Properties of lithium iron phosphate prepared by biomass-derived carbon coating for flexible lithium ion batteries

Properties of lithium iron phosphate prepared by biomass-derived carbon coating for flexible lithium ion batteries

Electrochimica Acta 300 (2019) 18e25 Contents lists available at ScienceDirect Electrochimica Acta journal homepage: www.elsevier.com/locate/electac...

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Electrochimica Acta 300 (2019) 18e25

Contents lists available at ScienceDirect

Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta

Properties of lithium iron phosphate prepared by biomass-derived carbon coating for flexible lithium ion batteries Hye-Jung Kim a, Geun-Hyeong Bae a, Sang-Min Lee b, Jou-Hyeon Ahn c, *, Jae-Kwang Kim a, ** a b c

Department of Solar & Energy Engineering, Cheongju University, Cheongju, Chungbuk 28503, South Korea Chungbuk Technopark 76, Yeongudanji-ro, Ochang, Cheongju, Chungbuk 28116, South Korea Department of Materials Engineering and Convergence Technology and RIGET, Gyeongsang National University, 501 Jinju-daero, Jinju, 52828, South Korea

a r t i c l e i n f o

a b s t r a c t

Article history: Received 11 October 2018 Received in revised form 17 December 2018 Accepted 10 January 2019 Available online 12 January 2019

This study highlights the effects of a biomass-derived carbon coating on the properties of flexible lithium iron phosphate polymer batteries. Pure LiFePO4 (LFP) and carbon-coated LiFePO4 (C-LFP) cathode materials are synthesized by a modified mechanical activation process. LiFePO4 was uniformly coated with biomass-derived carbon by the addition of orange peel during the synthesis. C-LFP with a particle size of approximately 90 nm, having a homogeneous particle size distribution with particles smaller than that of pure LFP (140 nm), is obtained with 6.0 wt% carbon content. The electrochemical properties are evaluated at room temperature by cyclic voltammetry and charge-discharge performance analysis using porous carbon current collector with a lithium metal anode and gel polymer electrolyte that is prepared by the phase inversion method. C-LFP exhibits a high discharge capacity of 147.3 mAh g1 (corresponding to 87% of the theoretical capacity) at 0.5 C-rate and 139.8 mAh g1 at 1 C-rate. Good rate capability and stable cycle performance were realized for flexible lithium polymer cells employing the biomass derivedcarbon coated LiFePO4. © 2019 Elsevier Ltd. All rights reserved.

Keywords: LiFePO4 Carbon coating Orange peel Biomass Flexible lithium polymer batteries

1. Introduction Lithium batteries are considered as an attractive power source for various applications, such as cellular phones, notebook computers, electric vehicles (EVs), and energy storage systems (ESSs) [1]. Currently, lithium metal oxides (LiMO2, M ¼ Co, Mn, Ni, Al) are the most commonly used as cathode material in lithium batteries. However, they are relatively expensive and have a lower stability of crystal structure. Much research has been devoted to finding low cost, effective replacements for LiMO2, especially for batteries destined for transportation applications. Lithium iron phosphate has been recognized as a good alternative cathode material for replacing LiMO2 since Padhi et al. first reported the redox reaction of lithium iron phosphate (LiFePO4) in 1997 [2]. LiFePO4 is a less expensive, abundant, environmentally friendly, and stable material with a high theoretical capacity of 170 mAh g1. It provides an operation potential of about 3.4 V versus Liþ/Li and exhibits stable

* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (J.-H. Ahn), [email protected] (J.-K. Kim). https://doi.org/10.1016/j.electacta.2019.01.057 0013-4686/© 2019 Elsevier Ltd. All rights reserved.

