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Synthesis of nano-Na3V2(PO4)2F3 cathodes with excess Na+ intercalation for enhanced capacity
Applied Materials Today 19 (2020) 100554
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Synthesis of nano-Na3 V2 (PO4 )2 F3 cathodes with excess Na+ intercalation for enhanced capacity Manhua Peng, Xiayan Wang ∗ , Guangsheng Guo ∗ Center of Excellence for Environmental Safety and Biological Effects, Beijing Key Laboratory for Green Catalysis and Separation, Department of Chemistry and Chemical Engineering, Beijing University of Technology, Beijing 100124, PR China
a r t i c l e
i n f o
Article history: Received 18 October 2019 Received in revised form 5 December 2019 Accepted 30 December 2019 Keywords: Sodium-ion battery Vanadium cathode High capacity Na+ intercalation
a b s t r a c t Vanadium-based materials are promising as cathodes for sodium-ion batteries owing to their superior cycling behavior and rate performance. However, there is a critical need for the rational implementation of Na+ intercalation to enhance the capacity. Herein, nano-Na3 V2 (PO4 )2 F3 materials were synthesized via bulk dismemberment by gas released by the decomposition of excess citric acid to trigger some degree of lattice distortion and defects. Owing to the weakened electrostatic repulsion between sodium ions in Na3 V2 (PO4 )2 F3 with an enlarged interlayer spacing, one more Na+ ion could be readily embedded to form Na4 V2 (PO4 )2 F3 . A new reaction plateau at 1.38 V/1.56 V accompanied the V3+ /V2+ redox reaction, inducing a capacity of 250 mA h·g−1 , which is higher than that of highly crystallined Na3 V2 (PO4 )2 F3 . A noteworthy reversible redox phenomenon involving three sodium ions was confirmed during the cycling process. The nano-Na3 V2 (PO4 )2 F3 electrode exhibited a capacity retention of 72% at 1.3–4.5 V after 20 cycles. Thus, this novel synthesis routine can provide vanadium-based cathode materials with enhanced Na+ intercalation properties, which can be applied to improve the capacity of sodium-ion batteries. © 2020 Elsevier Ltd. All rights reserved.
1. Introduction Owing to the abundance a low price of sodium, sodium-ion batteries (SIBs) can serve as backup power supplies for large-scale electronic devices [1–3]. However, a key issue facing SIBs is their low energy densities relative to those of lithium-ion batteries. Moreover, in crystals, it is more difficult to deintercalate sodium ions than lithium ions [4–6]. Increasing number of electrons transferred by a transition metal via the introduction of multiple valence states is an effective way to improve the battery capacity [7–9]. In this regard, vanadium, which has multiple valence states, could allow enhanced sodium ion intercalation [9]. Recently, vanadiumbased materials such as Na3 V2 O2-x (PO4 )2 F1+x(0≤x≤2) , NaVPO4 F, Na3 V2 (PO4 )3 , NaVOPO4 , and Na7 V4 (P2 O7 )4 (PO4 ) have been developed as cathode materials for SIBs [10–19]. Among these materials,
Abbreviations: SIB, sodium ion battery; SG, material material prepared by a modified sol–gel reaction; SS, material material prepared by a standard solid-state reaction; XRD, X-ray diffraction; SEM, scanning electron microscopy; TEM, transmission electron microscopy; XPS, X-ray photoelectron spectroscopy; EIS, impedance spectroscopy; GITT, galvanostatic intermittent titration technique; TGA, Thermogravimetric analysis. ∗ Corresponding authors. E-mail addresses: [email protected] (X. Wang), [email protected] (G. Guo). https://doi.org/10.1016/j.apmt.2020.100554 2352-9407/© 2020 Elsevier Ltd. All rights reserved.
pseudo-layered Na3 V2 (PO4 )2 F3 has a stable electronic structure and a higher operating voltage. However, the reported theoretical capacity of this material is only 130 mA h·g−1 because only two sodium ions can be reversibly deintercalated, corresponding to the V3+ /V4+ redox reaction at the voltage of 2.0–4.5 V [20–27]. Therefore, to improve the theoretical capacity of vanadium-based electrode materials, it is necessary to further increase the number of electrons transferred and to enhance sodium ion deintercalation. One optional method to increase the theoretical capacity of Na3 V2 (PO4 )2 F3 is to take advantage of a broader voltage window, such as 1.0–4.5 V instead of 2.0–4.5 V. But the deintercalation or embedding of excess sodium ions in Na3 V2 (PO4 )2 F3 materials can be realized by doping with Cr or Al at a broader voltage window of 1.0–4.5 V. Nevertheless, doping with such inert elements could cause the capacity to decrease. Thus, using pure-phase Na3 V2 (PO4 )2 F3 or doping with active elements may be a more favorable approach. However, it is difficult to intercalate a sodium ion in pure-phase of Na3 V2 (PO4 )2 F3 to form Na4 V2 (PO4 )2 F. Ceder and co-workers suggested that the intercalation behavior is mainly limited by the strong electrostatic repulsion between sodium ions rather than by the valence states of vanadium [9,28]. In view of the above analysis, this work focused on embedding excess sodium ions in pure-phase Na3 V2 (PO4 )2 F3 with the aim of enhancing the capacity. To achieve this goal, we synthesized nano-vanadium materials with increasing the interlayer spacing
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Fig. 1. Schematic of the modified sol–gel process for preparing nano-Na3 V2 (PO4 )2 F3 .
Fig. 2. (a) XRD pattern of the SS and SG materials. (b, c) SEM and (e, f) TEM images of the SS and SG materials (insets: images at higher magnification). (d) Unit cell of Na3 V2 (PO4 )2 F3 and the corresponding lattice spacings.
via a modified sol–gel method, which involved bulk dismemberment by gas originating from the decomposition of excess citric acid in the precursor. It is expected that the intercalation of three sodium ions could be achieved by effectively weakening the electrostatic repulsion in Na3 V2 (PO4 )2 F3 by increasing the interlayer spacing. The intercalation mechanism was systematically investigated by microstructural characterization and electrochemical measurements.
