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Lattice Water for the Enhanced Performance of Amorphous Iron Phosphate in Sodium-ion Batteries Soo Yeon Lim, Ji Hoon Lee, Sangryun Kim, Jaeho Shin, Wonchang Choi, Kyung Yoon Chung, Dae Soo Jung, and Jang Wook Choi ACS Energy Lett., Just Accepted Manuscript • DOI: 10.1021/acsenergylett.7b00120 • Publication Date (Web): 06 Apr 2017 Downloaded from http://pubs.acs.org on April 7, 2017

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Lattice Water for the Enhanced Performance of Amorphous Iron Phosphate in Sodium-ion Batteries Soo Yeon Lim,† Ji Hoon Lee,† Sangryun Kim,† Jaeho Shin,† Wonchang Choi,‡ Kyung Yoon Chung,‡ Dae Soo Jung§ and Jang Wook Choi*† †

Graduated School of Energy, Environment, Water, and Sustainability (EEWS) and KAIST

Institute (KI) NanoCentury, Korea Advanced Institute of Science and Technology (KAIST), 291 Daehak-ro, Yousung-gu, Daejeon 34141, Republic of Korea ‡

Center for Energy Convergence Research, Green City Technology Institute, Korea Institute of

Science and Technology (KIST), Hwarangno 14 gil 5, Seongbuk-gu, Seoul 02792, Seoul, Republic of Korea §

Energy & Environmental Division, Korea Institute of Ceramic Engineering & Technology

(KICET), 101 Soho-ro, Jinju-si, Gyeongsangnam-do 52581, Republic of Korea

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ABSTRACT

Here, we report amorphous iron phosphate with lattice water, namely FePO4·xH2O (x ~ 2.39) as a promising sodium-ion battery (SIB) cathode. After carbon coating, micron-sized FePO4·xH2O exhibits higher reversible capacity than its counterpart without lattice water (130.0 mAh g-1 vs. 50.6 mAh g-1 at 0.15C rate), along with clearly enhanced rate capability and cyclability. The superior electrochemical performance of FePO4·xH2O is attributed to the lattice water that facilitates sodium-ion diffusion via enlarged channel dimensions and the screening of the electrostatic interactions between sodium-ions and host anions. The amorphous phase is also advantageous in accommodating the stress created in the host framework during sodium-ion (de)intercalation. The presence of lattice water also protects the oxidation state of Fe from reductive surface carbon coating and slightly lowers the operation voltage via reduced inductive effect. The current study provides a useful insight into how to design SIB electrode materials particularly focusing on facile sodium-ion diffusion.


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TOC GRAPHICS

MAIN TEXT Targeting large-scale energy storage systems (ESSs), sodium ion batteries (SIBs) have been extensively investigated as viable alternatives to lithium ion batteries (LIBs).1-3 In pursuit of large-scale applications, SIBs have conspicuous advantages over LIBs such as low cost and worldwide accessibility of raw materials. However, the larger ionic diameter and heavier mass of the Na ion (1.06 Å, 22.99 g mol-1) in comparison with those of the Li ion (0.76 Å, 6.94 g mol-1) result in inferior electrochemical performance; the (de)intercalation of Na ions gives rise to severe structural distortions of the original phases that are accompanied by complicated phase transitions.4 As a result, most SIB electrode materials exhibit worse cycling and rate performance than their LIB counterparts. In an effort to overcome these limitations, a vast number of cathode materials have lately been investigated. These include various layered oxides5-14and polyanionic phases15-25 containing different transition metals (TMs). Fortunately, a long history of LIB research has allowed investigations in SIBs to take advantage of previous experience accumulated during the past few decades.


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As phosphates with the olivine framework have received great attention for LIB cathodes due to their well-defined operation voltage (~3.4 V vs. Li/Li+), decent charge-discharge kinetics, and safe characteristics,26-29 the sodium versions of phosphate have also been intensively studied. In particular, in the case of iron (Fe) as the TM, sodium iron phosphate (NaFePO4 or NFP) can hold two distinct crystal phases, namely the maricite phase and olivine phase.30 Conventional syntheses preferentially produce the maricite phase because it is thermodynamically more stable than the olivine phase.31 However, the ionic channels in the maricite phase are closed, rendering the material electrochemically inactive. In contrast, as in LIBs, olivine-NFP shows reversible (de)intercalation of Na ions. Unfortunately, scalable syntheses of olivine-NFP are not available. In fact, the production of olivine-NFP requires delicate control in chemical or electrochemical synthesis30-31 due to its metastable character. Although olivine-NFP could be prepared by ionexchange from olivine-LFP, the process is time-consuming and difficult to scale-up.30 Furthermore, the synthesis of pure olivine-LFP requires sophisticated control in multiple steps, as (1) particles can be grown in an uncontrolled manner at temperatures higher than 600 °C, (2) Fe3+-containing impurity phases, including LiFePO4OH, Li3Fe2(PO4)3, FePO4, LiFeP2O7, and Fe2O3, can be formed, and (3) undesired side-products, such as Li2CO3 and Li3PO4, can be generated on the surfaces of host particles.32 While crystalline NFP has these drawbacks, the battery community has also paid attention to amorphous FP as an alternative cathode phase for SIBs. In fact, it was found33-35 that the amorphous nature provides sufficient ionic channels for Na ion diffusion, although the channel pathways are not clearly defined. The amorphous phases carry additional advantages of low cost and easy synthesis. Nevertheless, electronic/ionic conductivity of amorphous FP needs to be improved to offer practically meaningful electrochemical performance. Thus, most studies have


