the Brunauer-Emmett-Teller (BET) and Barrett-Joyner-Halon (BJH) methods.

### Synthetic Procedure of RF Resin-Silica Composite and SiC Nanofibers

Gelator D-1, whose molecular structure is shown in Figure S1, was synthesized according to the literature (Li et al., 2013; Zhang and Xu, 2014). A typical synthesis route of the resin-silica composite was as following: D-1 (200 mg, 0.31 mmol) and resorcinol (180 mg, 1.63 mmol) were dissolved in deionized water (35 mL) and ethanol (5.0 mL) mixed solution at 40◦C under vigorous stirring. After that, concentrated ammonia aqueous solution (0.6 mL), formaldehyde (0.23 mL), and tetraethylorthosilicate (TEOS, 0.8 mL, 3.6 mmol) were added into the solution in sequence. The reaction mixture was then stirred at 40◦C for 10 h followed by heating to 80◦C and standing in static for 24 h. After filtration and washing with water, the as-prepared sample was brown-reddish powders and denoted as **S1**. Thereafter, sample **S1** was annealed at 350◦C for 2.0 h and then at 1,400◦C for 4.0 h with a heating rate of 2.0◦C· min−1 in Ar atmosphere to gain the as-synthesized SiC product, denoted as **S2**. After the sample was cooled down to room temperature naturally, it was immersed into10 wt% HF aqueous solution for 24 h at room temperature to completely remove the resident SiO2, and finally washed with deionized water for three times.

### The Electrochemical Measurements

Electrochemical tests were performed using CR2016 coin-type cells assembled in an Ar-filled glove box. For preparation of the working electrode, **S2** (80 mg), acetylene black (AB, 10 mg) and binder polyvinylidene fluoride (10 mg, dissolved in Nmethylpyrrolidone) were mixed together. The resultant slurry was then uniformly coated on a Cu foil current collector and dried overnight under vacuum. Electrochemical cells were assembled using **S2** electrode as the cathode, metallic lithium foil as the anode, Celgard 2325 porous film as the separator, and 1.0 M LiPF6 solution (dissolved in a mixed solvent of ethylene carbonate and diethyl carbonate, 1:1 by volume) as the electrolyte. The cells were charged and discharged between 3.0 and 0.01 V at room temperature.

### RESULTS AND DISCUSSION

In this work, the RF resin-silica composite **(S1)** was used as the source to fabricate aimed product-SiC as mentioned in the experimental section. The template and RF resin in **S1** was firstly converted to C at lower temperature (∼600◦C) when heated in Ar. When the temperature rose up to 1,400◦C, the silica in **S1** reacts with the C to produce SiC. The obtained sample **S2** was firstly characterized using WAXRD and Raman spectroscopy. As shown in **Figure 1**, in the WAXRD pattern (**Figure 1A**), peaks displayed at 2θ = 35.6, 41.4, 60.1, and 71.9◦ are indexed to (111), (200), (220), and (311) planes of cubic β-SiC phase (JCPDS: 29- 1129). At the same time, a broad peak centered at 2θ = 22.9◦ is ascribed to the (200) plane of graphite. There are not any diffraction lines that could be assigned to SiO2, which suggests that the SiO2 compound has been safely removed. In the Raman spectrum (**Figure 1B**), the D and G bands appear at 1,331 and 1,581 cm−1 , respectively. The value of IG/ID is 0.95, indicating that the carbon in sample **S2** is predominantly amorphous. Due to the relatively low intensities of SiC Raman shifts compared to those of C, it is difficult to define the SiC Raman peaks in **Figure 1B**. TGA analysis, shown in **Figure 2**, discloses that ∼7.5 wt% of free carbon presents in the sample **S2**. These results show that **S1** has been successfully converted to crystalline SiC with small amount of residual carbon.

FE-SEM and TEM images of the samples **S1** and **S2** are shown in Figure S2 and **Figure 3**. Sample **S1** are right-handed helical nanofibers with width, wall thickness and helical pitch of 80–150, 15–30, and 800–1,200 nm, respectively. The lengths of the nanofibers are several microns as observed in Figure S2. After pyrolysis, the sample **S2** exhibits twisted fiber-like morphology, as shown in **Figures 3a,b**. The average length and diameter of the nanofibers are several microns and 30–50 nm.

In the HRTEM image (**Figure 3c**), the interplanar distance is 0.25 nm, which is consistent with the lattice fringes of cubic β-SiC (111) face. The selected area electron diffraction (SAED) pattern (**Figure 3d**) implies that the obtained SiC is highly crystalline. These observations are consistent with the results of WAXRD analysis in **Figure 1**. The N2 adsorptiondesorption isotherms and the BJH pore size distribution plots calculated from the adsorption branch for the sample **S2** are shown in **Figure 4**. It shows type-IV isotherms with H3 hysteresis loop, which are the characteristic of microporous and mesoporous materials. These micro- and mesopores are produced mainly due to the breakup of C-C bonds and formation of crystalline SiC during the pyrolysis process as well as the dissolution of SiO2 compound during the HF solution treatment. The BET specific surface area and BJH adsorption pore volume are 326.4 m2 ·g −1 and 0.86 cm3 ·g −1 , respectively.

Cycling performance of the **S2** electrode was evaluated, as shown in **Figure 5A**. The first discharge and charge capacities are 309.3 and 221.9 mAh g−1 , respectively, giving an initial Coulombic efficiency of 71.7%, however, which quickly increases and keeps at above 97% in the following cycles. The discharge capacity quickly decreases to 195.4 mAh g−1 in the 8th cycle, thereafter stabilizes at about 170 mAh g−1 and rises slowly to 205.4 mAh g−1 in the 250th cycle. Apparently, the relatively low capacity of the **S2** electrode is primarily attributed to its low carbon content and the electrochemically inertial nature of SiC. Moreover, the cycling performance of the **S2** electrode tested at a higher current density of 0.3 A g−1 is also shown in **Figure 5B**. In the initial 40 cycles, the discharge capacity decreases from 247.9 to 101.8 mAh g−1 . After that, it starts to increase gradually, and reversible capacities of 254.5 mAh g−1 at the 250th cycle and 540.1 mAh g−1 at the 500th cycle are obtained. It seems that high current density will accelerate the activation of SiC electrode. Clearly, the SiC electrode exhibits higher capacity than the commercial graphite electrode (∼370 mAh g−1 ) and also better cycling stability than the common Si electrode (Zhang et al., 2014).

It's interesting to observe that the capacity increases after long cycles at the current density of 0.3 A g−1 , as displayed in **Figure 5B**. We firstly reported similar capacity ascending curve for the SiC/C composite (containing 17 wt% carbon, at the current density of 0.1 A g−1 ), which was attributed to a gradually activation of electrode in the lithiation/delithiation process similar to some metal oxides (Cheng et al., 2013; Wang C. et al., 2015; Sun et al., 2016; Li J. et al., 2017; Shao et al., 2017a). Zhang and Xu also reported the capacity increasing phenomenon for the nanocrystalline SiC from the second cycle. They proposed the reduction of Si4+ to Si0 and the formation of Li-C and Li-Si alloys in the electrode (Zhang and Xu, 2014). This is more rather correct explanation, because a high current density will

from 0.1 A g−1

to 5 A g−1

; (D) Nyquist plot for S2 electrode in the frequency range from 100 KHz to 10 MHz before and after cycling.

fasten the reaction of Li with SiC nanomaterials to form Li-Si and Li-C alloys. Consequently, after tens times of cycling, the SiC electrode will slowly convert to Si/C nanostructure composite, because the re-formation of SiC from Si/C composite needs at rather high temperature. In this case, previously electrochemical inert SiC electrode gradually transfer to electrochemically active Si/C composite electrode, which will interpret why the electrode capacity keeps growing at 0.3 A g−1 current density. In order to prove our hypothesis, the **S2** electrode was cycled at 5 A g−1 as shown in Figure S3. Its capacity decreases to about 76 mAh g −1 at the 16th cycle, and then gradually increases to 128 mAh g −1 at the 40th cycle. It is found that the capacity enhancement is much earlier and faster and the electrodes cycled at 0.1 and 0.3 A g−1 . It is very interesting to find that after SiC partially transfer to Si/C composite, it still maintains original nanofiber structure. This may explain why the SiC electrode could possess a rather good cycling stability. FE-SEM images of the **S2** electrode after cycling test were then taken and showed in **Figure 6**. The nanofibers keep original morphology, and no changes in lengths or diameters are observed. Figure S4 shows the XRD patent of the electrode after cycling at 5 A g−1 for 40 times. It could be seen that the bump for C species slightly increases and the diffractions for the SiC decreases, which indicates the change of component in the SiC/C composite. Out of our expectation, the signals for Si were not observed in the patent, which may because that the Si decomposed from SiC is not crystalline but amorphous.

Unlike Si-based electrodes, which suffer from huge volume changes during the lithiation/delithiation process, resulting in the pulverization of Si nanostructure and consequently shortening the cycling properties of the batteries. SiC nanostructures have

at 0.1 A g−1 ; (b) 500 cycles at 0.3 A g−1 .

rather high hardness, the partial decomposition of SiC into to Si and C composites would not alter the base structure of SiC, meanwhile, the decomposed Si nanoclusters could well distribute into the C matrix. This may explain why after longtime cycling the Si/C composite conversed from SiC/C still

\**1C* = *0.37 A g*−1 *;* \*\**the data wasn't marked in Timmons et al. (2007).*

preserves nanofiber structure. The Raman spectra of **S2** electrode after cycling have been measured and are showed in Figure S5. The values of IG/ID are 0.86 at the current density of 0.1 A g−1 after 250 cycles and 0.94 at the current density of 0.3 A g−1 after 500 cycles, respectively. Compared with the sample before cycling (the IG/ID value is 0.95), the graphitization change is small. It is reasonable to suggest that the decomposed C from SiC is also amorphous, because under the electrochemical reactions it is not easy to form crystalline carbon nanostructure during lithiation/delithiation.

Compared with our previously reported results, the C content in the SiC/C composites exhibited great effect on the capacity, as listed in **Table 1**. At low current density (0.1 A g−1 ), pure carbon shows the biggest capacity value, and the small amount of SiC will lower the capacity because of its electrochemical inertness. When SiC is the major moiety, the capacity decreases with the content of C after long cycles. However, at higher current density (over 0.3 A g−1 ), the SiC/C composites show higher capacity than pure carbon after long cycles, because high current density will accelerate the activation of SiC electrode. The SiC transforms into Si/C composite keeping original SiC nanostructure (Han et al., 2011). The theoretical capacity of Si is much higher than that of C (Ren et al., 2013, 2014; Sun et al., 2015; Li H. et al., 2017), therefore, SiC/C composites would show both higher capacity than C materials and superior stability than Si materials. The relative lower capacity values obtained in current work is due to the lower original C content, which may slowly increase with the cycling number.

The rating capability of the **S2** electrode was evaluated at various current densities from 0.1 to 5 A g−1 , as shown in **Figure 5C**. Reversible capacities of 165.1, 111.7, 76.2, 62.9, 51.5, 110.1, and 160.1 mAh g−1 are obtained, respectively, indicating excellent rating capability. To study the conductivity of the **S2** electrode, the electrochemical impedance spectra (EIS) were also collected and shown in **Figure 5D**, which illustrates an impedance of 180 before cycling, 120 after 100 cycles, and 50 after 500 cycles. The rapidly decreased impedance implies that the SiC electrode after cycling also possesses a good electrical conductivity and a rapid charge-transfer reaction for Li ion insertion and extraction. This result is consistent with our above proposal that SiC would gradually decompose into Si and C during electrochemical Li insertion/extraction.

