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On the Mechanism of the Improved Operation Voltage of Rhombohedral Nickel Hexacyanoferrate as Cathodes for Sodium-Ion Batteries Zhuan Ji, Bo Han, Haitao Liang, Chenggang Zhou, Qiang Gao, Kaisheng Xia, and Jinping Wu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b11070 • Publication Date (Web): 18 Nov 2016 Downloaded from http://pubs.acs.org on November 22, 2016

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ACS Applied Materials & Interfaces

On the Mechanism of the Improved Operation Voltage of Rhombohedral Nickel Hexacyanoferrate as Cathodes for Sodium-Ion Batteries Zhuan Ji, Bo Han*, Haitao Liang, Chenggang Zhou*, Qiang Gao, Kaisheng Xia, Jinping Wu Sustainable Energy Laboratory, Faculty of Materials Science and Chemistry, China University of Geosciences, Wuhan 430074, Hubei, China P.R. E-mail: [email protected] (B. H.)

[email protected] (C. Z.)

Abstract We reported a rhombohedral Na-rich nickel hexacyanoferrate (r-NiHCF) with high discharge voltage, which also possesses long cycle stability and excellent rate capability when serving as the cathode material of Na-ion batteries. First-principles calculations suggest that the high working voltage of r-NiHCF is correlated to the asymmetric residence of Na+ ions in the rhombohedral framework in parallel with the low charge density at the Fe2+ ions. In both aqueous and ether-based electrolytes, r-NiHCF exhibits higher voltage than that of cubic NiHCF. Rate and cycle experiments indicate that, r-NiHCF delivers a specific capacity of 66.8 mAh g-1 at the current density of 80 mA g-1, which is approximate to the theoretical capacity of r-NiHCF. A capacity retention of 96% can be achieved after 200 cycles. The excellent stability of r-NiHCF can be assigned to the absence of rhombohedral-cubic phase transition and negligible volume variation during electrochemical redox, as proven by the ex-situ XRD patterns at different depth of charge/discharge and the DFT calculations, respectively. Keywords: Na-ion batteries; Cathode Material; Rhombohedral nickel hexacyanoferrate; High discharge voltage; Density functional theory study; Mechanism;

Introduction Advanced electric energy storage (EES) technologies represent an important branch for electricity peak shaving and smart grid managements,1, 2 where secondary batteries

ACS Paragon Plus Environment


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have been realized to be one of the most promising choices.3-5 For large-scale stationary energy storage, beyond the requirement of high energy density, long operation life, high rate capability as well as low deployment cost are always demanded.6-8 The rich abundance of sodium, in contrast with lithium, suggests the feasibility of sodium-based batteries for future EES. In fact, high-temperature sodium-sulfur batteries have already been commercialized.9-11 However, the high operating temperature (300-350°C) of sodium-sulfur batteries would lead to severe safety risks. Developing room-temperature sodium-ion batteries (SIBs) has become a growing interest in recent years.12-16

To apply SIBs in EES, the most critical issue is to figure out an appropriate cathode material, which determines the energy density, stability and rate capability of SIBs. Many types of cathodes, including transition-metal oxides,17, compounds19,

20

and organic compounds,21,

22

18

polyanionic

have been reported. Particularly,

low-cost Prussian blue analogues (PBA, common formula: AaMbFe(CN)6·xH2O, where A represents alkali metal ion and M stands for transition metal ion) has been widely accepted as a promising candidate. The rigid and unimpeded intrinsic channels of PBA give rise to rapid transportation for alkali ions during electrochemical redox, indicating autogenous rate performances.23, 24 In general, most of the studied PBA occur as cubic phase with both outer M ions (typically Fe2+, Co2+ and Mn2+) and the six-fold coordinated Fe2+ or Fe3+ ions participating in redox (dual mode),25-37 providing a high theoretical specific capacity of up to 170 mAh g-1. i.e. the manganese hexacyanoferrate (MnHCF) reported by Goodenough et al26 implements a high initial discharge capacity of 123 mAh g-1; the iron hexacyanoferrate (FeHCF) reported by Guo et al28 delivers 170 mAh g-1 initial discharge capacity; the cobalt hexacyanoferrate (CoHCF) reported by Moritomo et al25 hosts an initial high discharge capacity of 135 mAh g-1. However, these cubic phase PBA cathode materials always encounter fast capacity loss during cycling. Delicate controlling of the contents of vacancies and zeolite water,27-29 or designing core-shell architectures, 38-40

could help improving the cycling durability of cathodes.