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relied on nano-sizing, morphology control, and composite formation with carbon nanomaterials.34, 36-38 In the present investigation, we have found that as-prepared amorphous FP usually exists in a hydrated form, which exhibits far better electrochemical performance after carbon surface coating than the non-hydrated counterpart. We reveal that the origin of the improved performance lies in the lattice water in the FP framework that facilitates Na ion diffusion via increased channel dimensions and charge screenings against the lattice anions.39-44 In previous studies34,

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dealing with amorphous FP, the materials were consistently dehydrated after

calcination. Although Masquelier et al. noticed the beneficial role of lattice water for amorphous FP,45-46 it was considered only for LIBs. In order to elucidate the role of lattice water, both hydrated and dehydrated FP was prepared as illustrated in Figure 1. For the hydrated form, commercially available FePO4·xH2O (x ~ 2.72) was carbon-coated by a ball-milling process (denoted as c-FePO4·xH2O with x ~ 2.39). The size distributions of active particles were obtained using electrophoretic light scattering technique, and the mean diameters of both samples were similar near 450 nm (Figure S1, Supporting Information). For the dehydrated control sample, the same initial FP powder was dehydrated by an annealing step at 300 °C. It should be noted that the heat-treatment stayed at 300 °C for only 5 min to prevent particle agglomeration. The heat-treatment was also conducted under argon flow to prevent the oxidation of the FP particles. After dehydration, this control sample was carbon-coated by the same ball-milling process. The presence of lattice water in FePO4·xH2O was qualitatively verified by Fourier transform infrared (FT-IR) spectroscopy spectra that exhibited bending and stretching vibrations


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of –OH group at 1650 cm-1 and near 3400 cm-1, respectively (Figure S2, Supporting Information). By contrast, both vibration peaks diminished in the spectrum of the dehydrated counterpart. Scanning electron microscope (SEM) characterization informs the particle size of each FP analogue, as displayed in Figure 2a-d. The individual particle sizes of FePO4·xH2O and FePO4 are both 50~100 nm. After carbon-coating, the particle sizes of both cases increased to 300~500 nm but remained similar to each other. The increase in particle size is attributed to agglomeration between active particles as well as between active particles and super-P during the ball-milling step. Overall, the dehydration step had a negligible effect on the particle size, so the comparison between c-FePO4·xH2O and c-FePO4 must be able to elucidate the effect of the lattice water. The X-ray diffraction (XRD) patterns (Figure S3, Supporting Information) of the carboncoated FP analogues confirm that they preserved their amorphous character throughout the dehydration and carbon-coating processes. The amount of lattice water and carbon in each FP analogue was analyzed by measuring the weight loss during thermogravimetric analysis (TGA) under air condition in the temperature ranges of 120~380 °C and 380~630 °C, respectively (Figure 2e and Figure S4, Supporting Information).47 At above 630 °C, carbon decomposition is negligible and the decomposition of FP begins.48 From these measurements, it was found that cFePO4·xH2O and FePO4·xH2O contain 17.8 wt% and 24.5 wt% lattice water, respectively, corresponding to 2.39 and 2.72 in stoichiometry with respect to each FP formula unit. The lower water content of c-FePO4·xH2O is ascribed to partial evaporation of lattice water during the carbon-coating step. On the other hand, both c-FePO4·xH2O and c-FePO4 turned out to contain around 19 wt% of carbon. It was also confirmed that both of the dehydrated compounds, cFePO4 and FePO4, do not have any lattice water at all.


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Since the electronic conductivity is critical for electrochemical performance, only the carbon-coated samples, c-FePO4·xH2O and c-FePO4 were examined as SIB cathode materials in our battery tests (Figure 3). To this end, coin-type half-cells were prepared, in which thin disks of sodium metal were used as both the reference and counter electrodes and 1M sodium perchlorate (NaClO4) in ethylene carbonate (EC) and propylene carbonate (PC) with a 1:1 volume ratio was used as the electrolyte. As expected, the materials without carbon-coating (FePO4·xH2O and FePO4) displayed poor electrochemical performance due to the insulating character of these materials (Figure S5, Supporting Information). In all of the electrochemical measurements, the samples were first sodiated, as the pristine forms do not contain Na. Figure 3a displays the first and second sodiation/desodiation profiles of c-FePO4·xH2O and c-FePO4. The effect of lattice water was clearly revealed by their specific capacities. At a current rate of 20 mA g-1, in the first cycle, c-FePO4·xH2O exhibited sodiation/desodiation capacities of 130.0/129.1mAh g-1 with Coulombic efficiency (CE), defined as desodiation capacity/sodiation capacity, of 99.3%. All specific capacities hereafter are based on the mass of the active material including lattice water but excluding the carbon coating, to clarify the degree of FP involved in the reaction. By contrast, at the same condition, c-FePO4 delivered only 50.6/49.1 mAh g-1 with 97.0% of CE. The reversible capacity of c-FePO4·xH2O is quite remarkable, as this value corresponds to 94.1% of its theoretical value (138.2 mAh g-1). Even though the presence of the lattice water sacrifices the specific capacity by increasing the molecular weight, the observed reversible capacity (130.0 mAh g-1) of the micron-sized particles is comparable with those of nanostructured FP in the composite forms with carbon nanotubes (120 mAh g-1),34 mesoporous carbon (151 mAh g-1),36 and reduced graphene oxide (130.5 mAh g-1).38 For reference, in contrast with the value (94.1%) of c-FePO4·xH2O in our case, these FP composites delivered