2. Materials and methods
2.2. Characterization The phases of the synthesized materials were characterized by X-ray diffraction (XRD) D8 Advance DaVinci, Bruker, Germany) equipped with a Cu target ( = 1.5418 Å) in the scanning range of 2 = 10–90 ◦ C. Scanning electron microscopy (SEM; Hitachi S4300, Japan) and transmission electron microscopy (TEM; Tecnai G2 F20, FEI, USA.) were utilized to analyze the morphologies of the materials. X-ray photoelectron spectroscopy (XPS, ThermoFischer, ESCALAB 250Xi) was used to determine the chemical environment of the transition metal at different cutoff voltages. Thermogravimetric analysis (TGA, France setaram company, LABSYS EVO) were utilized to analyze thermal decomposition of materials.
2.1. Material synthesis 2.3. Electrochemical performance Nano-Na3 V2 (PO4 )2 F3 was synthesized by a modified sol–gel reaction (denoted as SG material). The reagents C6 H8 O7 , H3 PO4 , NH4 VO3 , NaF, and NH4 F (stoichiometric ratio of 1.5:1.0:1.0:1.7:0.6) were dissolved in deionized water to form a homogeneous solution. Then, the solution was heated at 85 ◦ C for 6 h until completely dried. The obtained powder was placed in a tube furnace and heated at 300 ◦ C for 2 h in an argon atmosphere. The temperature was then increased to 600–700 ◦ C at a high speed of 30 ◦ C min−1 and the material was heated for a further 30 min. For comparison, highly crystallined Na3 V2 (PO4 )2 F3 was prepared by a standard solid-state reaction (denoted as SS material) according to a literature procedure [29].
To test the electrochemical performance of the synthesized electrode materials, the electrode material, acetylene black, and a binder of PTFE were mixed in a ratio of 70:20:10 to make an electrode plate. After the electrode plate was dried in an oven at 100 ◦ C for 10 h, it was cut into small wafers with masses of approximately 3–5 mg. Batteries were assembled in an argon-filled glove box using an electrolyte of 1 mol·L−1 NaClO4 in propylene carbonate with 2% fluoroethylene carbonate, a glass fiber film (grade GF/D, Whatman, USA), and a pure sodium plate as the negative electrode material. Electrochemical testing was performed at room temperature using a NEWARE battery test system (China). GITT was executed
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Fig. 3. (a, b) First cycle and second cycles of SS and SG materials in a voltage window of 1.2–4.5 V (1 C corresponds to 130 mA h·g−1 ). (c) Differential curve of dQ/dV. (d) Cyclic voltammograms. (e) Cycling stability. (f) Charge–discharge curves.
by NEWARE system. The battery was first charged for 5 min and then rest for 0.5 h. The whole charging and discharging process was performed at a current rate of 0.1 C. Electrochemical impedance spectroscopy (EIS) was performed on a ZAHNER system (Germany) using frequencies from 1 mHz to 1 MHz at an AC voltage of 5 mV.
3. Results and discussion The structural characteristics of the SS and SG materials were thoroughly investigated. To obtain a loose and expanded Na3 V2 (PO4 )2 F3 structure, an increased amount of chelating agent (citric acid) was added during gel formation (Fig. 1). The decomposition of citric acid could release a gas (carbon dioxide, water) that then dismembered the bulk into small pieces. This process is analogous to the “steamed bread” process, in which yeast works like citric acid. Thermogravimetric analysis confirmed the occurrence of this phenomenon because the weight loss reached 30% between 150 and 300 ◦ C, which mainly derived from the decomposition of citric acid (Fig. S1). Moreover, an additional weight loss of 15% was observed from 300 to 500 ◦ C, probably owing to further carbonization of citric acid and crystal water (Fig. S2). Because of these decomposition processes, the prepared SG material is fluffy,
which is different from the powdered SS material formed by the standard solid-state reaction (Fig. S3). The XRD patterns (Fig. 2a) confirm that both pure-phase SG and SS materials were synthesized. However, the characteristic peaks corresponding to the (002), (212) (113), and (202) facets of the SG material were broader than those of the SS material. Furthermore, the XRD peaks of the SG material were slightly shifted toward lower angles, which confirms that the lattice spacing was enlarged. (Fig. 2a, Fig. S4). SEM and TEM images (Fig. 2b, c and e, f) show that the sizes of the SS and spherical SG materials were approximately 1 m and 100 nm in diameter, respectively. Furthermore, after the heat treatment during the sol–gel process, many defects were formed on the surface of the spherical SG particles (Fig. 2f), whereas the SS material has a relative well-defined crystal plane (Fig. 2e). Typically, sodium ions are located in the interlayer of the (002) plane (Fig. 2d), where they migrate mainly along the xy axis. The SG material had a wider (002) lattice spacing (d = 0.556 nm) than that (d = 0.521 nm) of the SS material (Fig. S5). The intercalation behavior is limited mainly by the strong electrostatic repulsion between sodium ions affected by the compactness of sodium ions arrangement. The obtained a structure of materials including the lattice distortion, lattice defects and wider (002) lattice distance was confirmed by
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Es 2
Fig. 4. (a) XRD patterns a SG material electrode in the initial state, when charged to 4.5 V, and when discharged to 1.3 and 1.2 V. (b) Magnification of the XRD patterns in the region shown in the black box in (a). (c) Unit cell parameters. (d) XPS spectrum of vanadium in different charge–discharge states. (e)
nmvm 2
Es 2
Et
value corresponding to GITT
data from the charge to discharge state, where D is the diffusion coefficient, 18 of voltage values were selected. i is the current, nm is the mole number, Vm is the electro molar volume, S is the contact area of the electrode and electrolyte, and factor
Es 2 Et
4
S
is constant a. Therefore, D = a
Et
can be directly compared to diffusion coefficient via the
, which is easily obtained. The data in the black box correspond to the state with a capacity of approximately 150 mA h·g−1 during discharge [32]. (f) EIS curves
for various charge–discharge state. The semicircle and sloping line regions correspond to electron transfer and ionic diffusion resistance, respectively.