### CONCLUSIONS

Novel SiC nanofibers derived from resorcinol-formaldehyde resin/silica composites were designed and synthesized successfully. When used as anode materials for LIBs, they exhibited superior cycling stability, good rating capability and low impedance. A special capacity increasing phenomenon was observed after long cycles, which was ascribed to the partial decomposition of SiC nanostructure into Si/C composites during the lithiation/delithiation process. In addition, the effect of C content in the SiC/C composite was compared and discussed. The results in this work disclose that the nano-sized SiC materials are promising anode candidate for long-life LIBs due to their high nanostructure stability.

### AUTHOR CONTRIBUTIONS

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

### ACKNOWLEDGMENTS

This work was supported by the National Natural Science Foundation of China (Nos. 21574095, 51473106), the Priority Academic Program Development of Jiangsu High Education Institutions (PAPD), the Natural Science Foundation of Fujian Province, China (Grant No. 2016J01069).

### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2018.00166/full#supplementary-material

### 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 © 2018 Sun, Shao, Zhang, Li, Wu and Yang. 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 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.

# N/S Co-doped Carbon Derived From Cotton as High Performance Anode Materials for Lithium Ion Batteries

Jiawen Xiong1 , Qichang Pan1 , Fenghua Zheng1 , Xunhui Xiong1 , Chenghao Yang1,2 \*, Dongli Hu3 and Chunlai Huang3

*1 Guangzhou Key Laboratory for Surface Chemistry of Energy Materials, New Energy Research Institute, School of Environment and Energy, South China University of Technology, Guangzhou, China, 2 Guangdong Engineering and Technology Research Center for Surface Chemistry of Energy Materials, New Energy Research Institute, School of Environment and Energy, South China University of Technology, Guangzhou, China, 3 Jiangsu Key Lab of Silicon Based Electronic Materials, Jiangsu GCL Silicon Material Technology Development Co., Ltd, Xuzhou, China*

### Edited by:

*Qiaobao Zhang, Xiamen University, China*

### Reviewed by:

*Jinhu Yang, Tongji University, China Yunxiao Wang, Institute for Superconducting and Electronic Materials, Australia Hui Xia, Nanjing University of Science and Technology, China*

\*Correspondence:

*Chenghao Yang<EMAIL_ADDRESS>
### Specialty section:

*This article was submitted to Physical Chemistry and Chemical Physics, a section of the journal Frontiers in Chemistry*

> Received: *30 January 2018* Accepted: *08 March 2018* Published: *26 April 2018*

#### Citation:

*Xiong J, Pan Q, Zheng F, Xiong X, Yang C, Hu D and Huang C (2018) N/S Co-doped Carbon Derived From Cotton as High Performance Anode Materials for Lithium Ion Batteries. Front. Chem. 6:78. doi: 10.3389/fchem.2018.00078* Highly porous carbon with large surface areas is prepared using cotton as carbon sources which derived from discard cotton balls. Subsequently, the sulfur-nitrogen co-doped carbon was obtained by heat treatment the carbon in presence of thiourea and evaluated as Lithium-ion batteries anode. Benefiting from the S, N co-doping, the obtained S, N co-doped carbon exhibits excellent electrochemical performance. As a result, the as-prepared S, N co-doped carbon can deliver a high reversible capacity of 1,101.1 mA h g−1 after 150 cycles at 0.2 A g−1 , and a high capacity of 531.2 mA h g −1 can be observed even after 5,000 cycles at 10.0 A g−1 . Moreover, excellently rate capability also can be observed, a high capacity of 689 mA h g−1 can be obtained at 5.0 A g−1 . This superior lithium storage performance of S, N co-doped carbon make it as a promising low-cost and sustainable anode for high performance lithium ion batteries.

Keywords: lithium-ion batteries, anode materials, sustainable, cotton, N/S co-doped carbon

### INTRODUCTION

In the past decade, lithium-ion batteries (LIBs) have been widely used as power sources for computing, communications, consumer batteries (3C battery), and electric vehicle. And LIBs with a wide range of applications which due to its have high working voltage, high energy density, and long service life (Goodenough and Kim, 2010; Goodenough and Park, 2013; Zheng et al., 2015). Currently, graphite is mostly used as anode material for commercial LIBs due to its good electronic conductivity, low cost and outstanding cycling stability. However, graphite can not meet the increased energy and power density of high performance LIBs due to the low specific capacity and poor rate performance (Huang et al., 2016; Wang et al., 2017; Pan et al., 2018). On the other hand, the low lithiation/de-lithiation potentials (< 0.3 V vs. Li+/Li) which resulting in seriously security issue (Guo et al., 2011). Therefore, it is necessary to develop novel anode materials to replace graphite anode for high-performance LIBs (Pan et al., 2017a).

Recently, amorphous carbon and hard carbon are attracted attentions as promising anode to replace graphite anode for LIBs, which due to these carbonaceous materials exhibit higher specific capacity and offer higher lithiation/de-lithiation potential (Wang et al., 2009; Casas and Li, 2012; Tang et al., 2012). Moreover, these carbonaceous materials with partially graphitic carbons which can accommodate Li+ in the disordered interlayers as well as in the micropores, and can exhibit excellent cycling stability and rate capability (Wu et al., 2003). On the other

**419**

hand, other carbonaceous materials such as carbon nanotube, graphene and fullerenes were also developed as anode for LIBs, and exhibit excellent electrochemical performance (Etacheri et al., 2015; Wang et al., 2015). However, in order to synthesize these carbonaceous materials which rely on hydrocarbon precursors, resulting in expensive cost and commercially nonviable. Therefore, it is necessary to explore a scalable and inexpensive precursor as carbon sources for these carbonaceous materials as anode materials for LIBs.

Nowadays, multitudinous biomass raw materials have been extensively used as precursor for carbonaceous materials and application in LIBs.  However, cotton attracted more attention and have been considered as the most promising compared with the other biomass materials due to its abundant and low cost. On the other hand, in China, cotton is planted around 550 Million tons per year and giant cotton-products are abandoned which from clothes, medical alcohol cotton, and so on. Moreover, lots of abandoned cotton-products will bring lots of problems, such as environmental pollution, safe question, and so on.

Herein, we addressed the above mentioned issues by prepared high performance carbon-based anode materials using cotton as precursor, as shown in **Figure 1**. Highly porous carbon with large surface areas were prepared from cotton via s sample method. And the N/S-coped carbon were further obtained using thiourea as nitrogen and sulfur sources. When evaluated as anode materials for LIBs, these carbon materials exhibit outstanding rate capability and long-term cycling stability.

### EXPERIMENTAL

### Material Preparation

1.5 g cotton was dipped in homogeneous Mg(NO3)2 solution (8 mol L−1 , 20 ml), then dried in the oven. After that, the obtained cotton were annealed at 800◦C under N2 atmosphere for 3 h with a heating rate of 5◦C min−1 . After cooling, the obtained cotton carbon (denoted as CC) were washed with 1 M HCl and distilled water several times, respectively, then dried in the oven.

To obtained the N, S co-doped carbon, the obtained CC were immersed in 100 ml thiourea solution with ratio of 10:1. After drying, the obtained powders were calcined at 800◦C under N2 atmosphere for 3 h. The N, S co-doped cotton carbon (denoted as NS-CC) powders were obtained after cooling.

### Material Characterizations

The XRD patterns of all samples were conducted on the Bruker D8 Advance (Germany) (Cu, Kα, λ = 1.5405 Å). Raman spectra were obtained on a JOBIN-Yvon HR800 Raman spectrometer. Morphology of the all samples were studied by SEM (FEI Quanta 200 FEG) and TEM (Tecnai G2 F20 S-TWIN, Japan). BET method and non-linear density functional theory (NLDFT) (ASAP 2020 Micromeritics) were applied to test the specific surface area and the pore size distribution.

### Electrochemical Measurements

The samples were executed in CR2025 coin cells. The mingled ratio of sample CC, sample NS-CC and carbon black and poly (vinylidenedifluoride) were 7:2:1 (the loaded active electrode materials are about 0.5 mg cm−2 ). The mixture was coating on copper foil to prepare electrodes. Metal lithium boil as counter electrode and 1 M LiPF6 dissolved in ethylene carbonate (EC) and dimethyl carbonate (DMC) (1:1, v/v) as electrolyte. CV and EIS were conducted on a CHI660A electrochemical workstation. Galvanostatic charge/discharge and cycling performance were executed at 25◦C based on the active electrode material corresponding specific capacity.

### RESULTS AND DISCUSSIONS

The morphology of the all samples was characterized by SEM firstly. **Figures 2A,C** exhibit the SEM images of cotton carbon, which shows micron size bulk materials and composed of nanosheets. **Figures 2B–D** shows the SEM images of the cotton carbon after N and S co-doped, which exhibits similar morphology to cotton carbon. Moreover, cotton carbon and S, N co-doped cotton carbon with a large number of pores in the nanosheets according to the HRSEM (**Figures 2C,D**). On the other hand, EDS element mapping of S, N co-doped cotton carbon were investigated, which indicated that C, N, O, and S elements exist S, N co-doped cotton carbon. Furthermore, S and N elements evenly distributed (**Figures 2G–J**) in the carbon matrix. Therefore, the results indicated that the N and S elements were successfully doped into cotton carbon after heat treatment the carbon in presence of thiourea. The microstructure of CC and NS-CC was further studied by TEM, and the results are shown in **Figures 2E,F**. It clearly illustrates that the samples are amorphous carbon with nano/meso porous structure (Chen et al., 2014). Furthermore, typical selected area electron diffraction further proved that the carbon were amorphous, which corresponding to TEM results (Zhu and Akiyama, 2016).

The XRD patterns of the CC and NSCC are showed in **Figure 3A.** There are no obvious peaks for magnesium compounds in CC sample, which indicated that the Mg(NO3)2 or MgO are removed completely by washed with dilute hydrochloric acid. And two broad peaks at around 23 and 43◦ can be observed both at CC and NS-CC samples, which can be attributed to (002) and (100) graphitic planes, respectively (Hou et al., 2015; Chen et al., 2017). The Raman spectra of the CC and NS-CC are shown in **Figure 3B**, two peaks at around 1,361 and 1,596 cm−1 were obtained, which corresponding to the D band and G band for carbon materials, respectively (Li et al., 2015; Gao et al., 2017). Furthermore, the D band arises from edges, defects, and disordered carbon, whereas the G band is ascribed to sp2 hybridized carbon (Pan et al., 2017b). Therefore, a high ID/IG band intensity ratio indicates the generation of large amounts of defects. The ID/IG ratio for NS-CC is higher than that of CC, which indicated that more vacancies and defects generated by doping N and S atoms into the carbon material. More importantly, more vacancies and defects are beneficial for the transmission of Li-ion and offer more active site for Li storage,

which resulting in improved electrochemical performance (Qie et al., 2015; Lu et al., 2017).