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only 67.8%, 85.0%, and 73.4% with respect to the theoretical capacities of FP, respectively. Apparently, the evaluation of the specific capacity in comparison with the theoretical value depends on the content of lattice water in formula unit and could thus vary to some degree, depending on the sample preparation. However, the higher activation of c-FePO4·xH2O compared to c-FePO4 and therefore the crucial role of lattice water remains unchanged. The dQ/dV profiles (Figure 3b) also reflect the presence of lattice water in c-FePO4·xH2O. c-FePO4 exhibited sodiation/desodiation peaks at 2.07 V/2.75 V and 2.18 V/2.73 V in the first and second cycle. By contrast, c-FePO4·xH2O displayed sodiation/desodiation peaks at lower potentials of 2.00 V/2.58 V and 2.08 V/2.63 V, respectively, in the first and second cycle (Table S1, Supporting Information). The lower redox potentials of c-FePO4·xH2O can be explained by the decreased inductive effect due to the presence of lattice water, details of which will be discussed later in Figure 6. c-FePO4·xH2O exhibited stable cycling performance (Figure 3c) when tested at different current densities with various voltage windows. For example, at 20 mA g-1 with 1.5 V~4.15 V, 94.4% of the reversible capacity in the first cycle (122.4 mAh g-1) was retained after 100 cycles. At 10 times higher current density of 200 mA g-1 with the same voltage window, the reversible capacity in the first cycle dropped to 70.7 mAh g-1. However, the capacity retention was improved to 119.9% after 100 cycles. The increased capacity with cycling is attributed to a certain gradual activation process related to ion and electron migration38,

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that takes place

during a period of cycles at high current densities. Deeper understanding requires further investigation. Also, at a fixed current density of 20 mA g-1 but with a more narrow voltage window of 2.0 V~3.5 V, the capacity retention improved to 99.1% after 100 cycles, with respect to its reversible capacity of 65.7 mAh g-1 in the first cycle. This cyclability of c-FePO4·xH2O is as


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good as that of previously reported nanostructured FP composites34,

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with costly carbon

nanomaterials. This result suggests that the intrinsically low electronic and ionic conductivities of amorphous FP can be addressed by ball-mill based carbon coating and lattice water even for micron-sized particles. The cyclability of c-FePO4·xH2O can also be interpreted in terms of stress alleviation related to bond breaking/formation. While Na ion insertion/desertion during the electrochemical process is associated with breaking and forming multiple bonds that gives rise to substantial stress arising from reactions with large-sized Na ions, c-FePO4·xH2O effectively releases that stress by utilizing its amorphous structure and lattice water. The superior Na ion diffusion kinetics by the presence of lattice water was unveiled more directly in the rate performance test (Figure 3d). In the case of c-FePO4·xH2O, when the current rate gradually increased from 10 mA g-1 to 1000 mA g-1, the specific reversible capacity decreased from 135.4 to 31.4 mAh g-1. However, upon returning to a current density of 20 mA g1

, the original capacity at the same current density was fully recovered. In more aggressive

testing conditions that engage abrupt switching between high (1200 and 1500 mA g-1) and low (20 mA g-1) current densities, 101.3% and 101.5% of the original capacities at 20 mA g-1 were recovered. An additional round of a series of current densities was applied to observe rate performance during prolonged cycles. During the entire 130 cycles, the specific capacities at high current densities decreased in the second round as compared with those in the first round. However, the specific capacity at 20 mA g-1 in the cycles 128~130 was 95.8% of the initial value at the same current density. On the contrary, c-FePO4 exhibited much worse rate performance; although it started with smaller capacities at lower current densities (i.e., 58.9 mAh g-1 at 10 mA g-1), it ceased to operate in the middle of abrupt switching between low and high current densities in the first round. Since then, c-FePO4 stayed inactive, implying that the poorly defined