the wider XRD peaks and SEM image. These structural characteristics indicate the loose arrangement of sodium ions to motivate intercalation behavior. For example: The atoms on the surface of particles have defects, the interaction is weaker than the bulk atom. The atomic arrangement on the surface of particles is looser and more disordered than the atomic arrangement of the bulk phase, thus reducing the electrostatic repulsion between sodium ions [30,31]. Therefore, these results indicate that the sol–gel process produces Na3 V2 (PO4 )2 F3 with a larger interlayer spacing, which is expected to be beneficial for reducing electrostatic repulsion and achieving sodium ion deintercalation. As shown in Fig. 3a and b, the SG material electrode shows obvious sodium ion intercalation and deintercalation behavior. During the initial discharge process, it exhibits significant additional capacity between 1.2 and 3.0 V with a potential plateau at 1.38 V. Moreover, the first discharge capacity of 250 mA h·g−1 suggests the insertion of an additional two sodium ions. In the second cycle, an additional capacity of 85 mA h·g−1 is obtained between 1.2
and 3.7 V during the charge process including a potential plateau at 1.56 V. In contrast, almost no additional sodium ions were embedded in the SS material electrode. As shown in Fig. 3b, the SS material has a capacity of only 25 mA h·g-1 after being charged to 3.7 V in the second cycle. Further, a new pair of redox peaks at 1.38 V/1.56 V appeared in the dQ/dV curves of the first two cycles for the SG material (Fig. 3c). A similar phenomenon was also observed by cyclic voltammetry (Fig. 3d). The SG material exhibited a reversible process with a capacity of ∼150 mA h·g−1 over 20 cycles at 0.2 C (Fig. 3e) and a capacity retention of 72%. The charge–discharge curves for different cycles also exhibited three analogous pairs of voltage plateaus, suggesting reversible behavior (Fig. 3f). The reaction mechanism of the SG material during the charge and discharge processes was analyzed by XRD and XPS. As shown in Fig. 4a, the initial phase was maintained after sodium ions were embedded in the SG material. However, the (111) lattice spacing obviously increased with the charge process and then decreased
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Fig. 5. (a) Series of cyclic voltammograms at different scanning speeds. (b) b values at different potentials (log I = blog V + log a; for a strictly diffusion-limited reaction, b = 0.5, whereas for a capacitive current, b = 1. (c) Contributions of surface-controlled pseudocapacitive currents at a scanning rate of 0.1 mV·s−1 [33,34].
with the discharge process (Fig. 4b). The unit cell parameter c increased and the unit cell parameter b decreased when charged to 4.5 V. When discharged to 1.3 and 1.2 V, cell parameter a did not change significantly, but cell parameter c decreased, which might be due to a strengthened interaction between the atoms after the insertion of sodium ions (Fig. 4c). The XPS analysis showed that the valence state of vanadium changed from V3+ to V4+ when the charge reached 4.5 V [27]. During the intercalation process, three sodium ions occupied the Na vacancies, accompanying the V3+ to V2+ reduction reaction (Fig. 4d). The galvanostatic intermittent titration technique (GITT) was applied to characterize the migration kinetics of sodium ions (Fig.
2
s values for both SG and SS S6). As shown in Fig. 4e, the E Et materials are all of the same order of magnitude when Na migrates
2
s value of the SS material along the 3.6 and 4.0 V plateaus. The E Et decreased by about two orders of magnitude at a discharge capacity of approximately 150 mA h·g−1 (black box, Fig. 4e). However,
2
s value of the SG material maintained the same order of the E Et magnitude. This finding demonstrated that the SG material has a greater sodium ion migration ability. Subsequently, the sodium ion migration ability was reduced at the 1.4 V plateau region by about an order of magnitude of that at 3.6 V, and was further reduced by about two orders of magnitude at 1.2 V at the end of discharge. During the entire charge–discharge process, the intercalation and deintercalation of sodium ions are accompanied by electron conduction. Thus, EIS was performed to characterize the change in the electronic conductivity in a battery. As shown in Fig. 4f, during the charge process, the semicircle in the EIS spectrum gradually increased in size with the release of sodium ions, indicating an increase in impedance but a decrease in electron conduction capacity. The electron-conducting ability was minimized at 4.5 V. On the contrary, during discharging to 1.4 V, the semicircle in the EIS spectrum gradually decreased in size. This change suggested that the resistance was gradually reduced, whereas the electronconducting capacity was enhanced. In brief, both the sodium ion conductivity and the electronic conductivity were gradually reduced with the intercalation of additional sodium ions from 1.5 to 1.2 V, which implies that the sodium ion intercalation process may be diffusion controlled. To further confirm a diffusion-controlled reaction mechanism for the intercalation of additional sodium ions, a series of cyclic voltammograms were obtained at different scanning speeds in the voltage window of 1.2–4.5 V (Fig. 5a and Fig. S7a). Compared to the SS material, the extra capacity of the SG material was mainly observed at 1.2–3.6 V. Based on the equation log (I) = blog (V) + log a, plots of log (V)–log (I) were linearly fitted at the selected potentials (Fig. S7b and c), and the b values were calculated (Fig. 5b). A b value of 0.5 corresponds to a strictly diffusion-limited reaction, whereas a b value of 1 indicates a capacitive current. During the reduction process, the b value is between 0.5 and 1.0 from 2.5 to 1.5 V, sug-
gesting a mixed process of diffusion and surface control. However, between 1.5 and 1.2 V, the b value is approximately 0.5, suggesting diffusion control. During the oxidation process, the b value is approximately 0.5 between 1.25 and 1.65 V, indicating diffusion control at the beginning of the charge process. In the potential range of 1.9–2.2 V the b value is approximately 0.8, which suggest that the capacity is mainly derived from the deintercalation of surface sodium ions. Therefore, increasing the specific surface area of the electrode material can enhance the release of sodium ions at this stage. The currents contributed surface control and the diffusion control can be quantified using the formula i(E)/v1/2 = k1 v1/2 + k2 [33,34]. The k1 and k2 values can be obtained by fitting plots of i(E)/v1/2 - k1 v1/2 (Fig. S7d and e). The current contributed by surface control is less than 20% at 1.2–1.8 V and can reach approximately 50% of the total current at 1.9–3.0 V. When the potential was 1.38 V/1.56 V, the main mechanism is the deintercalation of sodium ions into the crystal lattice (Fig. 5c). Thus, increasing the lattice spacing to reduce the sodium ion intercalation activation energy and shorten the diffusion path is beneficial for enhancing the deintercalation of sodium ions.