cell operation by providing a large over-voltage margin for decomposition of liquid electrolyte [2e4]. Optimized preparation processes facilitated to achieve the desired particles size and morphology are gaining importance to improve the properties of LiFePO4. The rate capability limitation of LiFePO4 has come from the poor electronic conductivity and the slow lithium ion transfer across the phase boundary of LiFePO4/FePO4 [2e6]. Various approaches for enhancing the electronic conductivity of LiFePO4 have been attempted, such as coating the particles with carbon. In 1999, Ravet et al. reported that LiFePO4 had a capacity of 160 mAh g1 at 1 C-rate at 80  C when coated with carbon [7]. Carbon coating is usually formed by carbonization of organic/polymeric compounds during the synthesis of LiFePO4 [8e20] or by mixing with carbon materials [21e26] and biomass materials [27e31]. Recent studies on the effect of the surface carbon structure on the electrochemical performance of LiFePO4 revealed that sp2-coordinated carbon resulted in better conductivity than sp3-coordinated carbon and in enhanced electrochemical performance [8,10,20]. Moreover, the carbon composite LiFePO4 electrode of high conductivity applied into flexible lithium ion batteries [32e34]. We previously reported the synthesis of LiFePO4 by a modified

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mechanical activation method [24,25]. The modification in the solid-state synthesis process allowed achieving phase-pure, small crystallites of the electrode material with a uniform carbon coating, which could obviously improve the electrochemical performance. Although methods of improving the electrochemical properties LiFePO4 have been studied by many research groups, the effects of biomass-derived carbon coating on the physical and electrochemical performance of LiFePO4 have not been defined, especially for application of the coated material in polymer batteries. Orange peel biomass is used as a carbon source for coating the surface of LiFePO4. The focus of the present study is to compare the physical and electrochemical performance of flexible polymer batteries employing pure LiFePO4 with that based on biomass-derived carbon-coated LiFePO4 prepared by a modified mechanical activation process. 2. Experimental Pure LiFePO4 (LFP) and carbon-coated LiFePO4 (C-LFP) were synthesized by a modified mechanical activation method from the precursors (Li2CO3, FeC2O4$2H2O, and NH4H2PO4, 99% purity from Aldrich) in stoichiometric quantities. The orange peel precursor was immersed in 100 mL of 7% KOH solution at room temperature for 24 h and then dried at 80  C for carbon coating on the surface LiFePO4, as shown in Fig. 1. The KOH was used as activating agent. Pure LiFePO4 was synthesized without orange peel powder. The precursors were firstly mixed with magnetic stirring in 60 wt% of triply distilled water at room temperature for 7 h followed by reduced pressure drying of the blend at 70  C for 2 h using a rotary evaporator at 60 rpm to yield a solid powder. After the rotary evaporator, they had a high-energy ball milling in a hardened steel vial with zirconia balls at room temperature for 15 h under argon atmosphere. Finally, thermal treatment was performed on the pellets at 600  C for 10 h under nitrogen atmosphere. A microporous membrane of P(VdF-HFP) (Kynar 2801) was prepared by the phase inversion method at room temperature [35]. The P(VdF-HFP) co-polymer was vacuum dried at 60  C before use. N-methyl pyrrolidone (NMP) and methanol were used as received. P(VdF-HFP) solution (16%) was prepared in NMP (as the solvent) by mechanical stirring for at least 12 h at room temperature. The resultant viscous homogeneous solution was cast on a Teflon mold with a doctor blade to form a membrane and then immersed in methanol/water (80/20, v/v) for 4 h. The collected membranes were oven dried at 60  C for 12 h and then vacuum dried at 70  C for 12 h before further use. The gel polymer electrolyte (GPE) was prepared by soaking the membrane for 3 min in a solution of 1 M LiPF6 in ethylene carbonate (EC)/dimethyl carbonate (DMC) (1:1 by vol.) (PanaX. Etec Co.). The crystal structure of the cathode material was defined by Xray diffraction (XRD: SIEMENS D5005) using Cu-Ka radiation (35 mA/40 kV) and FT-IR absorption spectra (VERTEX 80v, Bruker