The XPS measurement was conducted to confirmed that the presence of N and S elements in NS-CC. As shown in Figures S1A,B, the survey spectrum for NS-CC and CC exhibit two predominant peak at around 284.8 and 532 eV can be observed, which can be assigned to C and O. It certified for the NS-CC that N and S atoms were successfully doped corresponding with EDS element mapping. In Figures S1C,D, the C 1s XPS spectrum for both CC and NS-CC can be deconvoluted into five peaks, which corresponding to C = C, C-C/C = N, C-O, C = O and π-π ∗ (Liu et al., 2017), respectively. The high resolution N 1s spectrum as shown in **Figure 3E**, the N 1s speak can be fitted by three component peaks at around 401.9, 400.5, and 398.3 eV, which can be ascribed to graphitic N, pyrrolic N and pyridinic N, respectively (Ou et al., 2015). As for high resolution of S 2p spectrum (**Figure 3F**), there were five peaks attributed to -C-S-C- bond and -C-SOX-C- bond. Therefore, these results indicated that the N and S has been successfully incorporated into the carbon structure of NS-CC. And the content of N and S in NS-CC were confirmed for 3.0 and 1.4%, respectively. The heteroatoms N and S can effectively enlarge the interlayer space because of their lager radius than C atom, resulting in forming the defects and providing more active sites for Li-ions on the carbon materials (Xu et al., 2015, 2016; Xiong et al., 2016). Moreover, pyridinic N and quaternary N are favorable for Li+ and electrons and the doped S in the carbon materials can participate in the redox reactions contribute to the reversible capacity (Ma et al., 2018).

The nitrogen adsorption/desorption isotherms and the pore size distribution of CC and NS-CC are shown in **Figures 3C,D.** As shown in **Figure 3C** both of the two samples showed type IV isotherms (Islam et al., 2017). The BET specific surface area of CC and NS-CC are 1235.35 and 1326.20 m2 g −1 , respectively. The BET specific surface areas of NS-CC increased compare than the CC after N and S atoms doping. In addition, the highly porous structure of the two samples were further evaluated by Barrett-Joyner-Halenda (BJH) calculations (**Figure 3D**). It can be seen that the two samples exhibit a broad pore size distribution, and the pore size of the two samples were centered at around 2 and 5 nm, respectively. Therefore, NS-CC exhibits larger specific surface area which can provide more active sites for lithium ion storage. Moreover, the highly porous structure of the two samples can greatly shorten the diffusion distance of both electrons and ions, which resulting in improved rate performance (Hao et al., 2014).

The electrochemical performance of CC and NS-CC was first measured by cyclic voltammetry (CV) with voltage range of 3.0– 0.01 V. **Figure 4A** exhibits the CV curves of the NS-CC sample. During the first discharge cycle, three cathodic peaks at around 1.4 and 0.65 V can be obtained and disappeared in the subsequent cycles, which corresponding to formation of a solid-electrolyte interphase (SEI) film (Wang et al., 2011; Jiang et al., 2013) as well as some irreversible and side reactions associated with the decomposition of electrolyte (Yoshio et al., 2000). Moreover, the CV profiles almost overlapped after the initial scanning cycle, which indicates that the structural stability of the NS-CC electrode during the subsequently cycling. On the other hand, the CV curves of CC electrode (Figure S2A) are similar to the NS-CC electrode.

The charge and discharge profiles of CC and NS-CC electrode at 0.1 A g−1 with cutoff voltage window of 0.01–3.0 V were

studied. As shown in **Figure 4B** and Figure S2B, two plateaus at around 1.4 and 0.65 V can be observed during the first discharge process, which corresponding to the forming of SEI layer and decomposition of the electrolyte, in good agreement with the above CV results. The first discharge capacity of CC and NS-CC are as high as 3,275.1 and 4,179.1 mA h g−1 , while the initial reversible capacity are only 1,512.3 and 1,823.3 mA h g −1 , corresponding to low initial Coulombic efficiency of 46.2 and 43.6%, respectively. The large irreversible capacity loss for the CC and NS-CC electrode can be ascribed to formation of a SEI layer on the relatively large specific surface area (Jiang et al., 2013). Furthermore, a mass of reduction of oxygen functionalities on the carbon materials surface (Bhattacharjya et al., 2014) and reduction of electrolyte components on the active electrode of the CC and NS-CC electrode (Hu et al., 2007) also contributed to the irreversible capacity loss.

The cycling performance of the CC and NS-CC electrode were evaluated at 0.2 A g−1 (**Figure 4D**). NS-CC electrode exhibits excellent cycling stability and a high reversible capacity of 1101.1 mA h g−1 can be observed after 150 cycles, but for CC electrode, a lower reversible capacity of 637.1 mA h g−1 can be obtained. Rate performance is very important for LIBs, especially application in electric vehicles. Therefore, the rate performance of the samples were evaluated at various current densities from 0.1 to 5.0 A g−1 . As seen in **Figure 4C**, NS-CC electrode can deliver reversible capacities of 1,443, 1,035, 954, 884, 802, and 689 mA h g−1 at 0.1, 0.2, 0.5, 1.0, 2.0, and 5.0 A g−1 , respectively. In contrast, the reversible capacities are 1,020, 655, 541, 478, 427, and 370 mA h g−1 at 0.1, 0.2, 0.5, 1.0, 2.0, and 5.0 A g−1 for CC electrode. Obviously, the NS-CC electrode exhibits excellent rate capability. Therefore, the long-term cycling at high current density was also tested for NS-CC electrode, the results as shown

(E) N spectrum (F) S2p spectrum electrons of NS-CC.

in **Figure 4E**. NS-CC electrode can deliver a high reversible capacity of 531.2 mA h g−1 can be obtained even after 5,000 cycles at 10.0 A g−1 . However, CC electrode only delver a lower reversible of 283 mA h g−1 after 5,000 cycles at the same current density. Therefore, NS-CC electrode exhibits excellent rate performance and long-term cycling stability, which shows better electrochemical performance than previous reported carbonbased materials, as shown in Table S1. As a result, The NS-CC electrode delivered amazing electrochemical performance especially with ultrahigh specific capacity and rate capability which can be given rise to the following reasons: (1) The carbon materials interlayer spacing are expanded by N and S successfully co-doping which benefit Li-ions diffusion. (2) The marked large surface of carbon materials offer plentiful micropores and mesopores structure shorten the diffusion distance, sufficient contact between electrolyte and electrode and active sites for lithium ion storage. (3) The ample pyridinic N, pyrrolic N and -S-C-S- covalent bonds built adequate active sites to improve surface capacity contribution (Xia et al., 2017).

The electrochemical impedance spectroscopy for CC and NS-CC electrode were further investigated to understand the significantly improved electrochemical performance. As shown in Figure S3, the Nyquist plots of CC and NS-CC electrode have shown the typical characteristics of one semicircle and a sloping straight line (Liu et al., 2016). The diameter of the semicircle is reduced in the plots of the NS-CC electrode compared with that of the CC electrode, indicating the decreased charge-transfer resistance at the electrode/electrolyte interface after doping of N, S atoms into the carbon. On the other hand, the charge transfer resistance presents a decreasing trend along with the cycles for both CC and NS-CC electrode, which due to formation of stable SEI film and the process of activation after cycling (Pan et al., 2016).

In order to further understand the high-rate performance, the capacitive behavior of the NS-CC and CC electrode were investigated and their kinetics were also analyzed with CV measurements. **Figure 5A** and Figure S4A show the CV curves of NS-CC electrode at various scan rates ranging from 0.1 to 10 mV s−1 . All of the curves display a similar shape, two cathodic peaks and one anodic peak are evidently on each curve. The peak current is not proportional to the square root of the sweep rate (v), indicating that the charge/discharge process is comprised of

faradic and non-faradic processes. According to the equation of the relationship of i and v:

$$i = a\nu^b \tag{1}$$

or

$$
\log(i) = b \times \log(\nu) + \log(a)' \tag{1}
$$

Here, a and b are constants. The process is an ionic diffusion controlled behavior when b-value is equal to 0.5, while is Li+ capacitive behavior when b-value is equal to 1.0. **Figure 5B** presents log(i)-log(v) plots for NS-CC electrode on the CV curves at peak 1, 2 and 3 potentials and the b-value are 0.84, 0.75 and 0.77, respectively. And for CC the b-value are 0.75, 0.71, and 0.97, respectively (Figure S4B). Therefore, It can be seen that all these values of b indicate fast kinetics resulting from the pseudocapacitive effect. Moreover, To quantify the pseudocapacitive contribution, we can divide the current response i at a fixed potential V into pseudocapacitive (k1v) and diffusion-controlled contributions (k2v 0.5) by following equation (Muller et al., 2015).

$$i(V) = k\_1 \nu + K\_2 \nu^{1/2}$$

By calculating both k1 and k2 constants, the overall contribution of pseudocapacitor at various scan rates can be obtained. The detail pseudocapacitive contribution of NS-CC and CC electrode at 10 mV s−1 as shown in **Figure 5C** and Figure S4C, in which 78.9% as capacitive (red region). Therefore, all contribution ratios of the capacitive capacity at scan rates of 0.1, 0.2, 0.5, 1, 2, and 5 mV s−1 were also obtained. **Figure 5D** and Figure S4D shows contributions of the pseudocapacitive behaviors at various scan rates. The proportion of capacitive contribution for NS-CC electrode are 29.8, 31.7, 36.9, 43.5, 52.1, 64.7, and 78.9% at 0.1, 0.2, 0.5, 1, 2, 5, and 10.0 mV s−1 , respectively. Furthermore, the contribution for CC electrode are 27.1, 28.2, 32.3, 41.3, 50.1, 62.6, and 75.6% at 0.1, 0.2, 0.5, 1.0, 2.0, 5.0, and10.0 mv s−1 . As a

result, the capacitive contribution for CC is smaller than the NS-CC electrode, which can attribute to the NS-CC electrode with larger specific area after N and S atoms co-doping, resulting in enhanced pseudocapacitive contribution (Augustyn et al., 2014; Chao et al., 2016).

## CONCLUSION

In summary, highly porous carbon were prepared by using cotton as precursor with a sample method. Subsequently, the sulfurnitrogen co-doped carbon were obtained via heat treatment the carbon in presence of thiourea, which can induce the defects and the expanded interlayer of the carbon. Therefore, the expanded interlayer and defects can reduce the diffusion distance of Li ions as well as offer more active sites for lithium storage. As a result, S, N co-doped carbon exhibits excellent electrochemical performance when evaluated as anode materials for Lithium-ion materials. The as-prepared S, N co-doped carbon can deliver a high reversible capacity of 546.4 mA h g−1 even after 5,000 cycles at 10 A g−1 . Moreover, excellently rate capability also can be observed, a high capacity of 600 mA h g−1 can be obtained at 5.0 A g−1 . This superior lithium storage performance of S, N codoped carbon make it as a promising low-cost and sustainable anode material for lithium ion batteries.

## AUTHOR CONTRIBUTIONS

JX conducted the experiments CY is the supervisor of this research work. JX and QP helped writing. JX, FZ, XX, DH, and CH performed the characterization and data analysis. All authors involved the analysis of experimental data and manuscript preparation.

### ACKNOWLEDGMENTS

We gratefully acknowledge the financial support from the Science and Technology Planning Project of Guangdong Province, China (No. 2017B090916002), Guangdong Natural Science Funds for Distinguished Young Scholar (2016A030306010), Guangdong Innovative and Entrepreneurial Research Team Program (2014ZT05N200) and Fundamental Research Funds for Central Universities, China (2017 ZX010).