4. Conclusions Nano-Na3 V2 (PO4 )2 F3 was synthesized by a modified sol–gel method, in which bulk dismemberment produced some degree of lattice distortion and defects. Compared with highly crystallined Na3 V2 (PO4 )2 F3 , one more Na+ ion could be reversibly embedded in Na3 V2 (PO4 )2 F3 to form Na4 V2 (PO4 )2 F3 owing to the weakened electrostatic repulsion between sodium ions. Moreover, compared with the highly crystallined material, the nano-Na3 V2 (PO4 )2 F3 electrode had a higher capacity of 250 mA h·g−1 , and also exhibited a capacity retention of 72% at 1.3–4.5 V after 20 cycles. A relatively stable sodium intercalation plateau was observed at approximately 1.38 V/1.56 V, which corresponded to a mainly diffusion-controlled process. Thus, the increase in the lattice spacing Na3 V2 (PO4 )2 F3 likely reduced the sodium ion intercalation activation energy and shortened the diffusion path, thus enhancing the deintercalation of sodium ions. This research provides a novel approach for the design of pure-phase, nano-structures to enhance Na+ intercalation in vanadium-based cathode materials and thereby improve the capacity of SIBs.
Data availability statement The raw/processed data required to reproduce these findings can-not be shared at this time due to technical or time limitations. Data will be made available on request.
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Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Manhua Peng: Conceptualization, Methodology, Validation, Investigation, Writing - original draft, Visualization, Funding acquisition. Xiayan Wang: Conceptualization, Validation, Resources, Data curation, Writing - review & editing, Supervision, Funding acquisition. Guangsheng Guo: Validation, Methodology, Supervision. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Nos. 21808006 and 21625501), the Beijing Outstanding Young Scientist Program (BJJWZYJH01201910005017), and the Beijing Municipal High-level Innovative Team Building Program (IDHT20180504). Appendix A. Supplementary data Supplementary material related to this article can be found, in the online version, at https://doi.org/10.1016/j.apmt.2020.100554. References [1] W. Huang, A. Marcelli, D. Xia, Application of synchrotron radiation technologies to electrode materials for Li- and Na-ion batteries, Adv. Energy Mater. 7 (2017), 1700460, http://dx.doi.org/10.1002/aenm.201700460. [2] Y. Fu, Q. Wei, G. Zhang, S. Sun, Advanced phosphorus-based materials for lithium/sodium-ion batteries: recent developments and future perspectives, Adv. Energy Mater. 8 (2018), 1703058, http://dx.doi.org/10.1002/aenm. 201702849. [3] C. Delmas, Sodium and sodium-ion batteries: 50 years of research, Adv. Energy Mater. 8 (2018), 1703137, http://dx.doi.org/10.1002/aenm. 201703137. [4] J.-Y. Hwang, S.-T. Myung, Y.-K. Sun, Sodium-ion batteries: present and future, Chem. Soc. Rev. 46 (2017) 3529–3614, http://dx.doi.org/10.1039/ C6CS00776G. [5] N. Yabuuchi, K. Kubota, M. Dahbi, S. Komaba, Research development on sodium-ion batteries, Chem. Rev. 114 (2014) 11636–11682, http://dx.doi.org/ 10.1021/cr500192f. [6] Y. Huang, L. Zhao, L. Li, M. Xie, F. Wu, R. Chen, Electrolytes and electrolyte/electrode interfaces in sodium-ion batteries: from scientific research to practical application, Adv. Mater. 31 (2019), 1808393, http://dx. doi.org/10.1002/adma.201808393. [7] R. Chen, R. Luo, Y. Huang, F. Wu, L. Li, Advanced high energy density secondary batteries with multi-electron reaction materials, Adv. Sci. 3 (2016), 1600051, http://dx.doi.org/10.1002/advs.201600051. [8] P. Balaya, H. Li, L. Kienle, J. Maier, Fully reversible homogeneous and heterogeneous Li storage in RuO2 with high capacity, Adv. Funct. Mater. 13 (2003) 621–625, http://dx.doi.org/10.1002/adfm.200304406. [9] M. Bianchini, P. Xiao, Y. Wang, G. Ceder, Additional sodium insertion into polyanionic cathodes for higher-energy Na-ion batteries, Adv. Energy Mater. 7 (2017), 1700514, http://dx.doi.org/10.1002/aenm.201700514. [10] H. Si, L. Seidl, E.M.L. Chu, S. Martens, J. Ma, X. Qiu, U. Stimming, O. Schneider, Impact of the morphology of V2 O5 electrodes on the electrochemical Na+ -ion intercalation, J. Electrochem. Soc. 165 (2018) A2709–A2717, http://dx.doi.org/ 10.1149/2.0621811jes. [11] W. Shen, C. Wang, H. Liu, W. Yang, Towards highly stable storage of sodium ions: A porous Na3 V2 (PO4 )3 /C cathode material for sodium-ion batteries, Chem–Eur. J. 19 (2013) 14712–14718, http://dx.doi.org/10.1002/chem. 201300005. [12] M. Ling, F. Li, H. Yi, X. Li, G. Hou, Q. Zheng, H. Zhang, Superior Na-storage performance of molten-state-blending-synthesized monoclinic NaVPO4 F nanoplates for Na-ion batteries, J. Mater. Chem. A 6 (2018) 24201–24209, http://dx.doi.org/10.1039/C8TA08842J. [13] Y. Ni, G. He, Stable cycling of -VOPO4 /NaVOPO4 cathodes for sodium-ion batteries, Electrochim. Acta 292 (2018) 47–54, http://dx.doi.org/10.1016/j. electacta.2018.09.140.