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ionic channels do not sustain repeated sodiation/desodiation cycles, especially the aggressive cycles consisting of high and low current density swing. More detailed reasoning is difficult to extract due to the amorphous nature of the material. The superior performance of c-FePO4·xH2O can mainly be explained by the lattice water. When cathode materials are sodiated, in most cases, Na ions are first desolvated at the electrodeelectrolyte interface, and bare Na ions then diffuse along the channels in the host structure. In this situation, the presence of the lattice water facilitates Na ion diffusion in the framework via charge screening between Na ions and the host anions as well as increased channel dimensions,50-51 although channel dimensions are difficult to quantify using diffraction analysis due to the amorphous character. The lattice water could also serve as pillars in the channels and contribute to structural stability over cycling,52-53 which is consistent with the recent report by Masquelier et al.45 that pointed out the beneficial role of the lattice water in hydrated FP for LIB performance. Besides the enlarged channel size, the amorphous character of c-FePO4·xH2O may be useful for facile Na ion diffusion during repeated cycles. It was reported6,54 that many crystalline SIB cathodes suffer from large strain and multiple phase transitions during sodiation/desodiation, which are associated with the large size of the Na ion. While these are intrinsic problems of SIBs, amorphous structures can alleviate the drawbacks by taking advantage of their relatively disordered atomic arrangements, which is more suitable for accommodating the strains. Thus, the amorphous phase and lattice water have a synergistic effect on the observed cycling and rate performance of c-FePO4·xH2O. It was confirmed that the lattice water stayed stable during cycling. The stability of lattice water is a crucial factor because the release of lattice water to some degree into the electrolyte would most likely harm the overall cell performance. When water content was compared before


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and after 10 cycles at a current rate of 20 mA g-1 using TGA, the content of lattice water was found to be consistent (both 16.8 wt%) as displayed in Figure S6, Supporting Information. The TGA curves of the electrode samples appeared a bit different from those of the pristine powder (Figure 2) because of the presence of other electrode components. In order to monitor the oxidation states of both the carbon-coated FP electrodes, X-ray absorption near edge structure (XANES) analyses were carried out. Figure 4a and b present the Fe-K edge XANES spectra of FePO4·xH2O and FePO4 before and after the carbon-coating process, along with those of FeO and Fe2O3 that serve as references for Fe2+ and Fe3+, respectively. The spectra of both pristine FePO4·xH2O and FePO4 nearly overlap with those of Fe2O3, indicating that the oxidation states of Fe in both compounds are close to +3. In addition, the spectrum of the carbon-coated version, c-FePO4·xH2O, remained close to that of Fe2O3, indicating the oxidation state of Fe in c-FePO4·xH2O was well preserved after the carbon-coating process, close to +3. By contrast, the spectrum of c-FePO4 slightly shifted toward that of FeO after carbon coating, reflecting a partial reduction of Fe toward +2 state. Although both samples underwent the same carbon-coating process, this oxidation state difference implies that the presence of the lattice water protects Fe from well-known carbo-reduction;55 when lattice water is present in the enlarged lattice spacing, the electron transfer for the reduction of Fe becomes mitigated to some level compared to the other case without lattice water. Ex-situ XANES analyses of both c-FePO4·xH2O and c-FePO4 were conducted to investigate the redox behavior of those materials during the first and second cycles (Figures 5 and S7, Supporting Information). Both samples exhibited profile-shifts toward lower and higher energy values after sodiation and desodiation, respectively, indicating reduction and re-oxidation of Fe as well as the reversibility of the given redox reaction. The pre-edge regions (Insets of


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Figures 5 and S7, Supporting Information) of both samples that are reflective of the transition of the core 1s electron to an unoccupied 3d orbital followed the same trend as the edge regions, reconfirming the corresponding reduction/oxidation of Fe. However, the swing widths of the energy absorption edge between sodiation and desodiation states were more narrow for c-FePO4 than that for c-FePO4·xH2O, which is consistent with the lower specific capacities of c-FePO4 in the electrochemical characterization. In particular, the XANES results indicate that the more narrow swing widths of c-FePO4 are largely caused by the original spectrum position of c-FePO4 at the pristine state, which corresponds to the partial reduction of Fe due to the carbo-reductive environment during carbon coating, as addressed above. The presence of lattice water imposes a different level of the inductive effect toward the TM center and modifies the operation voltage as verified in Figure 3b. Manthiram et al. originally rationalized56 the inductive effect from the iono-covalent nature of the bond between TM center and coordinating polyanions; depending on the electronegativity difference between the TM and the polyanion, the bond character can be varied from ionic to covalent.57 For more ionic the bond is, the weaker the hybridization between 3d orbital of the TM and 2p orbital of the coordinating atom; consequently, the bonding-antibonding (σ- σ*) splitting in terms of energy level is smaller (the right-hand case in Figure 6a). As a result, the operation voltage that is determined by the distance from the Na/Na+ redox energy level to the antibonding energy level of the hybridized orbital becomes higher compared to the more covalent bond case. Taking this into consideration, the presence of lattice water weakens the inductive effect: when the lattice water is coordinated to the FP host via hydrogen bonding with the oxygen atoms in the PO4 tetrahedra, it is anticipated that the hydrogen bonding weakens the ionic character of the Fe—O (in PO4) bond by lowering the electronegativity of the O in PO4, leading to the lower operating


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voltage of c-FePO4·xH2O compared to that of c-FePO4. The slightly shorter bond length of Fe— O of c-FePO4·xH2O was revealed by extended X-ray absorption fine structure (EXAFS) analysis (Figure S8, Supporting Information). In conclusion, the combination of lattice water and amorphous structure significantly improves the electrochemical performance of FP, as verified by comparative analyses with the control case without lattice water. The presence of the lattice water facilitates Na ion diffusion in the framework and stabilizes the overall structure for prolonged cycling, leading to superior cycling and rate performance. The enhanced performance based on lattice water provides a useful insight into how to design SIB electrode materials and suggests that various hydrated materials with enlarged Na ion channels are worth exploring further.

ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge on the ACS Publication website at DOI: . Experimental description, particle size distributions, FT-IR spectra, XRD spectra, TG profiles, electrochemical results, and XANES profiles (PDF)

AUTHOR INFORMATION Corresponding Author E-mail: [email protected]


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Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT We acknowledge the financial support from Korea Institute of Science and Technology (KIST) Institutional Program (2E27090), Industrial Strategic technology development program, (Project No. 10050477, Development of separator with low thermal shrinkage and electrolyte with high ionic conductivity for Na-ion batteries) funded By the Ministry of Trade, industry & Energy (MI, Korea), the Energy Efficiency & Resources Core Technology Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP), which is granted financial resources from the Ministry of Trade, Industry & Energy, Republic of Korea (20152020104870), Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (NRF-2015R1D1A1A01058334), and the Project of flexible power module and system development for wearable devices funded by the Ministry of Trade, Industry and Energy (10065730). S. Y. L. Acknowledges the National Research Foundation of Korea for the NFR-2013-Global PhD Fellowship Program. The authors thank the Pohang Accelerator Laboratory, Korea for extending the synchrotron XRD beamline for characterization.

References. (1) Choi, J. W.; Aurbach, D., Promise and Reality of Post-Lithium-Ion Batteries with High Energy Densities. Nat. Rev. Mater. 2016, 1, 16013. (2) Ellis, B. L.; Nazar, L. F., Sodium and Sodium-Ion Energy Storage Batteries. Curr. Opin. Solid State Mater. Sci. 2012, 16, 168-177.


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(3) Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S., Research Development on SodiumIon Batteries. Chem. Rev. 2014, 114, 11636-11682. (4) Slater, M. D.; Kim, D.; Lee, E.; Johnson, C. S., Sodium‐Ion Batteries. Adv. Funct. Mater. 2013, 23, 947-958. (5) Berthelot, R.; Carlier, D.; Delmas, C., Electrochemical Investigation of the P2-NaxCoO2 Phase Diagram. Nat. Mater. 2011, 10, 74-80. (6) Delmas, C.; Braconnier, J. J.; Fouassier, C.; Hagenmuller, P., Electrochemical Intercalation of Sodium in NaxCoO2 Bronzes. Solid State Ionics 1981, 3-4, 165-169. (7) Yabuuchi, N.; Kajiyama, M.; Iwatate, J.; Nishikawa, H.; Hitomi, S.; Okuyama, R.; Usui, R.; Yamada, Y.; Komaba, S., P2-type NaxFe1/2Mn1/2O2 Made from Earth-Abundant Elements for Rechargeable Na Batteries. Nat. Mater. 2012, 11, 512-517. (8) Pan, H.; Lu, X.; Yu, X.; Hu, Y.-S.; Li, H.; Yang, X.-Q.; Chen, L., Sodium Storage and Transport Properties in Layered Na2Ti3O7 for Room-Temperature Sodium-Ion Batteries. Adv. Energy Mater. 2013, 3, 1186-1194. (9) Vassilaras, P.; Toumar, A. J.; Ceder, G., Electrochemical Properties of NaNi1/3Co1/3Fe1/3O2 as a Cathode Material for Na-ion Batteries. Electrochem. Commun. 2014, 38, 79-81. (10) Yoshida, H.; Yabuuchi, N.; Komaba, S., NaFe0.5Co0.5O2 as High Energy and Power Positive Electrode for Na-ion Batteries. Electrochem. Commun. 2013, 34, 60-63. (11) Chen, X.; Zhou, X.; Hu, M.; Liang, J.; Wu, D.; Wei, J.; Zhou, Z., Stable Layered P3/P2 Na0.66Co0.5Mn0.5O2 Cathode Materials for Sodium-ion Batteries. J. Mater. Chem. A 2015, 3, 20708-20714. (12) Bucher, N.; Hartung, S.; Franklin, J. B.; Wise, A. M.; Lim, L. Y.; Chen, H.-Y.; Weker, J. N.; Toney, M. F.; Srinivasan, M., P2–NaxCoyMn1–yO2 (y = 0, 0.1) as Cathode Materials in Sodium-Ion Batteries—Effects of Doping and Morphology To Enhance Cycling Stability. Chem. Mater. 2016, 28, 2041-2051. (13) Yuan, D.; He, W.; Pei, F.; Wu, F.; Wu, Y.; Qian, J.; Cao, Y.; Ai, X.; Yang, H., Synthesis and Electrochemical Behaviors of Layered Na0.67[Mn0.65Co0.2Ni0.15]O2 Microflakes as a Stable Cathode Material for Sodium-ion Batteries. J. Mater. Chem. A 2013, 1, 3895-3899. (14) You, Y.; Kim, S. O.; Manthiram, A., A Honeycomb-Layered Oxide Cathode for SodiumIon Batteries with Suppressed P3–O1 Phase Transition. Adv. Energy Mater. 2016, 7, 1601698. (15) Ellis, B. L.; Makahnouk, W. R. M.; Makimura, Y.; Toghill, K.; Nazar, L. F., A Multifunctional 3.5 V Iron-Based Phosphate Cathode for Rechargeable Batteries. Nat. Mater. 2007, 6, 749-753. (16) Barpanda, P.; Nishimura, S.-i.; Yamada, A., High-Voltage Pyrophosphate Cathodes. Adv. Energy Mater. 2012, 2, 841-859. (17) Barpanda, P.; Oyama, G.; Nishimura, S.-i.; Chung, S.-C.; Yamada, A., A 3.8-V EarthAbundant Sodium Battery Electrode. Nat. Commun. 2014, 5, 4358. (18) Lim, S. Y.; Kim, H.; Chung, J.; Lee, J. H.; Kim, B. G.; Choi, J.-J.; Chung, K. Y.; Cho, W.; Kim, S.-J.; Goddard, W. A., Role of Intermediate Phase for Stable Cycling of Na7V4(P2O7)4PO4 in Sodium Ion Battery. Proc. Natl. Acad. Sci. USA 2014, 111, 599-604. (19) Kim, H.; Park, I.; Seo, D.-H.; Lee, S.; Kim, S.-W.; Kwon, W. J.; Park, Y.-U.; Kim, C. S.; Jeon, S.; Kang, K., New Iron-Based Mixed-Polyanion Cathodes for Lithium and Sodium Rechargeable Batteries: Combined First Principles Calculations and Experimental Study. J. Am. Chem. Soc. 2012, 134, 10369-10372.