### SUPPLEMENTARY MATERIAL

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fchem. 2018.00078/full#supplementary-material

### REFERENCES

as anode material for high-rate and long life Li-ion batteries. J. Mater. Chem. A 5, 4576–4582. doi: 10.1039/C6TA10932B

**Conflict of Interest Statement:** DH and CH were employed by company Jiangsu Key Lab of Silicon Based Electronic Materials.

The other 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 © 2018 Xiong, Pan, Zheng, Xiong, Yang, Hu 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 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.

# High-Level Heteroatom Doped Two-Dimensional Carbon Architectures for Highly Efficient Lithium-Ion Storage

Zhijie Wang1,2†, Yanyan Wang2†, Wenhui Wang3†, Xiaoliang Yu4 , Wei Lv 2 , Bin Xiang1 \* and Yan-Bing He2 \*

*1 CAS Key Lab of Materials for Energy Conversion, Department of Materials Science and Engineering, Synergetic Innovation Center of Quantum Information Quantum Physics, University of Science and Technology of China, Hefei, China, 2 Engineering Laboratory for the Next Generation Power and Energy Storage Batteries, Engineering Laboratory for Functionalized Carbon Materials, Graduate School at Shenzhen, Tsinghua University, Shenzhen, China, 3 China Key Laboratory of Optoelectronic Devices and Systems of Ministry of Education and Guangdong Province, College of Optoelectronic Engineering, Shenzhen University, Shenzhen, China, 4 Center for Green Research on Energy and Environment Materials, National Institute for Materials Science, Tsukaba, Japan*

#### Edited by:

*Qiaobao Zhang, Xiamen University, China*

### Reviewed by:

*Huan Pang, Yangzhou University, China Chenghao Yang, South China University of Technology, China Bote Zhao, Georgia Institute of Technology, United States*

#### \*Correspondence:

*Bin Xiang<EMAIL_ADDRESS>Yan-Bing He<EMAIL_ADDRESS>
*†These authors have contributed equally to this work.*

#### Specialty section:

*This article was submitted to Physical Chemistry and Chemical Physics, a section of the journal Frontiers in Chemistry*

> Received: *16 February 2018* Accepted: *20 March 2018* Published: *05 April 2018*

#### Citation:

*Wang Z, Wang Y, Wang W, Yu X, Lv W, Xiang B and He Y-B (2018) High-Level Heteroatom Doped Two-Dimensional Carbon Architectures for Highly Efficient Lithium-Ion Storage. Front. Chem. 6:97. doi: 10.3389/fchem.2018.00097* In this work, high-level heteroatom doped two-dimensional hierarchical carbon architectures (H-2D-HCA) are developed for highly efficient Li-ion storage applications.  More importantly, the H-2D-HCA shows abundant heteroatom functionality, with sulfur (S) doping of 0.9% and nitrogen (N) doping of as high as 15.5%, in which the electrochemically active N accounts for 84% of total N heteroatoms. In addition, the H-2D-HCA also has an expanded interlayer distance of 0.368 nm. When used as lithium-ion battery anodes, it shows excellent Li-ion storage performance. Even at a high current density of 5 A g−1 , it still delivers a high discharge capacity of 329 mA h g−1 after 1,000 cycles. First principle calculations verifies that such unique microstructure characteristics and high-level heteroatom doping nature can enhance Li adsorption stability, electronic conductivity and Li diffusion mobility of carbon nanomaterials. Therefore, the H-2D-HCA could be promising candidates for next-generation LIB anodes.

Keywords: 2D carbon nanomaterials, hierarchical structure, high-level heteroatom doping, Li-ion batteries, high-rate capability

### INTRODUCTION

Lithium ion batteries (LIBs) have been regarded the most important power sources for portable electronic devices and promising candidates to power future electric vehicles (Armand and Tarascon, 2008; Geng et al., 2018). In order to meet the increasing demand for energy density and fast discharge-charge abilities, it is urgent to develop LIB electrode materials with higher specific capacities, better rate capabilities and excellent cycle stabilities (Arico et al., 2005; Zhang et al., 2018). Graphite has served as the most popular anode materials for its low price, appropriate working voltage platform and high Columbic efficiency (Yazami and Touzain, 1983; Kaskhedikar and Maier, 2009; Lu et al., 2017).

**428**

Unfortunately, it suffers from limited Li storage capacity (372 mA h g−1 , according to the intercalation mechanism with the formation of LiC6 composites) and poor rate performance, which cannot satisfy the practical application requirements (Yazami and Touzain, 1983; Lu et al., 2017). Therefore, developing advanced alternative materials to replace graphite has attracted great research interest in recent years (Kaskhedikar and Maier, 2009; Zhang Q. et al., 2016; Wu et al., 2017; Deng et al., 2018).

Current researches have already proved that carbon nanomaterials delivered better LIB performance than graphite, since abundant Li-ion storage sites and rapid ion diffusion channels can be provided (Zhou et al., 2003; Dai et al., 2012; Zhang Q. et al., 2016; Wang et al., 2018). In addition, successful structural modification could further enhance the electrochemical performance of nanocarbons (Wu et al., 2003; Landi et al., 2009). Heteroatom doping plays an important role in the modification because it can adjust the physical and chemical properties of carbon nanomaterials (Wu et al., 2011). For instance, both experimental and theoretical results demonstrated that nitrogen (N) doping can positively affect the electric conductivity and electrochemical activity of nanocarbons (Ma et al., 2012; Zheng et al., 2014). Especially pyrrolic N (N-5) and pyridinic N (N-6) are able to create active sites for Li-ion adsorption in the carbon framework. Hence, increasing the doping concentration of N-5 and N-6 is beneficial to the LIB performance of nanocarbon electrodes (Wang et al., 2011; Ma et al., 2012; Mao et al., 2012; Zheng et al., 2014). Besides, sulfur (S) heteroatoms can enlarge the interlayer distance of carbons because of the larger covalent radius (102 pm) compared with that of C (77 pm) (Qie et al., 2015; Xu et al., 2016). The enlarged interlayer distance facilitates the insertion-extraction of electrolyte ions, and thus is able to improve rate capabilities (Qie et al., 2015; Xu et al., 2016; Liang et al., 2018). It is worth noting that, Li-ion storage performance of carbon nanomaterials can be promoted not only by improving heteroatom doping concentration, but also by the synergistic effects between different kinds of dopants (Ai et al., 2014). Therefore, highlevel N, S co-doping can be an effective strategy to achieve high-performance nanocarbon anodes.

In the previous work, we constructed two-dimensional (2D) hierarchical carbon architectures (2D-HCA) with N, S codoping nature for superior LIB anodes (Wang et al., 2016). Here, we further developed high-level heteroatom doped 2D hierarchical carbon architectures (H-2D-HCA) with enhanced Li-ion storage performance by increasing the N concentration in 2D-HCA. The obtained H-2D-HCA contains a much higher heteroatom concentration of 16.4%, with 15.5% of N and 0.9% of S. Interestingly, the electrochemically active N, i.e., N-5 and N-6, accounts for the majority of total N atoms (47 and 37%, respectively). Furthermore, it has an expanded interlayer space of 0.368 nm compared with that of graphite, which can enhance the Li ion diffusion speed. Benefiting from such high doping level and the unique microstructure, the H-2D-HCA can be used for highly efficient Li-ion storage. To be specific, even at a high current density of 5 A g−1 , it still delivered a high specific capacity of 329 mA h g−1 after 1,000 cycles.

## EXPERIMENTAL METHODS

### Preparation of H-2D-HCA

Mg-Al layered double hydroxides (Mg-Al LDH) and Mg-Al layered double oxides (Mg-Al LDO) were prepared following our previous work (Wang et al., 2016). Typically, Mg(NO3)2·6H2O (12.82 g, 99.9%, Macklin), Al(NO3)2·9H2O (9.38 g, 99.9%, Macklin) and urea (90.09 g, 99.9%, Macklin) were dissolved into 500 ml water in a round-bottom flask. Then the solution was heated 100◦C and kept for 12 h under magnetic stirring. After that, the temperature was decreased to 94◦C and kept for another 12 h without stirring. The produced white powder was collected with vacuum filtration and freeze dried, then Mg-Al LDH was obtained. As-prepared Mg-Al LDH was calcined at 500◦C for 3 h in a muffle furnace and ground to get Mg-Al LDO. To adsorption OII, 150 ml water was filled into a three-mouth-flask. After driving the dissolved air with Ar gas flow, 0.6 g Orange II (>85%, Macklin) was dispersed into the DI water. Then, 1 g Mg-Al LDO was added to the solution to adsorb the dye for 48 h. Magnetic stirring and Ar gas flow were maintained during the overall adsorption process. Afterwards, the solution was vacuum filtrated and washed with water for several times, then the obtained powder (RLDH/OII) was freeze dried and ground. The obtained RLDH/OII was mixed with 0.6 g melamine (99%, Macklin), and the mixture was dispersed into 50 mL methanol in a beaker. After that, the solution was heated to 50◦C under magnetic stirring. The temperature was kept until the methanol was totally evaporated, and then the RLDH-OII-melamine mixture (RLDH-OII-M) was collected. The carbonization of RLDH-OII-M was performed at 800◦C for 2 h under Ar atmosphere with a heating rate of 2◦C/min. Then the obtained black powder was washed with HCl (6M) and NaOH (2M) respectively at 70◦C for several hours. After freeze drying process, the target product H-2D-HCA was obtained.

2D-HCA was synthesized as reference samples followed the above procedures except melamine mixing process. OII and OIImelamine mixture (mass ratio: 1:1) were carbonized and washed with the above-mentioned conditions to get the reference sample C-OII and high-level doped C-OII (H-C-OII), respectively.

## Microstructure Characterization

The morphology observation and elemental mapping analyzing were conducted by using a field electron microscopy (FESEM, ZEISS SUPRA55). The microstructures of the four samples were examined by a field high resolution transmission electron microscopy (HRTEM, FEI Tecnai G2 F30). A Rigaku D/max 2500/PC diffractometer (Cu-Kα radiation with λ = 1.5418 Å) was used to measure the powder X-ray diffraction (XRD) patterns of samples. A HORIBA Labram HR Evolution Raman spectrometer was used to measure the Raman spectra of the samples (the excitation wavelength of the leaser is 532 nm). The XPS analyses were carried out on a PHI 5000 VersaProbe II spectrometer using monochromatic Al K(alpha) X-ray source. The N2 adsorption/desorption isotherms were measured by using a Micromeritics ASAP 2020 automated adsorption apparatus at 77 K. The specific surface areas was determined based on Brunauer-Emmett-Teller (BET) equation, and the pore-size

distribution was calculated by utilizing density functional theory (DFT).

### Electrochemical Tests

Li-ion storage performance of these samples was tested with halfbattery methods by using coin-type cells (CR2032). To prepare the working electrodes, a slurry consisted a mixture of 70 wt % of H-2D-HCA (or other three carbon samples), 20 wt % of poly (vinylidene fluoride) (PVDF) and 10 wt % of acetylene black in N-methyl pyrrolidone was coated on a copper foil. After being dried at 80◦C for 12 h, the electrode was punched into disks with a diameter of 12 mm. The mass loading of active materials was ∼ 1.2 mg cm−2 . 1.0 M LiPF6 dissolved in ethylene carbonate/diethyl carbonate (EC/DEC, with a volume ratio of 1:1) was used as the electrolyte. Lithium foil was used as both the counter electrode and reference electrodes. Battery assembling was carried out in an Ar-filled glove box with both moisture and oxygen concentrations below 1 ppm. Cyclic voltammetry was measured using an electrochemical workstation (CHI 660E) in the potential window of 0.005–3 V vs. Li/Li+ with a scanning rate is 0.2 mV s−1 . The cycling performance and rate capabilities were tested at in the potential window of 0.005–3 V vs. Li/Li+ using a LAND battery tester (CT2001A).