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Applied Materials Today journal homepage: www.elsevier.com/locate/apmt
Synthesis of nano-Na3 V2 (PO4 )2 F3 cathodes with excess Na+ intercalation for enhanced capacity Manhua Peng, Xiayan Wang ∗ , Guangsheng Guo ∗ Center of Excellence for Environmental Safety and Biological Effects, Beijing Key Laboratory for Green Catalysis and Separation, Department of Chemistry and Chemical Engineering, Beijing University of Technology, Beijing 100124, PR China
a r t i c l e
i n f o
Article history: Received 18 October 2019 Received in revised form 5 December 2019 Accepted 30 December 2019 Keywords: Sodium-ion battery Vanadium cathode High capacity Na+ intercalation
a b s t r a c t Vanadium-based materials are promising as cathodes for sodium-ion batteries owing to their superior cycling behavior and rate performance. However, there is a critical need for the rational implementation of Na+ intercalation to enhance the capacity. Herein, nano-Na3 V2 (PO4 )2 F3 materials were synthesized via bulk dismemberment by gas released by the decomposition of excess citric acid to trigger some degree of lattice distortion and defects. Owing to the weakened electrostatic repulsion between sodium ions in Na3 V2 (PO4 )2 F3 with an enlarged interlayer spacing, one more Na+ ion could be readily embedded to form Na4 V2 (PO4 )2 F3 . A new reaction plateau at 1.38 V/1.56 V accompanied the V3+ /V2+ redox reaction, inducing a capacity of 250 mA h·g−1 , which is higher than that of highly crystallined Na3 V2 (PO4 )2 F3 . A noteworthy reversible redox phenomenon involving three sodium ions was confirmed during the cycling process. The nano-Na3 V2 (PO4 )2 F3 electrode exhibited a capacity retention of 72% at 1.3–4.5 V after 20 cycles. Thus, this novel synthesis routine can provide vanadium-based cathode materials with enhanced Na+ intercalation properties, which can be applied to improve the capacity of sodium-ion batteries. © 2020 Elsevier Ltd. All rights reserved.
1. Introduction Owing to the abundance a low price of sodium, sodium-ion batteries (SIBs) can serve as backup power supplies for large-scale electronic devices [1–3]. However, a key issue facing SIBs is their low energy densities relative to those of lithium-ion batteries. Moreover, in crystals, it is more difficult to deintercalate sodium ions than lithium ions [4–6]. Increasing number of electrons transferred by a transition metal via the introduction of multiple valence states is an effective way to improve the battery capacity [7–9]. In this regard, vanadium, which has multiple valence states, could allow enhanced sodium ion intercalation [9]. Recently, vanadiumbased materials such as Na3 V2 O2-x (PO4 )2 F1+x(0≤x≤2) , NaVPO4 F, Na3 V2 (PO4 )3 , NaVOPO4 , and Na7 V4 (P2 O7 )4 (PO4 ) have been developed as cathode materials for SIBs [10–19]. Among these materials,
Abbreviations: SIB, sodium ion battery; SG, material material prepared by a modified sol–gel reaction; SS, material material prepared by a standard solid-state reaction; XRD, X-ray diffraction; SEM, scanning electron microscopy; TEM, transmission electron microscopy; XPS, X-ray photoelectron spectroscopy; EIS, impedance spectroscopy; GITT, galvanostatic intermittent titration technique; TGA, Thermogravimetric analysis. ∗ Corresponding authors. E-mail addresses: [email protected] (X. Wang), [email protected] (G. Guo). https://doi.org/10.1016/j.apmt.2020.100554 2352-9407/© 2020 Elsevier Ltd. All rights reserved.
pseudo-layered Na3 V2 (PO4 )2 F3 has a stable electronic structure and a higher operating voltage. However, the reported theoretical capacity of this material is only 130 mA h·g−1 because only two sodium ions can be reversibly deintercalated, corresponding to the V3+ /V4+ redox reaction at the voltage of 2.0–4.5 V [20–27]. Therefore, to improve the theoretical capacity of vanadium-based electrode materials, it is necessary to further increase the number of electrons transferred and to enhance sodium ion deintercalation. One optional method to increase the theoretical capacity of Na3 V2 (PO4 )2 F3 is to take advantage of a broader voltage window, such as 1.0–4.5 V instead of 2.0–4.5 V. But the deintercalation or embedding of excess sodium ions in Na3 V2 (PO4 )2 F3 materials can be realized by doping with Cr or Al at a broader voltage window of 1.0–4.5 V. Nevertheless, doping with such inert elements could cause the capacity to decrease. Thus, using pure-phase Na3 V2 (PO4 )2 F3 or doping with active elements may be a more favorable approach. However, it is difficult to intercalate a sodium ion in pure-phase of Na3 V2 (PO4 )2 F3 to form Na4 V2 (PO4 )2 F. Ceder and co-workers suggested that the intercalation behavior is mainly limited by the strong electrostatic repulsion between sodium ions rather than by the valence states of vanadium [9,28]. In view of the above analysis, this work focused on embedding excess sodium ions in pure-phase Na3 V2 (PO4 )2 F3 with the aim of enhancing the capacity. To achieve this goal, we synthesized nano-vanadium materials with increasing the interlayer spacing
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Fig. 1. Schematic of the modified sol–gel process for preparing nano-Na3 V2 (PO4 )2 F3 .
Fig. 2. (a) XRD pattern of the SS and SG materials. (b, c) SEM and (e, f) TEM images of the SS and SG materials (insets: images at higher magnification). (d) Unit cell of Na3 V2 (PO4 )2 F3 and the corresponding lattice spacings.
via a modified sol–gel method, which involved bulk dismemberment by gas originating from the decomposition of excess citric acid in the precursor. It is expected that the intercalation of three sodium ions could be achieved by effectively weakening the electrostatic repulsion in Na3 V2 (PO4 )2 F3 by increasing the interlayer spacing. The intercalation mechanism was systematically investigated by microstructural characterization and electrochemical measurements.