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(20) Rui, X.; Sun, W.; Wu, C.; Yu, Y.; Yan, Q., An Advanced Sodium-Ion Battery Composed of Carbon Coated Na3V2(PO4)3 in a Porous Graphene Network. Adv. Mater. 2015, 27, 66706676. (21) Kim, H.; Park, C. S.; Choi, J. W.; Jung, Y., Defect-Controlled Formation of Triclinic Na2CoP2O7 for 4 V Sodium-Ion Batteries. Angew. Chem. Int. ed. 2016, 55, 6662-6666. (22) Kim, H.; Shakoor, R. A.; Park, C.; Lim, S. Y.; Kim, J.-S.; Jo, Y. N.; Cho, W.; Miyasaka, K.; Kahraman, R.; Jung, Y.; Choi, J. W., Na2FeP2O7 as a Promising Iron-Based Pyrophosphate Cathode for Sodium Rechargeable Batteries: A Combined Experimental and Theoretical Study. Adv. Funct. Mater. 2013, 23, 1147-1155. (23) Kim, H.; Yoon, G.; Park, I.; Hong, J.; Park, K.-Y.; Kim, J.; Lee, K.-S.; Sung, N.-E.; Lee, S.; Kang, K., Highly Stable Iron- and Manganese-Based Cathodes for Long-Lasting Sodium Rechargeable Batteries. Chem. Mater. 2016, 28, 7241-7249. (24) Barker, J.; Saidi, M. Y.; Swoyer, J. L., A Sodium-ion Cell Based on the Fluorophosphate Compound NaVPO4F. Electrochem. Solid State Lett. 2003, 6, A1-A4. (25) Kumar, P. R.; Jung, Y. H.; Kim, D. K., Influence of Carbon Polymorphism Towards Improved Sodium Storage Properties of Na3V2O2x (PO4)2F3-2x. J. Solid State Electrochem. 2017, 21, 223-232. (26) Padhi, A. K.; Nanjundaswamy, K.; Goodenough, J., Phospho‐Olivines as Positive‐ Electrode Materials for Rechargeable Lithium Batteries. J. Electrochem. Soc. 1997, 144, 11881194. (27) Yamada, A.; Chung, S.-C.; Hinokuma, K., Optimized LiFePO4 for Lithium Battery Cathodes. J. Electrochem. Soc. 2001, 148, A224-A229. (28) Yuan, L.-X.; Wang, Z.-H.; Zhang, W.-X.; Hu, X.-L.; Chen, J.-T.; Huang, Y.-H.; Goodenough, J. B., Development and Challenges of LiFePO4 Cathode Material for Lithium-Ion Batteries. Energy Environ. Sci. 2011, 4, 269-284. (29) Peng, L.; Zhao, Y.; Ding, Y.; Yu, G., Self-assembled LiFePO4 Nanowires with High Rate Capability for Li-ion Batteries. Chem. Commum. 2014, 50, 9569-9572. (30) Moreau, P.; Guyomard, D.; Gaubicher, J.; Boucher, F., Structure and Stability of Sodium Intercalated Phases in Olivine FePO4. Chem. Mater. 2010, 22, 4126-4128. (31) Lee, K. T.; Ramesh, T. N.; Nan, F.; Botton, G.; Nazar, L. F., Topochemical Synthesis of Sodium Metal Phosphate Olivines for Sodium-Ion Batteries. Chem. Mater. 2011, 23, 3593-3600. (32) Yamada, A.; Chung, S. C.; Hinokuma, K., Optimized LiFePO4 for Lithium Battery Cathodes. J. Electrochem. Soc. 2001, 148, A224-A229. (33) Xiong, H.; Slater, M. D.; Balasubramanian, M.; Johnson, C. S.; Rajh, T., Amorphous TiO2 Nanotube Anode for Rechargeable Sodium Ion Batteries. J. Phy. Chem. Lett. 2011, 2, 2560-2565. (34) Liu, Y.; Xu, Y.; Han, X.; Pellegrinelli, C.; Zhu, Y.; Zhu, H.; Wan, J.; Chung, A. C.; Vaaland, O.; Wang, C.; Hu, L., Porous Amorphous FePO4 Nanoparticles Connected by SingleWall Carbon Nanotubes for Sodium Ion Battery Cathodes. Nano Lett. 2012, 12, 5664-5668. (35) Denis, S.; Baudrin, E.; Touboul, M.; Tarascon, J. M., Synthesis and Electrochemical Properties of Amorphous Vanadates of General Formula RVO4 (R= In, Cr, Fe, Al, Y) vs. Li. J. Electrochem. Soc. 1997, 144, 4099-4109. (36) Fang, Y.; Xiao, L.; Qian, J.; Ai, X.; Yang, H.; Cao, Y., Mesoporous Amorphous FePO4 Nanospheres as High-Performance Cathode Material for Sodium-Ion Batteries. Nano Lett. 2014, 14, 3539-3543.