## RESULTS AND DISCUSSION

H-2D-HCA was synthesized based on a template-assistant method that we reported before (Wang et al., 2016). Mg-Al layered double oxides (Mg-Al LDO, LDO for short) were used as the templates. Orange II (OII for short, an N- and Scontaining organic dye) was used as both carbon precursors and heteroatom sources, and melamine were used as extra N sources. The preparation approach is illustrated in **Figure 1** Firstly, Mg-Al layered double hydroxides (Mg-Al LDH, LDH for short) were calcined to derive LDO. Afterwards, the LDO was dispersed into OII solution to adsorb this organic dye. During this process, the LDO was rehydrated to LDH (RLDH). More importantly, a morphology change has occurred simultaneously. This should be attributed to the morphology change occurred in the rehydration process of LDO (Wang et al., 2016). As a result, the obtained RLDH-OII composite had a 2D hierarchical structure, rather than a 2D smooth plating structure like LDH or LDO. After that, the RLDH-OII was mixed with melamine in methanol under magnetic stirring, and a RLH-OII-melamine mixture (RLDH-OII-M) was collected. Then, the RLDH-OII-M was heated at 800◦C to carbonize the organic, meanwhile, RLDH was calcined to LDO again. Finally, the obtained LDO-carbon (LDO-C) mixture was washed with NaOH and HCl to remove the LDO templates, and H-2D-HCA was achieved at last. Detailed procedures can be seen in the Experimental Methods section.

Both LDH and LDO have a hexagonal plating morphology with a smooth surface (Supplementary Figure 1). However, after OII adsorption and melamine mixing process, RLDH-OII and RLDH-OII-M show a hierarchical structure of some small-size sheets decorating on the surface of the hexagonal plates (Supplementary Figure 2). This should be attributed to the morphology change during the rehydration process of LDO (Wang et al., 2016). After carbonization and template removal process, the obtained H-2D-HCA has successfully maintained this kind of hierarchical structure, in which small carbon nanosheets growing on the surface of large-size hexagonal carbon nanoplates (**Figure 2A**). The diameter of H-2D-HCA is ∼2.2µm (**Figure 2B**) and the average thickness is ∼200 nm (Supplementary Figure 3). The small carbon nanosheets growing on the surface of H-2D-HCA have a size of ∼100 nm and a thickness of ∼10 nm (Supplementary Figure 3). The adjacent carbon nanosheets have formed many half-open pores with diameters of 10 of nanometers (Supplementary Figures 3, 4). These pores can act as reservoirs to storage electrolyte and guarantee good contact of the nanocarbon electrodes with electrolyte. 2D-HCA has similar hierarchical structure and the size with H-2D-HCA (Supplementary Figure 5), which was also reported in the previous work (Wang et al., 2016). This kind of hierarchical structure can alleviate aggregation issues of carbon nanomaterials, which caused by π-π interaction between carbon layers, and thus decrease the electrochemically active surface loss (Zhao et al., 2014; Yu et al., 2015). For comparison, H-C-OII and C-OII have irregular structure with size of about tens of micrometers (Supplementary Figure 5). These structures do not benefit electrolyte access to carbon electrodes and hinder the diffusion of Li ions.

The microstructure of H-2D-HCA was investigated by high-resolution transmission electron microscopy (HRTEM). Interestingly, H-2D-HCA has an amorphous nature and graphitic micro-crystallites with an average interlayer distance of 0.368 nm (**Figure 2C**), which is larger than that of graphite (Ou et al., 2017; Yang et al., 2018).The broaden peaks in the X-ray diffraction (XRD) pattern also confirms the amorphous nature of H-2D-HCA (Supplementary Figure 6). The peak located at 24.2◦ can be assigned to (002) diffraction peak of graphite (Xu F. et al., 2015; Zheng et al., 2015; Zhao et al., 2016; Yang et al., 2017). This peak corresponds to an interlayer distance of 0.367 nm (calculated based on the Bragg's law), which agrees well with the HRTEM results. The enlarged interlayer distance can promote Liion diffusion and help to improve the high rate performance (Qie et al., 2015; Xu et al., 2016; Liang et al., 2018). 2D-HCA has a similar amorphous structure with few graphitic layers, while H-C-OII and C-OII only show amorphous porous structure without graphite micro-crystallites could be observed (Supplementary Figure 7).

**Figure 2D** shows SEM elemental mapping. It revealed that H-2D-HCA contained not only C elements, but also abundant N and S heteroatoms. The doped N and S distribute uniformly throughout the carbon framework.

X-ray photoelectron spectroscopy (XPS) was utilized to evaluate the surface chemistry in H-2D-HCA. N and S heteroatom concentrations can be determined to be 15.5 and 0.9%, respectively (Supplementary Figure 8). The total N and S heteroatom doping level of 16.4% in H-2D-HCA, which is much higher than that of most N, S dual-doped carbon nanomaterials (Supplementary Table 1) (Ai et al., 2014; Sun et al., 2015; Xu G. et al., 2015; Zhou et al., 2015; Zhuang et al., 2015; Shan et al., 2016; Zhang et al., 2016a). The high-resolution spectrum of C1s can be fitted to four peaks (**Figure 3A**). The peak located

at 284.4 eV can be attributed to C-C/C=C of carbon; and the peak located at 285.1 eV reveal the presence of C-N and C-S (Xu G. et al., 2015; Zhuang et al., 2015; Zhang et al., 2016a). The peaks located at 286.2 and 288 eV can be assigned to C-O and C=O, respectively. The high-resolution N1s spectrum can be deconvoluted into three different peaks located at 398.3, 300.8, and 401.1 eV, corresponding to pyridinic-N (N-6), pyrrolic-N (N-5) and graphitic-N (N-Q), respectively (**Figure 3B**; Wang et al., 2011; Ma et al., 2012; Mao et al., 2012; Zheng et al., 2014; Li et al., 2017). N-6, N-5, and N-Q accounted for 37, 47, and 16% of total N atoms, respectively (**Table 1**). The electrochemically active N species, N-6 and N-5, occupied a large proportion of total N atoms (84%). The fine split peaks in high-resolution of S2p spectrum indicated the presence of C-S (163.5 eV), C=S (164.8 eV) and SOX group (167.7 eV) (**Figure 3C**; Zhang X. et al., 2016). As demonstrated by other work, S heteroatom can enlarge the interlayer distance of carbon nanomaterials because of the larger covalent radius (Qie et al., 2015; Xu et al., 2016).

The extended interlayer space can promote insertion-extraction speed of Li ions in carbon nanomaterials and thus was expected to improve the fast charge-discharge properties. The elemental composition of H-2D-HCA, 2D-HCA, H-C-OII, and C-OII are summarized in **Table 1** (the data of H-C-OII and C-OII were summarized based on XPS results shown in Supplementary Figures 9, 10, and the data of 2D-HCA were calculated from the Reference of Wang et al., 2016). Obviously, H-2D-HCA has the highest doping level and electrochemically active N proportion. Besides, because of the higher N concentration, H-2D-HCA has a much higher heteroatom doping level than 2D-HCA, and H-C-OII also has a much higher heteroatom doping level than C-OII (**Table 1**). This should be attributed to the use of melamine as extra N source during the synthesizing approach (Sheng et al., 2011).

The N2 adsorption/desorption isotherms of H-2D-HCA exhibit a combination of type I and type II characteristics, with a distinct hysteresis lop at relative pressure P/P0 ranging from 0.42 to 1 (**Figure 3D**). The specific surface area (SSA) can be calculated to be 535 m2 g −1 , which is lower than that of 2D-HCA (Wang et al., 2016). The lower SSA could be attributed to melamine decomposition during carbonization process, and this phenomenon was also reported by other previous work (Zheng et al., 2011). Pore-size distribution (PSD) curves reveal that the H-2D-HCA had both micropores which peaks located at 0.52, 1.23, and 1.71 nm, and mesopores which peaks located at 3.97 nm (**Figure 3E**). The presences of mesopores with size ranging from 6 to 30 nm can also be seen in the enlarged PSD curve (inset of **Figure 3E)**. The co-existence of micropores and mesopores in H-2D-HCA suggests its hierarchical porous

TABLE 1 | Surface physiochemical properties of various samples by XPS tests.

structure (Wang et al., 2008; Hu et al., 2015). The hierarchical porous structure benefits LIB performance of nanocarbon electrodes, in which the micropores offer abundant active sites for Li-ion storage and the mesopores can promote the rapid diffusion of Li ions (Zhao et al., 2012). By comparison, the pores in H-C-OII (Supplementary Figure 11) and C-OII (Wang et al., 2016) are mainly micropores (size less than 2 nm), with the much less presence of mesopores.

**Figure 3F** shows the Raman spectrum of H-2D-HCA. Two remarkable peaks of D band located at 1,341 cm−1 and G band at 1,581 cm−1 can be observed. As is known, G band corresponds to the zone center E2g mode related to phonon vibrations in sp2 carbon materials, and D band related to structural defects (such as edge, heteroatoms, etc.) and partially disordered structures of the sp2 domains framework (Ferrari et al., 2006). The high ID/IG ratio of 1.15 for H-2D-HCA and 1.11 for H-C-OII (Supplementary Figure 12) indicated that they have low graphitization degree and abundant defects (Ferrari et al., 2006). The defects mainly include the heteroatoms doped in the carbon framework. In contrast, because of the lower heteroatom doping level, 2D-HCA and C-OII have lower ID/IG ratio of 0.96 and 0.98, respectively (Supplementary Figure 12).

Li-ion storage properties of H-2D-HCA were evaluated using half-battery test methods. Cyclic voltammetry (CV) curves suggest that H-2D-HCA has similar Li-ion storage behaviors with other carbon nanomaterials. As shown in **Figure 4A**, in the first cycle, the pronounced irreversible anodic peak at around 0.5 V relates to the electrochemical decomposition of electrolyte and the formation of solid electrolyte interface (SEI) on the huge surface (Sun et al., 2015; Xu G. et al., 2015; Deng et al., 2016; Zhang et al., 2016b). The followed anodic peak near 0 V corresponds to the electrochemical intercalation of Li ions into graphitic structures. The cathodic peaks at 0.2 and 1.2 V can be assigned to Li ions extraction from graphitic layers and defect sites, respectively (Zhang et al., 2016b). The curves in the following cycles overlap together, suggesting the stable formation of SEI layers. When applied as anodes for LIBs, H-2D-HCA exhibits high specific capacity and excellent rate performance. As shown in **Figure 4B**, H-2D-HCA delivers an initial discharge capacity of 1,861 mA h g−1 at current density of 200 mA g−1 . When the current density increases to 500, 1,000, 2,000, 5,000 mA g −1 , the discharge capacities of H-2D-HCA were 756, 636, 504, and 348 mA h g−1 , respectively (measured from the middle cycle in each current density). After that, with the current density returns to 2,000, 1,000, 500, and 200 mA g−1 , the discharge capacity recovers to the initial capacity values with only little fade.