2. Materials and methods
2.2. Characterization The phases of the synthesized materials were characterized by X-ray diffraction (XRD) D8 Advance DaVinci, Bruker, Germany) equipped with a Cu target ( = 1.5418 Å) in the scanning range of 2 = 10–90 ◦ C. Scanning electron microscopy (SEM; Hitachi S4300, Japan) and transmission electron microscopy (TEM; Tecnai G2 F20, FEI, USA.) were utilized to analyze the morphologies of the materials. X-ray photoelectron spectroscopy (XPS, ThermoFischer, ESCALAB 250Xi) was used to determine the chemical environment of the transition metal at different cutoff voltages. Thermogravimetric analysis (TGA, France setaram company, LABSYS EVO) were utilized to analyze thermal decomposition of materials.
2.1. Material synthesis 2.3. Electrochemical performance Nano-Na3 V2 (PO4 )2 F3 was synthesized by a modified sol–gel reaction (denoted as SG material). The reagents C6 H8 O7 , H3 PO4 , NH4 VO3 , NaF, and NH4 F (stoichiometric ratio of 1.5:1.0:1.0:1.7:0.6) were dissolved in deionized water to form a homogeneous solution. Then, the solution was heated at 85 ◦ C for 6 h until completely dried. The obtained powder was placed in a tube furnace and heated at 300 ◦ C for 2 h in an argon atmosphere. The temperature was then increased to 600–700 ◦ C at a high speed of 30 ◦ C min−1 and the material was heated for a further 30 min. For comparison, highly crystallined Na3 V2 (PO4 )2 F3 was prepared by a standard solid-state reaction (denoted as SS material) according to a literature procedure [29].
To test the electrochemical performance of the synthesized electrode materials, the electrode material, acetylene black, and a binder of PTFE were mixed in a ratio of 70:20:10 to make an electrode plate. After the electrode plate was dried in an oven at 100 ◦ C for 10 h, it was cut into small wafers with masses of approximately 3–5 mg. Batteries were assembled in an argon-filled glove box using an electrolyte of 1 mol·L−1 NaClO4 in propylene carbonate with 2% fluoroethylene carbonate, a glass fiber film (grade GF/D, Whatman, USA), and a pure sodium plate as the negative electrode material. Electrochemical testing was performed at room temperature using a NEWARE battery test system (China). GITT was executed
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Fig. 3. (a, b) First cycle and second cycles of SS and SG materials in a voltage window of 1.2–4.5 V (1 C corresponds to 130 mA h·g−1 ). (c) Differential curve of dQ/dV. (d) Cyclic voltammograms. (e) Cycling stability. (f) Charge–discharge curves.
by NEWARE system. The battery was first charged for 5 min and then rest for 0.5 h. The whole charging and discharging process was performed at a current rate of 0.1 C. Electrochemical impedance spectroscopy (EIS) was performed on a ZAHNER system (Germany) using frequencies from 1 mHz to 1 MHz at an AC voltage of 5 mV.
3. Results and discussion The structural characteristics of the SS and SG materials were thoroughly investigated. To obtain a loose and expanded Na3 V2 (PO4 )2 F3 structure, an increased amount of chelating agent (citric acid) was added during gel formation (Fig. 1). The decomposition of citric acid could release a gas (carbon dioxide, water) that then dismembered the bulk into small pieces. This process is analogous to the “steamed bread” process, in which yeast works like citric acid. Thermogravimetric analysis confirmed the occurrence of this phenomenon because the weight loss reached 30% between 150 and 300 ◦ C, which mainly derived from the decomposition of citric acid (Fig. S1). Moreover, an additional weight loss of 15% was observed from 300 to 500 ◦ C, probably owing to further carbonization of citric acid and crystal water (Fig. S2). Because of these decomposition processes, the prepared SG material is fluffy,
which is different from the powdered SS material formed by the standard solid-state reaction (Fig. S3). The XRD patterns (Fig. 2a) confirm that both pure-phase SG and SS materials were synthesized. However, the characteristic peaks corresponding to the (002), (212) (113), and (202) facets of the SG material were broader than those of the SS material. Furthermore, the XRD peaks of the SG material were slightly shifted toward lower angles, which confirms that the lattice spacing was enlarged. (Fig. 2a, Fig. S4). SEM and TEM images (Fig. 2b, c and e, f) show that the sizes of the SS and spherical SG materials were approximately 1 m and 100 nm in diameter, respectively. Furthermore, after the heat treatment during the sol–gel process, many defects were formed on the surface of the spherical SG particles (Fig. 2f), whereas the SS material has a relative well-defined crystal plane (Fig. 2e). Typically, sodium ions are located in the interlayer of the (002) plane (Fig. 2d), where they migrate mainly along the xy axis. The SG material had a wider (002) lattice spacing (d = 0.556 nm) than that (d = 0.521 nm) of the SS material (Fig. S5). The intercalation behavior is limited mainly by the strong electrostatic repulsion between sodium ions affected by the compactness of sodium ions arrangement. The obtained a structure of materials including the lattice distortion, lattice defects and wider (002) lattice distance was confirmed by
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Es 2
Fig. 4. (a) XRD patterns a SG material electrode in the initial state, when charged to 4.5 V, and when discharged to 1.3 and 1.2 V. (b) Magnification of the XRD patterns in the region shown in the black box in (a). (c) Unit cell parameters. (d) XPS spectrum of vanadium in different charge–discharge states. (e)
nmvm 2
Es 2
Et
value corresponding to GITT
data from the charge to discharge state, where D is the diffusion coefficient, 18 of voltage values were selected. i is the current, nm is the mole number, Vm is the electro molar volume, S is the contact area of the electrode and electrolyte, and factor
Es 2 Et
4
S
is constant a. Therefore, D = a
Et
can be directly compared to diffusion coefficient via the
, which is easily obtained. The data in the black box correspond to the state with a capacity of approximately 150 mA h·g−1 during discharge [32]. (f) EIS curves
for various charge–discharge state. The semicircle and sloping line regions correspond to electron transfer and ionic diffusion resistance, respectively.