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Optics). Thermal analysis (TGA) of all materials was performed on an SDT-Q600 (USA) instrument under oxygen flow at a heating rate of 10  C min1. SEM and EDX mapping using a field emission scanning electron microscope (Philips XL30 S FEG) were employed to evaluate the particle size and homogeneity of the chemical composition. The nature of the coated carbon was analyzed using high-resolution transmission electron microscopy (HR-TEM) (JEM3010, JEOL) and energy dispersive X-ray using a TEM-EDX (JEM2010, JEOL) instrument. The surface area was examined from N2 sorption data (ASAP 2020 Analyzer) using the Brunauer-EmmettTeller (BET) method and the atomic ratio was determined by inductively coupled plasma (ICP) analysis (Atomscan 25, Optima 4300DV). The carbon content of the biomass-derived CeLiFePO4 was determined by elemental analysis (CHNS-932, LECO). The electrical conductivity was analyzed by the four-point probe method using a LORESTA-GP (MCP-T600) instrument. An electron probe microanalyzer (EPMA) was used to study the distribution of grains in the cathode film. To prepare the cathode, the LiFePO4 powder, carbon black, and poly(vinylidene fluoride) (PVdF: Aldrich) binder were mixed in a ratio of 83:7:10 by weight and the mixed viscous slurry in NMP was cast on porous carbon current collector and dried at 90  C under vacuum for 6 h. The dried film was cut into disks with an area of 0.95 cm2 and total mass of ~3.1 mg for use as the coin-type cell. The active material mass loading was 2.7 mg cm2. Two-electrode cointype cells were assembled with lithium metal as the anode, gel polymer electrolyte, and LFP or C-LFP as cathode. Cyclic voltammetry was carried out in the range of 2.2e4.4 V at a scan rate of 0.1 mV s1. Constant current tests were carried out in the range of 2.5e4.0 V using an automatic galvanostatic charge-discharge unit (WBCS3000 battery cycler) at room temperature. The rate performances were investigated at different current density rates ranging from 0.5 C (0.23 mA cm2) to 1 C (0.46 mA cm2). 3. Results and discussion The chemical composition ratio of the sample matched the theoretical molar ratio of Li:Fe:P (1.01:0.99:1.00, within the error range of ICP), and elemental analysis revealed that the C-LFP composite contained 6.0 wt% carbon. The XRD spectra of the synthesized samples are shown in Fig. 2. The structures of the two different samples were identified as belonging to the Pnma space group with the orthorhombic olivine structure. The diffraction patterns obtained for the two samples are in close agreement with the standard LiFePO4 spectra, with no indication of impurity phases such as Fe2P, FeP, and Li3PO4. A previous study demonstrated that the carbon source effectively provides a carbon-rich atmosphere that prevents oxidation of the Fe2þ ions to Fe3þ. The crystal size of carbon-coated LiFePO4 was smaller than that of pure LiFePO4 [22,36,37]. Fig. 3 shows the XRD peaks of the (200), (210), (211), (020), and (311) planes for pure LFP and C-LFP. The patterns show a

Fig. 1. Process for synthesis of biomass derived-carbon coated LiFePO4.

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Fig. 2. Comparison of XRD spectra of pure LFP and biomass-derived C-LFP with standard Pnma orthorhombic LiFePO4.

clear peak shift by about 0.009 after carbon coating, indicative of a change in the lattice parameter. The refined lattice parameters of pure LFP and C-LFP are listed in Table S1. The lattice constant of the different axes, except for the c-axis, decreased after coating with biomass-derived carbon. The FT-IR absorption spectrum of pure LFP and C-LFP are shown in Fig. S1. In the FT-IR spectra, the bands in the low wavenumber region of 400e700 cm1 correspond to the local surroundings of

lithium. The characteristic peaks of LiFePO4, including the doublet at 465505 cm1, have been indexed in the spectrum [38]. The peaks in the range of 900e1200 cm1 correspond to the stretching 1 modes of the PO3 (due to 4 unit. The absence of peaks at 422 cm Li3PO4), 759, and 1178 cm1 due to LiFeP2O7 indicates that the samples are free of this impurity [39]. The peaks in the FT-IR spectra of pure LFP are clearly broader than those of C-LFP. This broadening provides evidence of a decrease in the lifetime of the phonons, and thus indicates the existence of defects that break the periodicity of the lattice sites inside the crystallites of pure LFP. It is remarkable that in the high wavenumber range, for both materials, the IR bands appear at exactly the same position. This indicates that the chemical bonds of the PO4 tetrahedra forming the backbone of the crystal phase are stable. The TG/DTA curve of pure LFP (a) and C-LFP (b) are displayed in Fig. 4. In the temperature range from room temperature to 250  C, there was some weight loss in the TG curves ((a) and (b)) that corresponds to the elimination of water. The purity of iron in pure LFP was determined by heating pure LFP to around 600  C in oxygen, which caused a weight gain corresponding to the complete oxidation of all the iron to the ferric state. If all the iron in LFP is ferrous, then a weight gain of 5.0 wt% is expected [40]. A similar weight gain was observed for C-LFP at a temperature of about 400  C. A shift in the temperature of the weight gain is sometimes observed for coated samples. C-LFP synthesized under the same conditions used for pure LFP showed less weight gain than pure LFP in the upper region of the curve. This result indicates that carbon prevents the oxidation of ferrous ions to the ferric state. When the surface of LFP was coated with carbon, a rapid weight loss was observed due to decomposition of the carbon on LFP. The carbon content in synthesized LFP was around 6.0% based on the weight