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In contrast, utilizing only monomeric redox of the inner iron species by using electrochemically inert Ni2+ or Zn2+ as outer ions could considerably enhance the cathode stability.41-43 Guo et al41 reported an extremely stable “zero-strain” nickel hexacyanoferrate (NiHCF) with cubic lattice, which reserves 99.7% of initial capacity (66 mAh g-1) after 200 galvanostatic cycles; Dai et al42 reported mesoporous NiHCF with cubic lattice, which exhibits no capacity loss (initial capacity 55 mAh g-1) after 180 cycles at a high current density of 100 mA g-1. Apparently, although satisfactory cyclic performance can be obtained, the monomeric redox sacrifices almost half of the theoretical capacity of PBA. At this level, a feasible approach to improve the cathode energy density is to promote the operation voltage. In recent years, rhombohedral PBA has attracted increasing attentions.26, 28, 35-37 The concentration of defects and the corresponding coordinative water in PBA framework, particularly for the Na+ enriched phases, would reduce the lattice symmetry to form rhombohedral lattice, which generally causes voltage variation of the Fe3+/Fe2+ redox couple.35, 36 i.e. the charge/discharge potentials of cubic FeHCF (3.6V/3.4V for low-spin Fe3+/Fe2+, 2.9V/2.7V for high-spin Fe3+/Fe2+)28,

30

change to 3.3V/3.29V and 3.11V/3.0V in

rhombohedral FeHCF,36 respectively, implying an average effect of the inner and outer redox couples. The voltage variation becomes more visible in the rhombohedral MnHCF. The potential of low-spin Fe3+/Fe2+ increases from 3.31V/3.24V in cubic phase26 to 3.53V/3.44V in rhombohedral phase,35 which is almost overlapped with that of high-spin Mn3+/Mn2+. These experimental reports indicate that a rhombohedral lattice may give a chance to tune the operating voltage of the inner Fe3+/Fe2+ redox couple for higher energy density, which is particularly important for NiHCFs where only the inner couple contributes to the capacity.

In the present work, we reported a rhombohedral NiHCF by simple co-precipitation method, which delivers high operation voltage, excellent cyclic performance and superior rate capability. Using density functional theory (DFT) simulations, we identified that the high operation voltage can be ascribed to the asymmetric residence



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of Na+ ions in the rhombohedral framework in combination with the low charge density at the Fe2+ ions. Our results not only unravel the deep insight on the mechanism of the improved operation voltage of the rhombohedral NiHCF but provide an applicable candidate for developing high performance room-temperature SIBs.

Experimental Section

Preparation and characterization of NiHCFs. All analytical-grade chemicals utilized in the synthesis procedure were purchased from Aladdin China and were used without purification. The synthesis of rhombohedral and cubic NiHCF (denoted by r-NiHCF and c-NiHCF, respectively) followed a typical co-precipitation method. For r-NiHCF, NiCl2·6H2O (1.604 g), sodium citrate (5.475 g) and PVP (3.750 g) precursors were dissolved in ultrapure water (250 mL). The solution was then mixed with the sodium ferrocyanide solution (3.267 g of Na4Fe(CN)6·10H2O in 250 mL ultrapure water) with magnetic stirring till clarification. The mixture was let stand for 72 h at 30 °C. The light blue precipitate was centrifuged and washed by ultrapure water and ethanol alternatively for 3 times, followed by a 24 h desiccation at 60 °C to obtain the target product. For c-NiHCF, the differences are that the dosage of nickel precursors was increased to 2.377 g, while K3Fe(CN)6 (0.824 g) was employed as hexacyanoferrate source instead. Here, orange powder products were obtained. The morphological features of both products were recorded via FE-SEM (SU8010, Hitachi, Japan). The crystalline phases were determined via X-ray diffraction (XRD-6000, Shimadzu, Japan) using Cu Kα as the radiation source at the operation voltage of 40 kV and a scan rate of 0.05° min-1. The spectrum were collected by infrared (Thermo Fisher Scientific, Nicolet 6700, USA) and Raman spectroscopies (LabRAM HR800, Horiba JobinYvon, France). Elemental composition of C, N, H were determined by Elementar Analysensysteme (Vario EL Cube, GmbH, Germany), and the Ni, Fe, Na, K contents were determined by ICP-AES (IRIS Intrepid II XSP, Thermo Elemental, USA). The water contents in the materials were determined by TG-DSC (STA449F3,


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NETZSCH, Germany), where the temperature range was set to be 30-600 °C with a heating rate of 10 °C min-1 under N2 atmosphere.