the wider XRD peaks and SEM image. These structural characteristics indicate the loose arrangement of sodium ions to motivate intercalation behavior. For example: The atoms on the surface of particles have defects, the interaction is weaker than the bulk atom. The atomic arrangement on the surface of particles is looser and more disordered than the atomic arrangement of the bulk phase, thus reducing the electrostatic repulsion between sodium ions [30,31]. Therefore, these results indicate that the sol–gel process produces Na3 V2 (PO4 )2 F3 with a larger interlayer spacing, which is expected to be beneficial for reducing electrostatic repulsion and achieving sodium ion deintercalation. As shown in Fig. 3a and b, the SG material electrode shows obvious sodium ion intercalation and deintercalation behavior. During the initial discharge process, it exhibits significant additional capacity between 1.2 and 3.0 V with a potential plateau at 1.38 V. Moreover, the first discharge capacity of 250 mA h·g−1 suggests the insertion of an additional two sodium ions. In the second cycle, an additional capacity of 85 mA h·g−1 is obtained between 1.2
and 3.7 V during the charge process including a potential plateau at 1.56 V. In contrast, almost no additional sodium ions were embedded in the SS material electrode. As shown in Fig. 3b, the SS material has a capacity of only 25 mA h·g-1 after being charged to 3.7 V in the second cycle. Further, a new pair of redox peaks at 1.38 V/1.56 V appeared in the dQ/dV curves of the first two cycles for the SG material (Fig. 3c). A similar phenomenon was also observed by cyclic voltammetry (Fig. 3d). The SG material exhibited a reversible process with a capacity of ∼150 mA h·g−1 over 20 cycles at 0.2 C (Fig. 3e) and a capacity retention of 72%. The charge–discharge curves for different cycles also exhibited three analogous pairs of voltage plateaus, suggesting reversible behavior (Fig. 3f). The reaction mechanism of the SG material during the charge and discharge processes was analyzed by XRD and XPS. As shown in Fig. 4a, the initial phase was maintained after sodium ions were embedded in the SG material. However, the (111) lattice spacing obviously increased with the charge process and then decreased
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Fig. 5. (a) Series of cyclic voltammograms at different scanning speeds. (b) b values at different potentials (log I = blog V + log a; for a strictly diffusion-limited reaction, b = 0.5, whereas for a capacitive current, b = 1. (c) Contributions of surface-controlled pseudocapacitive currents at a scanning rate of 0.1 mV·s−1 [33,34].
with the discharge process (Fig. 4b). The unit cell parameter c increased and the unit cell parameter b decreased when charged to 4.5 V. When discharged to 1.3 and 1.2 V, cell parameter a did not change significantly, but cell parameter c decreased, which might be due to a strengthened interaction between the atoms after the insertion of sodium ions (Fig. 4c). The XPS analysis showed that the valence state of vanadium changed from V3+ to V4+ when the charge reached 4.5 V [27]. During the intercalation process, three sodium ions occupied the Na vacancies, accompanying the V3+ to V2+ reduction reaction (Fig. 4d). The galvanostatic intermittent titration technique (GITT) was applied to characterize the migration kinetics of sodium ions (Fig.
2
s values for both SG and SS S6). As shown in Fig. 4e, the E Et materials are all of the same order of magnitude when Na migrates
2
s value of the SS material along the 3.6 and 4.0 V plateaus. The E Et decreased by about two orders of magnitude at a discharge capacity of approximately 150 mA h·g−1 (black box, Fig. 4e). However,
2
s value of the SG material maintained the same order of the E Et magnitude. This finding demonstrated that the SG material has a greater sodium ion migration ability. Subsequently, the sodium ion migration ability was reduced at the 1.4 V plateau region by about an order of magnitude of that at 3.6 V, and was further reduced by about two orders of magnitude at 1.2 V at the end of discharge. During the entire charge–discharge process, the intercalation and deintercalation of sodium ions are accompanied by electron conduction. Thus, EIS was performed to characterize the change in the electronic conductivity in a battery. As shown in Fig. 4f, during the charge process, the semicircle in the EIS spectrum gradually increased in size with the release of sodium ions, indicating an increase in impedance but a decrease in electron conduction capacity. The electron-conducting ability was minimized at 4.5 V. On the contrary, during discharging to 1.4 V, the semicircle in the EIS spectrum gradually decreased in size. This change suggested that the resistance was gradually reduced, whereas the electronconducting capacity was enhanced. In brief, both the sodium ion conductivity and the electronic conductivity were gradually reduced with the intercalation of additional sodium ions from 1.5 to 1.2 V, which implies that the sodium ion intercalation process may be diffusion controlled. To further confirm a diffusion-controlled reaction mechanism for the intercalation of additional sodium ions, a series of cyclic voltammograms were obtained at different scanning speeds in the voltage window of 1.2–4.5 V (Fig. 5a and Fig. S7a). Compared to the SS material, the extra capacity of the SG material was mainly observed at 1.2–3.6 V. Based on the equation log (I) = blog (V) + log a, plots of log (V)–log (I) were linearly fitted at the selected potentials (Fig. S7b and c), and the b values were calculated (Fig. 5b). A b value of 0.5 corresponds to a strictly diffusion-limited reaction, whereas a b value of 1 indicates a capacitive current. During the reduction process, the b value is between 0.5 and 1.0 from 2.5 to 1.5 V, sug-
gesting a mixed process of diffusion and surface control. However, between 1.5 and 1.2 V, the b value is approximately 0.5, suggesting diffusion control. During the oxidation process, the b value is approximately 0.5 between 1.25 and 1.65 V, indicating diffusion control at the beginning of the charge process. In the potential range of 1.9–2.2 V the b value is approximately 0.8, which suggest that the capacity is mainly derived from the deintercalation of surface sodium ions. Therefore, increasing the specific surface area of the electrode material can enhance the release of sodium ions at this stage. The currents contributed surface control and the diffusion control can be quantified using the formula i(E)/v1/2 = k1 v1/2 + k2 [33,34]. The k1 and k2 values can be obtained by fitting plots of i(E)/v1/2 - k1 v1/2 (Fig. S7d and e). The current contributed by surface control is less than 20% at 1.2–1.8 V and can reach approximately 50% of the total current at 1.9–3.0 V. When the potential was 1.38 V/1.56 V, the main mechanism is the deintercalation of sodium ions into the crystal lattice (Fig. 5c). Thus, increasing the lattice spacing to reduce the sodium ion intercalation activation energy and shorten the diffusion path is beneficial for enhancing the deintercalation of sodium ions.