Fig. 3. Peak profiles obtained from (200), (210), (211), (020), and (311) planes of pure LFP and C-LFP in XRD pattern.

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Fig. 5. SEM images of (a) pure LFP and (b) C-LFP.

Fig. 4. TG and DTA curves of pure (a) LFP and (b) C-LFP at heating rate of 5  C min air at flow rate of 100 mL min1.

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loss in the TGA profile [41]. The carbon content of C-LFP closely matches the value determined by elemental analysis. The biomass-derived carbon coating on the LFP had a marked influence on the morphological properties of the samples. Fig. 5 shows the SEM images of pure LFP and C-LFP. The mechanical activation process has been reported to be particularly efficient for synthesizing small particles with homogenous morphology [12,22,36]. The samples of pure LFP and C-LFP synthesized in this study by the mechanical activation method also consisted of very small particles with sizes varying from 50 nm to 200 nm with a large number of particles. The average particle size was visibly smaller for C-LFP (90 nm) than for pure LFP (140 nm). This demonstrates that inclusion of a carbon source during the synthesis of LFP is advantageous for reducing the particle size. The carbon coating stabilizes the small crystals and prevents inter-particle contact to a great extent, thereby preventing enlargement of the particles at high temperatures due to coalescence [42]. The carbon coating on LFP was nanometer thick. The TEM images of pure LFP (a, b) and C-LFP (c, d) are shown in Fig. 6. The particle size of the samples in (a, c) could be clearly defined from the TEM image. Most of the particles in pure LFP and C-LFP were respectively ~150 nm and ~80 nm in size, though bigger particles were also seen, where the latter resulted from agglomeration. The average particle size of pure LFP and C-LFP was 140 and 90 nm, respectively. C-LFP comprised smaller particles with a uniform morphology that should provide a higher surface area. The larger particle size observed for pure LFP might be the

result of growth of the crystals [22]. A uniform carbon coating on the particle surface effectively enhances the electronic conductivity of olivine phosphates [10,19,43]. TEM analysis of LFP (Fig. 6(b)) showed only the bare particle surface (the LFP crystals appear as darker regions with a diagonal pattern) whereas for CLFP, a nanometer-sized web of amorphous carbon was observed (Fig. 6(d); the carbon coating appears as light regions with a wave pattern) surrounding the particle of LFP. For C-LFP, the carbon coating is seen to envelop the particle on nearly all surfaces, but the thickness of the coating is not uniform, varying from 4 to 8 nm. The porous carbon coating enveloping LFP resulted in a higher specific surface area of 18.4 m2 g1, as well as a higher electronic conductivity of 4.8  104 S cm1 for this material. Efficient carbon coating could also lead to improved charge transfer kinetics of the material and result in higher reversible capacity. The results of TEM-EDX analysis across the surface of pure LFP in the step-bystep mode are shown in Fig. S2. The average atomic weight ratios of pure LFP from the particle core to the surface were determined by EDX analysis. The average Fe, P, and O atomic weight ratios in pure LFP were 37, 21, and 42 wt% respectively, corresponding to the stoichiometric weight ratio. The elemental distribution of C-LFP was analyzed by SEM energy dispersive X-ray mapping. Focusing on the selected region shown in Fig. S3(a), the elemental mapping images of C, O, P, and Fe are shown in Fig. S3(be), respectively. The contour observed for all of these elements matched with that of the corresponding SEM image. Because carbon also matches with uniform distribution and image contrast as that of the other three elements, it is concluded that the LFP coated by carbon in a uniformly distributed morphology. The composition of C-LFP was near stoichiometric, with O, P, and Fe weight ratios of 39.7, 19.3, and 34.7 wt%, as determined by SEMEDX (f) and the carbon content (6.3 wt%) was almost in agreement