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(37) Xu, S.; Zhang, S.; Zhang, J.; Tan, T.; Liu, Y., A Maize-Like FePO4@MCNT Nanowire Composite for Sodium-Ion Batteries via a Microemulsion Technique. J. Mater. Chem. A 2014, 2, 7221-7228. (38) Liu, Y.; Xu, S.; Zhang, S.; Zhang, J.; Fan, J.; Zhou, Y., Direct Growth of FePO4/Reduced Graphene Oxide Nanosheet Composites for the Sodium-Ion Battery. J. Mater. Chem. A 2015, 3, 5501-5508. (39) Nam, K. W.; Kim, S.; Lee, S.; Salama, M.; Shterenberg, I.; Gofer, Y.; Kim, J.-S.; Yang, E.; Park, C. S.; Kim, J.-S.; Lee, S.-S.; Chang, W.-S.; Doo, S.-G.; Jo, Y. N.; Jung, Y.; Aurbach, D.; Choi, J. W., The High Performance of Crystal Water Containing Manganese Birnessite Cathodes for Magnesium Batteries. Nano Lett. 2015, 15, 4071-4079. (40) Sai Gautam, G.; Canepa, P.; Richards, W. D.; Malik, R.; Ceder, G., Role of Structural H2O in Intercalation Electrodes: The Case of Mg in Nanocrystalline Xerogel-V2O5. Nano Lett. 2016, 16, 2426-2431. (41) Sun, X. Q.; Duffort, V.; Mehdi, B. L.; Browning, N. D.; Nazar, L. F., Investigation of the Mechanism of Mg Insertion in Birnessite in Nonaqueous and Aqueous Rechargeable Mg-Ion Batteries. Chem. Mater. 2016, 28, 534-542. (42) Tepavcevic, S.; Liu, Y. Z.; Zhou, D. H.; Lai, B.; Maser, J.; Zuo, X. B.; Chan, H.; Kral, P.; Johnson, C. S.; Stamenkovic, V.; Markovic, N. M.; Rajh, T., Nanostructured Layered Cathode for Rechargeable Mg-Ion Batteries. Acs Nano 2015, 9, 8194-8205. (43) Senguttuvan, P.; Han, S.-D.; Kim, S.; Lipson, A. L.; Tepavcevic, S.; Fister, T. T.; Bloom, I. D.; Burrell, A. K.; Johnson, C. S., A High Power Rechargeable Nonaqueous Multivalent Zn/V2O5 Battery. Adv. Energy Mater. 2016, 6, 1600826. (44) Tepavcevic, S.; Xiong, H.; Stamenkovic, V. R.; Zuo, X. B.; Balasubramanian, M.; Prakapenka, V. B.; Johnson, C. S.; Rajh, T., Nanostructured Bilayered Vanadium Oxide Electrodes for Rechargeable Sodium-Ion Batteries. Acs Nano 2012, 6, 530-538. (45) Masquelier, C.; Reale, P.; Wurm, C.; Morcrette, M.; Dupont, L.; Larcher, D., Hydrated Iron Phosphates FePO4.nH2O and Fe4(P2O7)3.nH2O as 3 V Positive Electrodes in Rechargeable Lithium Batteries. J. Electrochem. Soc. 2002, 149, A1037-A1044. (46) Delacourt, C.; Poizot, P.; Bonnin, D.; Masquelier, C., Lithium-Insertion Mechanism in Crystalline and Amorphous FePO4⋅ nH2O. J. Electrochem. Soc. 2009, 156, A595-A605. (47) Hong, Y. S.; Ryu, K. S.; Park, Y. J.; Kim, M. G.; Lee, J. M.; Chang, S. H., Amorphous FePO4 as 3 V Cathode Material for Lithium Secondary Batteries. J. Mater. Chem. 2002, 12, 1870-1874. (48) Liu, H.-C.; Ho, W.-H.; Li, C.-F.; Yen, S.-K., Electrochemical Synthesis of FePO4 for Anodes in Rechargeable Lithium Batteries. J. Electrochem. Soc. 2008, 155, E178-E182. (49) Mathew, V.; Kim, S.; Kang, J.; Gim, J.; Song, J.; Baboo, J. P.; Park, W.; Ahn, D.; Han, J.; Gu, L.; Wang, Y.; Hu, Y.-S.; Sun, Y.-K.; Kim, J., Amorphous Iron Phosphate: Potential Host for Various Charge Carrier Ions. NPG Asia Mater. 2014, 6, e138. (50) Novak, P.; Imhof, R.; Haas, O., Magnesium Insertion Electrodes for Rechargeable Nonaqueous Batteries - a Competitive Alternative to Lithium? Electrochim. Acta 1999, 45, 351367. (51) Novak, P.; Desilvestro, J., Electrochemical Insertion of Magnesium in Metal-Oxide and Sulfides from Aprotic Electrolytes. J. Electrochem. Soc. 1993, 140, 140-144. (52) Nam, K. W.; Kim, S.; Yang, E.; Jung, Y.; Levi, E.; Aurbach, D.; Choi, J. W., Critical Role of Crystal Water for a Layered Cathode Material in Sodium Ion Batteries. Chem. Mater. 2015, 27, 3721-3725.