4. Conclusions Nano-Na3 V2 (PO4 )2 F3 was synthesized by a modified sol–gel method, in which bulk dismemberment produced some degree of lattice distortion and defects. Compared with highly crystallined Na3 V2 (PO4 )2 F3 , one more Na+ ion could be reversibly embedded in Na3 V2 (PO4 )2 F3 to form Na4 V2 (PO4 )2 F3 owing to the weakened electrostatic repulsion between sodium ions. Moreover, compared with the highly crystallined material, the nano-Na3 V2 (PO4 )2 F3 electrode had a higher capacity of 250 mA h·g−1 , and also exhibited a capacity retention of 72% at 1.3–4.5 V after 20 cycles. A relatively stable sodium intercalation plateau was observed at approximately 1.38 V/1.56 V, which corresponded to a mainly diffusion-controlled process. Thus, the increase in the lattice spacing Na3 V2 (PO4 )2 F3 likely reduced the sodium ion intercalation activation energy and shortened the diffusion path, thus enhancing the deintercalation of sodium ions. This research provides a novel approach for the design of pure-phase, nano-structures to enhance Na+ intercalation in vanadium-based cathode materials and thereby improve the capacity of SIBs.
Data availability statement The raw/processed data required to reproduce these findings can-not be shared at this time due to technical or time limitations. Data will be made available on request.
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Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Manhua Peng: Conceptualization, Methodology, Validation, Investigation, Writing - original draft, Visualization, Funding acquisition. Xiayan Wang: Conceptualization, Validation, Resources, Data curation, Writing - review & editing, Supervision, Funding acquisition. Guangsheng Guo: Validation, Methodology, Supervision. Acknowledgements This work was financially supported by the National Natural Science Foundation of China (Nos. 21808006 and 21625501), the Beijing Outstanding Young Scientist Program (BJJWZYJH01201910005017), and the Beijing Municipal High-level Innovative Team Building Program (IDHT20180504). Appendix A. Supplementary data Supplementary material related to this article can be found, in the online version, at https://doi.org/10.1016/j.apmt.2020.100554. References [1] W. Huang, A. Marcelli, D. Xia, Application of synchrotron radiation technologies to electrode materials for Li- and Na-ion batteries, Adv. Energy Mater. 7 (2017), 1700460, http://dx.doi.org/10.1002/aenm.201700460. [2] Y. Fu, Q. Wei, G. Zhang, S. Sun, Advanced phosphorus-based materials for lithium/sodium-ion batteries: recent developments and future perspectives, Adv. Energy Mater. 8 (2018), 1703058, http://dx.doi.org/10.1002/aenm. 201702849. [3] C. Delmas, Sodium and sodium-ion batteries: 50 years of research, Adv. Energy Mater. 8 (2018), 1703137, http://dx.doi.org/10.1002/aenm. 201703137. [4] J.-Y. Hwang, S.-T. Myung, Y.-K. Sun, Sodium-ion batteries: present and future, Chem. Soc. Rev. 46 (2017) 3529–3614, http://dx.doi.org/10.1039/ C6CS00776G. [5] N. Yabuuchi, K. Kubota, M. Dahbi, S. Komaba, Research development on sodium-ion batteries, Chem. Rev. 114 (2014) 11636–11682, http://dx.doi.org/ 10.1021/cr500192f. [6] Y. Huang, L. Zhao, L. Li, M. Xie, F. Wu, R. Chen, Electrolytes and electrolyte/electrode interfaces in sodium-ion batteries: from scientific research to practical application, Adv. Mater. 31 (2019), 1808393, http://dx. doi.org/10.1002/adma.201808393. [7] R. Chen, R. Luo, Y. Huang, F. Wu, L. Li, Advanced high energy density secondary batteries with multi-electron reaction materials, Adv. Sci. 3 (2016), 1600051, http://dx.doi.org/10.1002/advs.201600051. [8] P. Balaya, H. Li, L. Kienle, J. Maier, Fully reversible homogeneous and heterogeneous Li storage in RuO2 with high capacity, Adv. Funct. Mater. 13 (2003) 621–625, http://dx.doi.org/10.1002/adfm.200304406. [9] M. Bianchini, P. Xiao, Y. Wang, G. Ceder, Additional sodium insertion into polyanionic cathodes for higher-energy Na-ion batteries, Adv. Energy Mater. 7 (2017), 1700514, http://dx.doi.org/10.1002/aenm.201700514. [10] H. Si, L. Seidl, E.M.L. Chu, S. Martens, J. Ma, X. Qiu, U. Stimming, O. Schneider, Impact of the morphology of V2 O5 electrodes on the electrochemical Na+ -ion intercalation, J. Electrochem. Soc. 165 (2018) A2709–A2717, http://dx.doi.org/ 10.1149/2.0621811jes. [11] W. Shen, C. Wang, H. Liu, W. Yang, Towards highly stable storage of sodium ions: A porous Na3 V2 (PO4 )3 /C cathode material for sodium-ion batteries, Chem–Eur. J. 19 (2013) 14712–14718, http://dx.doi.org/10.1002/chem. 201300005. [12] M. Ling, F. Li, H. Yi, X. Li, G. Hou, Q. Zheng, H. Zhang, Superior Na-storage performance of molten-state-blending-synthesized monoclinic NaVPO4 F nanoplates for Na-ion batteries, J. Mater. Chem. A 6 (2018) 24201–24209, http://dx.doi.org/10.1039/C8TA08842J. [13] Y. Ni, G. He, Stable cycling of -VOPO4 /NaVOPO4 cathodes for sodium-ion batteries, Electrochim. Acta 292 (2018) 47–54, http://dx.doi.org/10.1016/j. electacta.2018.09.140.
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