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Fig. 6. TEM images of (a) and (b) pure LFP, (c) and (d) C-LFP.

with the elemental analysis result of 6.0 wt%. To prepare flexible LFP electrode, the mixed slurry of NMP solvent, LFP powder, carbon black powder and PVdF binder was spread uniformly on the porous carbon current collector of 250 mm thickness. It is then sucked under the electrode and vacuum dried. Fig. 7(a) displays the photo image of the flexible LFP electrode and Fig. 7(b) exhibits that the LFP is well absorbed between the carbon layers. Fig. 7(cee) shows the electron probe microanalysis (EPMA) elemental mapping data for the C-LFP electrode. The elements in the C-LFP electrode were found to be evenly distributed along the surface of the electrode, where the distribution of Fe (d) and P (e) corresponded with the selected region of Fig. 7(c). From the elements detected by EPMA, it can be concluded that each particle is composed of more than one crystallite. In the present work, we employed a porous poly(vinylidene fluoridecohexafluoropropylene) {P(VdF-HFP)} membrane as the host pore for preparation of the gel polymer electrolyte. The membrane was prepared by applying the phase inversion process parameters optimized in our laboratory [35]. Fig. 7(f) displays the SEM image of the porous membrane that consists of pores with an average diameter of 90 nm. The electrolyte uptake of the polymer membrane within a time period of 1 min was 240% when soaked in organic liquid electrolyte. The porous membrane has good wettability by the liquid electrolyte and exhibits high mechanical strength of 38 MPa and self-standing properties. The thus-prepared gel polymer electrolyte exhibited a sufficiently high ionic conductivity of 6.5  104 S cm1 at 25  C. The CV data for pure LFP and C-LFP during the first cycle are presented in Fig. S4. A single redox reaction pair, characteristic of

Liþ ion insertion and extraction, was observed, and the area under the oxidation and reduction peaks remained nearly the same, showing that an equal quantity of lithium ions was reversibly extracted and inserted into the materials. The main differences between pure LFP and C-LFP are in the redox current and the potential separation (DV), appearing as a shoulder in the voltammogram. The redox current for C-LFP (0.351 mA) was more than double that of pure LFP (0.163 mA). The potential separation between the anodic and cathodic peaks (DV) was 1.5 V and 0.4 V for pure LFP and C-LFP, respectively. The smaller redox current and larger potential separation (DV) for the carbon-free LFP results from its poor electronic conductivity and limited lithium ion diffusivity due to the larger particle size compared to that of C-LFP. These kinetic limitations also lead to the shoulders in the CV curves. The improved electron conductivity and ion diffusion capabilities of CLFP could result in higher reversibility of the electrochemical reactions over repeated cycling as well. Enhancement of the chargetransfer kinetics of carbon-coated LFP has been reported earlier based on CV measurements [12]. The samples of LFP were evaluated as cathode materials by analysis of their electrochemical performance in lithium polymer cells at room temperature. The pure LFP and C-LFP conventional cells with Al substrate and [email protected] 2200 separator show 141.2 and 158.3 mAh g1 discharge capacities at 0.1 C-rate, respectively (Fig. S5). The initial charge-discharge performance of the cells with the pure LFP and C-LFP cathodes at a current density of 0.23 mA cm2 and 0.46 mA cm2 (corresponding to 0.5 and 1 Crate) are compared in Fig. 8(a). The difference between the charge and discharge flat voltages (DV) of pure LFP and C-LFP was minimal,

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Fig. 7. Photographic (a) and cross-section SEM image (b) of flexible LFP electrode using porous carbon current collector. EPMA elemental mapping of flexible LFP electrode film (cee) and (f) SEM image of phase inversion P(VdF-HFP) membrane. (c) Selected area, (d) Fe, and (e) P.