Long-term cyclability at high current density has been tested to further examine Li-ion storage performance of H-2D-HCA. For comparison, 2D-HCA, H-C-OII and C-OII have been tested at the same condition. First, all the samples were firstly activated at a low current density of 200 mA g−1 for two cycles, and then the current density was directly increased to 5 A g−1 . The voltage profiles of these samples in the first two cycles are shown in Supplementary Figure 13. H-2D-HCA delivered an initial discharge capacity and charge capacity of 1,924 and 1,824 mA h g−1 , together with an initial columbic efficiency (ICE) of 56.4% (**Figure 4C**). Because of the large surface area and porous structure, the H-2D-HCA delivered much higher specific capacity than graphite. The irreversible capacity loss in the first cycle is due to the SEI formation and the irreversible insertion of Li ions into micropores. H-2D-HCA has the highest ICE in these four samples (Supplementary Table 2). However, it was still far from satisfaction for practical application. Reducing surface area, coating with dense carbon layer, optimizing pore structure and pre-lithiation may have positive effects on improving the ICE. In addition, it is noted that the ICE of H-2D-HCA is higher than that of 2D-HCA, and the ICE of H-C-OII is higher than that of C-OII, suggesting that increasing N doping level may improve the ICE of nanocarbon anodes. When the current density increased to 5 A g−1 , H-2D-HCA delivers a discharge capacity of 660 mA h g−1 and a CE of 77.9%. The CE rises to a value of higher than 97% in the followed cycles (**Figure 4C**). Then the discharge capacity of H-2D-HCA drops to 450 mA h g−1 in 25 cycles, and increases afterwards, reaches 550 mA h g−1 in the 196th cycles. After that, the discharge capacity remains stable only with slight decay. Even after 1,000 cycles, H-2D-HCA can still deliver a high discharge capacity of 329 mA h g−1 , and a high CE of 99.2% (**Figure 4C**). In contrast, 2D-HCA delivered a first discharge capacity of 713 mA h g−1 and a corresponding CE of 65% at 5 A g−1 . The higher initial capacity of 2D-HCA can be attributed to the larger SSA. The following discharge capacity of 2D-HCA gradually decreases with cycling. After 1,000 cycles, the discharge capacity fades to 222 mA h g−1 , which is much lower than that of H-2D-HCA (**Figure 4C**). As for H-C-OII and C-OII, they discharge capacity respectively are 88 and 78 mA g −1 after 1,000 cycles, only about a quarter of that of H-2D-HCA (**Figure 4C**). The corresponding CE of 2D-HCA, H-C-OII, and C-OII were shown in Supplementary Figure 14, which were

worse than that of H-2D-HCA. In addition, the electrochemical performance of H-2D-HCA is better than most of N, S co-doped carbon nanomaterials (summarized in Supplementary Table 1).

In view of above results, benefiting from the unique microstructure, H-2D-HCA delivers highly efficient Li-ion storage performance. We performed first principle calculations to further understand the effect of high-level heteroatom doping on the electrochemical performance of carbon nanomaterials. The influences of dopant types and interlayer distance on Li adsorption energy, electrical conductivity and Li ion diffusion barriers were studied in this work.

To evaluate the stability of Li ions adsorbed on the carbon systems, the adsorption energy was calculated with the followed equation:

$$E\_{abs} = E\_2 - E\_1 - E\_{Li},\tag{1}$$

Where Eabs is the adsorption energy, E2 is the total energy of the geometry optimized structure with the absorbed Li atoms, E1 is the energy of different carbon systems, and ELi is the energy of single Li atom in bulk form (Zhou et al., 2004). Different carbon systems with an expanded interlayer distance of 0.368 nm were considered and their Eabs have been calculated: pure carbon (P-C), N-6 doped carbon (N-6/C), N-5 doped carbon (N-5/C), N-Q doped carbon (N-Q/C), and N, S co-doped carbon (N/S/C). For comparison, Eabs of abovementioned carbon systems with an interlayer distance of 0.34 nm were also calculated. In the 0.368 nm spacing case, Eabs for the P/C, N-5/C, N-6/C, N-Q/C and N/S/C are −0.08, −2.74, −3.28, 0.58, and −1.68 eV, respectively (the detailed structures are presented in Supplementary Figure 15). Whereas in the 0.34 nm spacing case, the corresponding Eabs are 0.18, −2.61, −3.25, 0.25 and −0.94 eV, respectively (**Figure 5A**). As is known, the more negative Eabs value means more stable Li absorption (Ma et al., 2012). Obviously, the carbon system with 0.368 nm interlayer distance have enhanced Li absorption stability. In addition, N-6 and N-5 doped carbon systems have lower Eabs. This indicates that increasing doping concentration of N-5 and N-6 can obviously improve Li-ion storage ability of carbon nanomaterials. Furthermore, the carbon systems with co-doping of N and S are also have lower Eabs compared with pure carbon, indicating more stable Li adsorption.

Another factor that affects Li-ion storage performance is the electrical conductivity of nanocarbon electrodes. In general, materials conductivity is determined by their density of states (DOS) at the Fermi level. We calculated the DOS of carbon systems of pure one and N, S co-doped one (structures can be seen in Supplementary Figures 15, 16). Comparing with the pure carbon system, the N, S co-doped carbon system has higher DOS value (3.0 eV) than the pure carbon (0.63 eV) at Fermi

level (**Figure 5B**), demonstrating its better electric conductivity. The diffusion of kinetics of Li ions also plays a dominant part in the electrochemical performance of nanocarbon anodes. The diffusion kinetics are examined by investigating Li ion diffusion barriers (activation energy). With the climbing-image nudged elastic band (CI-NEB) method, we studied a specific diffusion path, in which Li ions migrate between two adjacent hollow sites (**Figure 5C**). Convergence tests show that five intermediate images are adequate to accurately describe the activation energy barriers. The calculated activation energy barriers for 0.34 and 0.368 nm are 90.1 and 71.2 m eV, respectively (**Figure 5D**). This suggests that Li ions can diffusion much easier in carbon nanomaterials with extended interlayer distance. Theoretical calculations indicate that the expanded interlayer distance and the doped high-level N and S heteroatoms in the carbon framework can effectively enhance the Li absorption stability, Li diffusion mobility and electronic conductivity, which facilitates the transportation of Li ions and electrons, and thus improve Li-ion storage performance of H-2D-HCA.

### CONCLUSIONS

In summary, by using template-assistant methods, we have developed high-level heteroatom doped two-dimensional carbon architectures (H-2D-HCA) for highly efficient Li-ion storage. The hierarchical structure can alleviate the electrochemical active surface loss, and the porous structure provided rapid diffusion channels for Li ions. More importantly, the highly concentrated heteroatoms of 0.9% sulfur and 15.5% nitrogen are able to created abundant electrochemical active sties to storage Li ions. Also the increased N doping concentration improved the ICE of carbon nanomaterials. Furthermore, the expanded interlayer distance can promote insertion-extraction speed of Li ions. First principle calculations confirmed the enhanced Li absorption stability as well as the electronic conductivity by co-doping of N and S heteroatoms and the accumulated diffusion mobility of Li ions owing to expanded interlayer space. Benefiting from these unique microstructure characteristics and high-level heteroatom doping nature, H-2D-HCA exhibited enhanced Li-ion storage performance. Even at a high current density of 5 A g−1 , it can exhibit a high discharge capacity of 329 mA h g−1 after 1,000 cycles. Because of such superior electrochemical performance, H-2D-HCA can be promising electrode candidates for fast charge-discharge Li-ion batteries.

### AUTHOR CONTRIBUTIONS

ZW and YW did the materials preparation and characterization, as well as electrochemical tests. WW did the first principle calculations. BX and Y-BH supervised this work. All the authors discussed the results and wrote the manuscript.

### FUNDING

This work was supported by the National Key Basic Research Program of China (2014CB932400), the National Natural Science Foundation of China (51672156 and 51232005), the Guangdong special support program (2015TQ01N401), Guangdong Province Technical Plan Project (2017B010119001 and 2017B090907005), Production-study-research cooperation project of Dongguan City<PHONE_NUMBER>213), Shenzhen Technical Plan Project

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high-performance sodium ion battery anode. Mater. Today Energy 8, 37–44. doi: 10.1016/j.mtener.2018.02.001

**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.

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# Effect of Nb and F Co-doping on Li1.2Mn0.54Ni0.13Co0.13O2 Cathode Material for High-Performance Lithium-Ion Batteries

Lei Ming1 , Bao Zhang1 , Yang Cao1,2, Jia-Feng Zhang1 \*, Chun-Hui Wang1 , Xiao-Wei Wang1 and Hui Li 1

*1 School of Metallurgy and Environment, Central South University, Changsha, China, 2 Medical Engineering Center, Xiangya Hospital of Central South University, Changsha, China*

#### Edited by:

*Qiaobao Zhang, Xiamen University, China*

### Reviewed by:

*Wang Renheng, Nanyang Technological University, Singapore Lingjun Li, Changsha University of Science and Technology, China Keyu Xie, Northwestern Polytechnical University, China*

### \*Correspondence:

*Jia-Feng Zhang<EMAIL_ADDRESS>
#### Specialty section:

*This article was submitted to Physical Chemistry and Chemical Physics, a section of the journal Frontiers in Chemistry*

> Received: *30 January 2018* Accepted: *08 March 2018* Published: *05 April 2018*

#### Citation:

*Ming L, Zhang B, Cao Y, Zhang J-F, Wang C-H, Wang X-W and Li H (2018) Effect of Nb and F Co-doping on Li1.2Mn0.54Ni0.13Co0.13O2 Cathode Material for High-Performance Lithium-Ion Batteries. Front. Chem. 6:76. doi: 10.3389/fchem.2018.00076* The Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x (*x* = 0, 0.01, 0.03, 0.05) is prepared by traditional solid-phase method, and the Nb and F ions are successfully doped into Mn and O sites of layered materials Li1.2Mn0.54Co0.13Ni0.13O2, respectively. The incorporating Nb ion in Mn site can effectively restrain the migration of transition metal ions during long-term cycling, and keep the stability of the crystal structure. The Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x shows suppressed voltage fade and higher capacity retention of 98.1% after 200 cycles at rate of 1 C. The replacement of O2− by the strongly electronegative F− is beneficial for suppressed the structure change of Li2MnO3 from the eliminating of oxygen in initial charge process. Therefore, the initial coulombic efficiency of doped Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x gets improved, which is higher than that of pure Li1.2Mn0.54Co0.13Ni0.13O2. In addition, the Nb and F co-doping can effectively enhance the transfer of lithium-ion and electrons, and thus improving rate performance.

Keywords: Li1.2Mn0.54Ni0.13Co0.13O2 , Nb and F co-doping, cathode material, coulombic efficiency, electrochemical property

### INTRODUCTION

Lithium-ion batteries (LIBs) have been broadly used in the portable electronics, and regarded as the most promising energy storages for hybrid electric vehicles (HEVs) and electric vehicles (EVs) (Tarascon and Armand, 2001; Armand and Tarascon, 2008; Chiang, 2010). It is generally believed that the cathode materials are the primary factors for the improvements of the lithium-ion batteries. However, conventional cathode materials, such as LiCoO2, LiMn2O4, and LiFePO4 show low specific capacity and unsatisfied energy density, which will limit further practical application in the energy storage system (Ding et al., 2011; Luo et al., 2012; Zheng et al., 2015a). Among the developed cathode materials, lithium-rich layered material attracts great attentions of scientists, due to their high capacity of above 250 mAh g−1 , high operating voltage and low cost compared with other cathode materials (Thackeray et al., 2005; Park et al., 2007). It is a pity that lithium-rich layered materials have some intrinsic drawbacks of huge initially irreversible capacity, poor rate capability, the continous capacity and voltage decay during long-term cycling, which block their practical applications (Ellis et al., 2010). The poor rate capability is due to the poor electrical conductivity of Li2MnO3 component in lithium-rich materials, while the low initial coulombic efficiency is related to the elimination of O2− make the change of structure of Li2MnO3 in charge process (Johnson et al., 2007; He et al., 2012). Additionally, the capacity and voltage decay is caused by the structure transformation and formation of passivation layer during cycling (Lu and Dahn, 2002; Armstrong et al., 2006).