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(53) Kim, S.; Lee, S.; Nam, K. W.; Shin, J.; Lim, S. Y.; Cho, W.; Suzuki, K.; Oshima, Y.; Hirayama, M.; Kanno, R.; Choi, J. W., On the Mechanism of Crystal Water Insertion during Anomalous Spinel-to-Birnessite Phase Transition. Chem. Mater. 2016, 28, 5488-5494. (54) Galceran, M.; Saurel, D.; Acebedo, B.; Roddatis, V. V.; Martin, E.; Rojo, T.; CasasCabanas, M., The Mechanism of NaFePO4 (De)Sodiation Determined by In Situ X-ray Diffraction. Phys. Chem. Chem. Phys. 2014, 16, 8837-8842. (55) Coustan, L.; Favier, F., Microwave-Assisted Decoration of Carbon Substrates for Manganese Dioxide-Based Supercapacitors. J. Electrochem. Soc. 2015, 162, A5133-A5139. (56) Manthiram, A.; Goodenough, J., Lithium Insertion into Fe2(SO4)3 Frameworks. J. Power Sources 1989, 26, 403-408. (57) Melot, B. C.; Tarascon, J.-M., Design and Preparation of Materials for Advanced Electrochemical Storage. Acc. Chem. Res. 2013, 46, 1226-1238.


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Figure 1. Schematic diagram illustrating synthetic routes of amorphous c-FePO4·xH2O and cFePO4. The small circles inside particles indicate lattice water.


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Figure 2. SEM images of amorphous FePO4 analogues: (a) FePO4·xH2O, (b) c-FePO4·xH2O, (c) FePO4, and (d) c-FePO4. (e) Thermogravimetric analysis (TGA) curves of c-FePO4·xH2O and cFePO4 under air atmosphere.


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ACS Energy Letters

Figure 3. (a) Galvanostatic profiles of the first and second cycles of c-FePO4·xH2O and c-FePO4. (b) The differential capacity (dQ/dV) curves of c-FePO4·xH2O and c-FePO4 obtained from the galvanostatic profiles in the first and second cycles in (a). (c) The capacity retentions of cFePO4·xH2O at different rates and voltage windows. (d) The comparison of the rate capability between c-FePO4·xH2O and c-FePO4 at different current densities. Two rounds of current density variation were provided, and in each round, the current density was changed in the sequence of 10→20→50→100→200→400→500→1000→20→1200→20→1500→20 mA g-1.


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Figure 4. Normalized Fe K-edge XANES spectra: (a) FePO4·xH2O and (b) FePO4 before and after the carbon-coating process. FeO and Fe2O3 were used as references of Fe2+ and Fe3+ states, respectively.


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ACS Energy Letters

Figure 5. Ex-situ Fe K-edge XANES spectra of (a) c-FePO4·xH2O and (b) c-FePO4 after the first sodiation and desodiation. FeO and Fe2O3 were used as references of Fe2+ and Fe3+ states, respectively. Insets: magnified pre-edge regions of (a) and (b).


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Figure 6. (a) The energy diagram showing how the bond character between center transition metal and polyanion affects the inductive effect and consequently the operation voltage. (b) Schematic illustration on the effect of lattice water on the inductive effect in c-FePO4·xH2O.


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