i.e., 0.14 V and 0.1 V at 0.5 C, respectively, indicating good kinetics of the redox reaction. However, the DV value of pure LFP and C-LFP increased to 0.81 V and 0.27 V at 1 C, signifying an increase in the cell resistance. Pure LFP showed the highest cell resistance. Clear flat voltage plateaus were observed at ~3.50 V for charging and ~3.35 V for discharging. The highest specific capacity was achieved for C-LFP, exhibiting respective charge capacities of 150.2 mAh g1 and 140.3 mAh g1 and discharge capacities of 147.3 mAh g1 and 139.8 mAh g1 at 0.5 C and 1 C, corresponding to 87% and 82% utilization of the LFP active material. The discharge capacities for pure LFP were 126.6 and 105.3 mAh g1 at 0.5 C and 1 C, respectively. A gradual decrease in the discharge capacity and increase in DV with increasing C-rate was evident, as is generally the case for LFP cathodes, attributed to limited lithium ion diffusion to meet the fast reaction kinetics at higher current density, especially at low temperature. The comparatively better rate capability shown by biomass derived C-LFP evidently arises from the small particles, high surface area, uniform morphology, and better conductivity achieved with the biomass-derived carbon coating. The cycle

performance of the lithium cells up to 50 cycles obtained with the cathodes of pure LFP and C-LFP at different C-rates are presented in Fig. 8(b). Good cycling ability was shown by the cathodes at all Crates evaluated herein. C-LFP (with the carbon coating) showed >95% and 89% retention of its initial discharge capacity at the 50th cycle at 0.5 and 1 C-rate, respectively. Pure LFP performed equally well at 0.5 C-rate with >91% retention of its initial discharge capacity at the 50th cycle. However, the retention was lowered to about 86% at the higher current density of 1 C. The capacity fade per cycle, calculated based on the discharge capacity of the cell initially and after 50 charge-discharge cycles, also supports this observation (the capacity fade per cycle of pure LFP and C-LFP was 0.22 and 0.29% at 0.5 C-rate and 0.14 and 0.27% at 1 C-rate, respectively). Thus, carbon-coated LFP showed slightly better performance in terms of the initial discharge capacity and capacity fade rate. This was achieved by synthesizing a phase-pure, LFP active material with small particles (~90 nm) having a very effective biomassderived carbon coating.

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is an effective technique for achieving high performance of flexible LFP battery, which results in nano-meter sized particles of the active material with a thin coating of carbon and increased electrical conductivity. Acknowledgements This research was supported by the Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (2017M1A2A2087577 and 2018R1A4A1024691). This research was partially supported by the Cheongju University Research Scholarship Grants in 2017. Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi.org/10.1016/j.electacta.2019.01.057. References

Fig. 8. Initial charge-discharge curves (a) and cycle performance (b) of Li/gel polymer electrolyte/LFP cells with cathodes of pure LFP and biomass-derived C-LFP (room temperature, 0.5, and 1 C-rate, 2.5e4.0 V).

4. Conclusion Olivine-LFP materials (pure LFP and C-LFP) were synthesized by a modified mechanical activation process followed by firing, and orange peel was used as the carbon precursor in the case of C-LFP. The effects of coating LFP with biomass-derived carbon on the physical and electrochemical performance of the resulting composite were investigated by X-ray diffraction, FT-IR, TGA, SEM, TEM, and galvanostatic charge-discharge analyses. The initial addition of dried orange peel, along with the other starting materials, resulted in nano-meter sized particles of the active material with a thin coating of carbon. Biomass derived-C-LFP exhibits a uniform morphology with the smallest average particle size of 90 nm. Thermogravimetric analysis demonstrated that the stability of CLFP derived from the carbon coating is better than that of pure LFP. Comparison of the CV data for the uncoated and carbon-coated samples clearly demonstrated the better reaction kinetics of CLFP. Evaluation of the electrochemical capacity in flexible lithium cells at room temperature showed an active material utilization of 87% and 82% at 0.5 and 1 C-rate, respectively, with C-LFP. Stable cycle performance was also achieved, even at the higher current densities of 0.5 C and 1 C. The relatively poorer performance of pure LFP is attributed to the larger particles and poor electrical conductivity, which adversely affect the rate of lithium ion diffusion in the material. This study shows that biomass derived-carbon coating

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