To solve above problems, many methods are proposed, such as surface coating, ion doping, and particle size reducing. Generally, surface coating could effectively suppress the side reaction between lithium-rich layered material and electrolyte and elimination of the oxygen vacancies, thus improving initial coulombic efficiency and cycling stability (Li et al., 2014, 2015, 2016). The reducing particles size could shorten the pathway to enhance the rate capability (Zheng et al., 2017). However, above approaches cannot effectively suppressed the voltage decay during cycling. Bulk cationic doping with Al, Mg, Cr, Zr, and Ru could effectively suppress the migration of transition metal (TM) during cycling, and mitigate the capacity and voltage decay (Kim et al., 2006; Jiao et al., 2007; Luo and Dahn, 2011; Sathiya et al., 2013; Xu et al., 2014). In addition, O2− site is replaced by anions, such as Cl− and F−, which could be beneficial for suppressing the structure change of Li2MnO3 from the eliminating of oxygen in initial charge process, and thus Li+ could return to the material lattice in subsequent charge-discharge process (Kang and Amine, 2005; Park et al., 2005). Therefore, the bulk doping also increase the electronic conductivity, and improve the rate performance of material bulk.

Therefore, we described the incorporation of Nb5+ and F− into the Mn site and O site of Li1.2Mn0.54Ni0.13Co0.13O2 (LMNCO), respectively. The Nb5+ and F− co-doping suppress the TM migration during cycling and alleviate Li+ loss during the elimination of O2− in the charging process. In this paper, the doped Li-rich layered oxide (Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x) is synthesized by high temperature solid phase method, the effects of Nb5+ and F− co-doping on the initial coulomb efficiency, rate performance, cycle performance and work voltage are discussed in detail.

### EXPERIMENTAL

### Materials Synthesis

The fabrication process of Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x (x = 0, 0.01, 0.03, 0.05) by high temperature solid phase method is schemed in **Figure 1**. Typically, the stoichiometric Li(CH3COO)·2H2O, Mn(CH3COO)2·4H2O, Ni(CH3COO)2·4H2O, Co(CH3COO)2·4H2O, Nb2O5, LiF and citric acid were mixed with 50 wt% of deionized water by ball-milling for 8 h (all chemicals of 99% purity). The mole ratio of Li(CH3COO)·2H2O, Mn(CH3COO)2·4H2O, Ni(CH3COO)2·4H2O, Co(CH3COO)2·4H2O, Nb2O5, LiF are 1.2: 0.54-x: 0.13: 0.13: x/2: 6x (x = 0, 0.01, 0.03, 0.05), respectively. Then the mixtures were dried at 80◦C for 12 h, and ground into fine particles. Finally, the mixture powders were calcined at 550◦C for 5 h, follow at 850◦C for 15 h in air to get a set of Li-rich layered oxide materials Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x. The Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x materials with x = 0, 0.01, 0.03, 0.05 are shorted as LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5, respectively.

### Characterization

All materials were characterized through an X-ray diffraction (XRD: Rigaku, D/max 2500v/pc) with Cu Kα radiation. The scanning electron microscopy (SEM: Philips, FEI Quanta 200 FEG) and transmission electron microscopy (TEM: JEM-2010, JEOL) were applied to observed the microstructure and the structure of all materials. The elemental chemical states of all

materials were analyzed by X-ray photoelectron spectroscopy (XPS, PHI 5000VersaProbe).

### Electrochemical Evaluation

Electrochemical performance of all Li-rich layered oxide Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x were tested using CR2032 coin cell. The electrode preparation process was consisted of three steps as follow. Firstly, 80% active material (LMNCO), 10% acetylene black, and 10% polyvinylidene fluoride (PVDF) binder were mixed with NMP solvent. Secondly, as prepared viscous cathode slurry was cast on aluminum foil. Thirdly, the foil was dried at 90◦C under vacuum for 12 h. Then it was punched into 12 mm diameter disks with the loading of active cathode mass in the range of 3–4 mg cm−2 . The coin cells were assembled in an argon-filled dry box. The lithium metal and the Celgard 2500 were used as anode material and the separator, respectively. 1 M LiPF6 in ethylene carbonate/diethyl carbonate (V/V = 1:1) was used as electrolyte. The galvanostatic chargedischarge measurements were carried out on LAND CT2001A battery testing system (Wuhan, China). Cyclic voltammetry (CV) measurements were performed by IM6 electrochemical testing station at scan rates of 0.1 mV s−1 between 2.0 and 4.8 V. Electrochemical impedance spectroscopy (EIS) was conducted by IM6 electrochemical testing station between 100 kHz and 0.01 Hz by applying perturbation AC voltage signal of 5 mV.

### RESULTS AND DISCUSSION

**Figure 2** shows the XRD patterns of Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x materials (x = 0, 0.01, 0.03, and 0.05). As seen Figure S1A, the XRD pattern of LMNCO-NF0 material belongs to the layered α-NaFeO2 structure with space group R3m (Figure S1A). There is a weak diffraction peak around 20–25◦ in the XRD pattern of the LMNCO-NF0, corresponding to the short-range cation ordering of Li+ and Mn4+ in the transition metal layers, as illustrated for Li2MnO3 structure in Figure S1B (Jarvis et al., 2011). The adjacent peaks of (006)/(012) and (108)/(110) show obvious separation, indicating the perfect layer structure of LMNCO-NF0 (Gong et al., 2004). Meanwhile, the intensity ratio of I(003)/I(104) is the indication of mixing degree for transition-metal ions in the lithium layer (Zheng et al., 2015b). For LMNCO-NF0, the I(003)/I(104) value reach 1.6, suggesting low mixing degree of transition-metal ions in the lithium layer. After doping, the XRD patterns of all samples are similar to that of LMNCO-NF0, and adjacent peaks of (006)/(012), (108)/(110) and I(003)/I(104) value remain significantly unchange. The XRD patterns of LMNCO-NF1 and LMNCO-NF3 show highly pure phase, implying that Nb and F ions are successfully doped into the crystal lattice. But when the doping amount increases, the impurity phases of Li3NbO4, Nb2O5, and LiF are observed in the XRD pattern of LMNCO-NF5, due to the solid solubility of Nb and F elements in the Li1.2Mn0.54Co0.13Ni0.13O2 material is beyond the limitation. In addition, for all doped materials, the diffraction peaks of doped samples slightly shift to lower 2θ compared to that of the LMNCO-NF0, indicating that Nb and F doping can enlarge the interlayer spacing. Furthermore, the lattice parameters of all samples are calculated by Rietveld refinement, and the results are listed in **Table 1**. It is clearly seen that the values of lattice parameter a and c get higher along with increasing amount of Nb and F ion doping. While the increasing ratio of c/a for doped sample represents low mixing degree for transition-metal ions in the lithium layer. This phenomenon also suggests that cell volume is enlarged after doping with Nb5+ and F−, which is beneficial for the diffusion of the Li+ ions (Jafta et al., 2012). In addition, the average crystallite size of the LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 are calculated by using Scherrer equation [βcos(θ) = kλ/D], where β is full-width at half-maximum (FWHM) of the XRD peak and k is a constant (0.9) as given in Table S1. This suggests that the addition of Nb2O5 and LiF could suppressed the growth of crystallite size.

In order to analyze the effect of Nb and F co-doping on chemical composition of LMNCO, the stoichiometric amounts

of metal element in all samples have been determined by ICP analysis, and the result are listed in Table S2. As seen in Table S2, the molar ratio of Li: Ni: Co: Mn in pristine LMNCO is 1.213: 0.133: 0.132: 0.543, which is close to the theoretical ratio of 1.2: 0.133: 0.133: 0.54. The molar ratio of Li: Ni: Co: Mn of all doped samples (LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5) does not vary compare to pristine LMNCO, suggesting that the Nb and F co-doping not affect the chemical composition of the samples.

The SEM images of Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x materials (x = 0, 0.01, 0.03, and 0.05) are shown in Figure S2. The particle sizes of all doped samples (LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5) are slightly smaller than that of LMNCO-NF0. Particularly, the particles sizes of doped samples decrease with the increasing amount of Nb5+ and F− elements, suggesting

TABLE 1 | The crystallographic parameters of LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5, respectively.

Nb2O5 and LiF can exhibit space steric effect, thus effectively suppressing the growth of particles. The tap density of pure LMNCO and all doped samples (LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5) are listed in Table S3. The tap density the pure LMNCO and all doped samples (LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5) are 1.51, 1.50, 1.52, and 1.49, respectively. The tap density of all doped samples (LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5) not vary compared to pure LMNCO. In addition, EDS elements mapping test has been performed on the LMNCO-NF3. As seen Figure S2E, the element mappings clearly demonstrate that Ni, Co, Mn, Nb, and F elements are all homogeneously distributed in LMNCO-NF3 structure, which confirm that Nb and F are doped into the bulk material.

**Figure 3** show TEM, HETEM and selected area electron diffraction (SAED) patterns of the LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5. The interplanar spacing of the lattice fringes (003) gradually expand along with the increase content of Nb and F co-doping, which is consistent with the XRD calculated results. This phenomenon is ascribed to that the Nb5+ ions doped into the Mn4+ sites and the F− ions occupy the pack oxygen sites. The (SAED) patterns (**Figures 3A3-D3** reveal that all the samples belong to the hexagonal symmetry of the local structure.

In order to investigate the effect of Nb and F ion codoping on the oxidation states of some elements (Ni, Mn, Co,

Nb, F, O) for all samples, **Figure 4** shows the XPS spectra of Li1.2Mn0.54−xNbxCo0.13Ni0.13O2−6xF6x materials (x = 0, 0.01, 0.03, and 0.05), respectively. For the doped samples, the Ni, Mn, Co binding energy peak all shift to higher binding energy compare to that of LMNCO-NF0, resulting from the density of electron clouds reduce around LMNCO. The Nb5+ ion doping into Mn4+ site will reduce the electron clouds. The electronegative of F− is stronger than that of O2−, and the electron clouds of all transition metal elements tend to bond with F −, suggesting that the F− is successfully doping into the site of pack oxygen. The F 1s and Nb 3d binding energy peaks are not detected in XPS spectrum of LMNCO-NF0. However, after doped Nb and F elements, it is obviously observed the binding energy peak of F 1s and Nb 3d for the LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5. The peak intensities of F1s and Nb 3d increase along with the increment amount of the Nb and F doping. The above analysis suggests that Nb and F successfully doped into the Mn site and O site of Li1.2Mn0.54Co0.13Ni0.13O2.

The initial charge/discharge curves of LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 between 2.0 V and 4.8 V at 0.1 C. are showed in the **Figure 5**. All samples exhibit two plateaus. One plateau below 4.5 V is related to lithium extraction form the layered LiMO2, and the other plateau above 4.5 V corresponds to the lithium-ion extraction from the Li2MnO3 component and accompanied by the extraction of oxygen (Johnson et al., 2004). The charge/discharge capacities of LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 are 354.8/231.8, 325.1/254.1, 337.3/269.3, and 318.9/239.9, respectively. Therefore, for LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5, the initial coulombic efficiency reaches 78.2, 79.84, and 75.2%, respectively, which are higher than that of LMNCO-NF0(65.3). Furthermore, as can be seen in Table S6, the coulombic efficiency of LMNCO-NF3 is better than these previously reported articles (Wang et al., 2013; Chao et al., 2014; Jin et al., 2014; He et al., 2015; Yin et al., 2015). It is confirmed that the covalency of the metaloxygen bond and electronegativity of the dopant ions have significant influence on the degree of oxygen loss from the lattice, easily mitigating the structure change of the Li2MnO3 during charging process (Wang and Manthiram, 2013). Therefore, the initial coulombic efficiency is improved by Nb and F effective co-doping.

**Figure 6** shows the charge/discharge curves of all samples at different rates of 0.1, 0.5, 1, and 5C. It is seen that the rate performance is enhanced by moderate amount of Nb and F co-doping. The discharge capacities of LMNCO-NF1 are 254.8, 245.2, 213.5, 162 mAh g−1 at 0.1 C, 0.5 C, 1 C, and 5 C, while the discharge capacities of LMNCO-NF3 are 269.8, 257.3, 235.3, and 173.3 mAh g−1 respectively, which are higher than that of LMNCO-NF0 (231.2, 201.9, 149.9, and 70.4 at 0.1C, 0.5C, 1C,

and 5C, respectively). The improved rate performance is related to the inequitable valent doping of Nb and F element in Mn and O sites, respectively, which can increase the oxygen vacancies in material surface, enhancing the electronic conductivity of host material eventually. However, the discharge capacities of LMNCO-NF5 are 240.6, 206.7, 158, and 94.8 mAh g−1 at 0.1C, 0.5C, 1C, and 5C, respectively. The LMNCO-NF5 show relatively poor rate performance and lower initial coulombic efficiency, owing to that the excessive Nb and F co-doping can form a thick Li3NbO4/Nb2O5/LiF layer, and lengthen the lithium ion diffusion path.

**Figure 7** shows the cycling performance and selected chargedischarge cures of LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5, respectively. The discharge capacity of LMNCO-NF0 is only 115.6 mAh g−1 with capacity retention of only 76.5% after 200 cycles at 1 C. In comparison, the discharge capacity of LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 are 191.8, 221.5, and 140.9 mAh g−1 with capacity retention of 88.3, 94.2, and 89.3%, respectively, which are higher than that of LMNCO-NF0. The fading capacity of LMNCO-NF0 is attributed to the side reaction between organic electrolyte and electrode to form inactive surface layers, and the unfavorable structure change during cycling process. As surface corrosion of the sample can trigger the dissolved Mn, resulting in the capacity fade, Nb and F co-doping can effectively suppress Mn dissolution (Table S4), stabilize the surface structure. The improved cycling performance is ascribed to Nb and F co-doping. Specifically, the binding between O and Nb (Table S5), which is stronger than the Mn-O bond, and the stronger binding of Nb-O can effectively suppress the loss of oxygen from the lattice (Deng and Manthiram, 2011).

Meanwhile the strongly electronegative of F− ions can keep the structure stability of Li2MnO3. Partial substitution of O2− anions by F− is also proposed as a way to stabilize the layered structures with the formation of strong M-F bonds (Li et al., 2015).

The inhibitory effect of structural transformation of Nb and F co-doped could be seen from the charge-discharge curves. As seem **Figure 7B**, the corresponding discharge profile of LMNCO-NF0 at 1C exhibited obvious voltage fade after cycling process. The average voltage only is 2.83 V with the average voltage

of 77.5% at first cycle after 200 cycles. During the long-term cycling process, the migration of transition metal ions will result in the transformation of layered structure to spinel structure, accompanied by the continous voltage fade as the consequence (**Figure 8E**). While, after Nb and F co-doping, the voltage decay are effectively suppressed. The average voltages of LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 can reach 3.12 V, 3.23 V, and 3.01 V after 200 cycles, respectively, The Nb5+ and F− doped into the bulk material which could suppress the migration of TM ions to Li layer, and thus suppress the layered to spinel phase transformation (**Figure 8F**).

In order to further confirm the effect of Nb and F co-doping on the electrochemical performance, the cyclic voltammetry (CV) measurements from 2.0 to 4.8 V at a scan rate of 0.1 mVs−1 were carried out, as shown in Figure S3. For all samples, there are two anodic peaks at 4.08 and 4.62 V during the initial cycle. The anodic peak at 4.08 V belongs to lithium deintercalation from LMO structure accompanied by the oxidation of Ni2+/Ni4+ and Co3+/Co4+, and the other peak at 4.62 V corresponds to lithium deintercalation from Li2MnO3 structure, which is associated with the Mn4+ activation process (Yu and Zhou, 2012). However, the anodic peak at high potential of 4.62 V is ascribed to the irreversible reaction about removal of Li2O from the Li2MnO3 component, which will disappear in the following cycles (Ohzuku et al., 2011). During the reduction process, the cathodic peak at 3.75 V corresponds to the redcution of Ni2+/Ni4+ and Co3+/Co4+, and the activation of the Li2MnO3 component. there is the peak at about 3.20 V in following cathodic reaction (Jin et al., 2014). In the subsequent cycles, the peak at 4.6 and 3.9 V of all materials disappears, and the peak at 3.7 V emerges. In addition, for LMNCO-NF0, the reduction peak below 3.0 V is observed after 100 cycles, and the intensity gradually increase with the increase of the cycle, suggesting the formation of the spinel phase during the cycling. However, the reduction peak below 3.0 V for LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 are hardly observed after 100 cycles, and the reduction peak below 3.0 V appeared when the number of cycles up to 200 cycles. This result indication that the phase transformation from layered into spinel structure is effectively suppressed after Nb and F co-doping.

The results of HRTEM investigations of LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 after 100 and 200 cycles are displayed in **Figure 8**. As seen **Figure 8A**, the interplanar spacing of LMNCO-NF0 after 100 and 200 cycles show (003) planes of layer phase, and the (112) planes of the spinel phase. Moreover, LMNCO-NF0 was found to the formation of local amorphous and spinel phase domains after 100 cycles, and the local amorphous areas increase after 200 cycles. In contrast local amorphous domains after 100 cycles do not appear. While local amorphous domains and spinel phase of all doped materials after 200 cycles were observed. Meanwhile the local amorphous areas of all doped materials after 200 cycles are less than that of LMNCO-NF0. Furthermore, the (220) and (440) planes of the spinel phase not be observed after 100 cycles, but after 200 cycles emerges. The Nb and F co-doping can effectively mitigate the migration of TM

TABLE 2 | Impedance parameters derived using equivalent circuit model (Figure S4B) for LMNCO-NF0, LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 electrodes before cycling (fully discharged).

ions, and suppressed the voltage fade during high voltage cycling.

To gain insight into the effect of Nb and F co-doping on kinetic behavior of lithium ion diffusion, the EIS was carried out. Figure S4 and **Figure 9** show Nyquist plots of all samples before and after cycling, respectively. All the Nyquist plots are consist of the depressed semicircles and a slope. The high frequency intercept at the real axis belongs to the ohmic resistance (R) of interaction between electrolyte and electrode, and the semicircle in the middle frequency region corresponds to charge-transfer resistance (Rct). The slope in the low frequency region is related to the diffusion of lithium ion in the bulk material (Koga et al., 2013; La Mantia et al., 2013; Wang et al., 2015, 2017). The values of Rct and R of all materials before cycling are simulated by the equivalent circuit and the result were listed in **Table 2**. In addition, the lithium-ion diffusion coefficients of all samples are calculated by the following equations (Fey et al., 2001; Levi and Aurbach, 2004; Lin et al., 2013) and also listed in **Table 2**.

$$\mathbf{D}\_{\mathrm{Li}}{}^{+} = \frac{\mathbb{R}^{2}\mathbb{T}^{2}}{2\mathbf{n}^{4}\mathbb{F}^{4}\mathbb{C}^{2}\_{\mathrm{Li}^{2}}\sigma^{2}} \tag{1}$$

$$\boldsymbol{Z}^{'} = \boldsymbol{R}\_{\Omega} + \boldsymbol{R}\_{\text{ct}} \boldsymbol{\sigma} \boldsymbol{\alpha}^{-1/2} \tag{2}$$

Here, T is 298 K, R is gas constant (8.314 J K−1 mol−1 ), A is the surface area of the electrode, F is the Faraday constant (96,485 C mol−1 ), n is the number of electrons involved in reaction, C is the concentration of lithium ion. Where ω is the angular frequency in the low frequency region and σ is the Warburg coefficient. The graph of Z′ against ω −1/2 in the low frequency region is a straight line with the slope of σ.

As seen **Table 2**, the Rct of LMNCO-NF1, LMNCO-NF3, and LMNCO-NF5 are much lower than that of LMNCO-NF0, suggesting that the Nb and F co-doping will increase electronic conductivity of bulk material and improve the kinetics of lithiumion diffusion, due to Li slab space is enlarged. Therefore, all doped materials have higher capacity and excellent rate performance.

Meanwhile, further explain the effect of the Nb and F doped on the cycling performance were performed by EIS after the 1st, 100th, and 200th cycles. For all samples, the value of Rct is simulated by the equivalent circuit and the lithium-ion diffusion coefficients are calculated by the equation (1 and 2), while the results are listed in the Table S7. The charge transfer resistance (Rct) of LMNCO-NF0 increases continuously, while lithium ion diffusion coefficient (DLi+) decreases rapidly and remains only half of the pristine value after 200 cycles, resulting from the severe structure change during long-term cycling. While the charge transfer resistance (Rct) and lithium ion diffusion coefficient (DLi+) of all doped materials exhibit a little variation and keep acceptable values after 200 cycles, indicating the excellent reaction kinetic. These suggest that Nb and F co-doping could keep the capacity stability which could be ascribed to that doping of Nb and F ions into the bulk layered materials could suppress the change structure from a layered into a spinel structure during cycling.

### CONCLUSION

Nb and F co-doped Li1.2Mn0.54Ni0.13Co0.13O2 have been fabricated by using traditional solid phase method. The Nb5+ and F− ions are successfully doping into the Mn4+ site and O2− sites, respectively, which is beneficial for suppress the loss of the oxygen and the mixed migration of transition metal ions. Therefore, Nb and F co-doping can enhance initial coulombic efficiency of 81.4% and rate performance with discharge capacity with 269.8, 257.3, 235.3, and 173.3mAh g−1 at the discharge rates of 0.1, 0.5, 1 C, 5 C and cycling performance with discharge capacity of 221.5 mAh g−1 after 200 cycles at 1 C, as well as suppress voltage fade of Li1.2Mn0.54Ni0.13Co0.13O2 during cycling. It is convinced that the Li1.2Mn0.54Ni0.13Co0.13O2 after Nb and F co-doping can satisfy the requirements of the electric vehicle and the renewable energy storage, and become advanced lithium ion cathode materials for application of Li-ion battery.