Electrochemical Tests. The cathode slurry, composed of 75 wt% of r-NiHCF or c-NiHCF active material, 15 wt% of SuperP conductive carbon and 10 wt% of poly(vinylidene fluoride) (PVDF) binder, was dispersed by 3-methyl-pyrrolidone (NMP) and next casted onto aluminum current collector, followed by vacuum drying at 70 oC for 12 h. The cyclic voltammetry (CV) patterns of the prepared r-NiHCF and c-NiHCF cathodes were first scanned (VMP3 potentiastats, Bio-logic, France) in 1 M Na2SO4 aqueous electrolyte through a three-electrode system, where a commercial saturated calomel electrode (SCE) and a platinum electrode were set as reference and counter electrode, respectively. The potentials were converted relative to the standard hydrogen electrode (SHE). The electrochemical performances of r- and c-NiHCF serving as Na-ion batteries cathodes were evaluated with standard CR2025 coin cells. 1 mol L-1 NaClO4 in EC/DEC (1:1, v:v), sodium metal (16 mm in diameter) and glass-fiber membrane (GF/A, Whatman) were selected as electrolyte, anode and separator, respectively. Coin cell assembly was conducted in argon atmosphere at ambient temperature. The assembled cells were aged for 12h before test to ensure full absorption of electrolyte into the electrodes. The Galvanostatic tests were carried out using an Arbin tester (BT2000, USA) at the potential window of 2.0-3.9 V (vs. Na+/Na).

Computational methods. To obtain the crystalline structure of r-NiHCF, we started from the lattice of r-MnHCF provided by Goodenough et al35 and carried out the Rietveld refinement44, 45 using the FullProf package46. Using the refined structure of r-NiHCF and the reported lattice of c-NiHCF, we further conducted DFT+U calculations, which can well deal with the on-site self-interaction error among the Fe-d and Ni-d electrons, to reveal their electronic structures and energetics through the Vienna ab initio simulation package (VASP) with the additional Hubbard-U terms of UFe=3.0 eV and UNi=3.4 eV, respectively.47, 48 Spin-polarized generalized gradients



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approximation (GGA) using the Perdew–Burke–Ernzerhof (PBE) exchange correlation functional was utilized to deal with the exchange-correlation effects.49 The projector augmented wave (PAW) pseudopotentials were employed to represent the electron-ion interactions while the plane-wave basis sets were utilized to describe the valence orbitals with the kinetic energy cutoff of 450 eV. The Brillouin zone integration was sampled using the Monkhorst-Pack scheme (3×3×1 for r-NiHCF and 2×2×2 for c-NiHCF). The structural relaxation was performed using the conjugated gradient minimization method with the maximum force in each atom less than 0.05 eV/Å and the energy convergence accuracy set to be 10-5 eV. Population analysis was performed on the basis of the Bader charge division scheme.50

Results and Discussion

Although synthesized at the same condition, r- and c-NiHCF exhibit completely different morphological features. In detail, r-NiHCF appears as aggregated granular particles with the sizes around 30-50 nm (Fig. 1a), while c-NiHCF occurs as typical cubes with the sizes varying in the range of 300-600 nm (Fig. 1b). Combining the results from ICP-AES (Table S1) and elemental analysis (Table S2), the chemical formulas of r- and c-NiHCF products are determined to be Na1.46Ni[Fe(CN)6]0.83□ 0.17·2.2H2O

and K0.04Ni[Fe(CN)6]0.67□0.33·4.0H2O respectively. Here, □ stands for

[Fe(CN)6]3- or [Fe(CN)6]4- vacancies where water molecules accumulate and part of them would coordinate with the Ni2+ ions to complement the octahedrons in the framework.51 TGA measurements (Fig. 1c) show that, from room temperature to 300 o

C, c-NiHCF has a weight loss of 26.35%, which is much higher than that of r-NiHCF

(12.87%). The weight losses can be assigned to the water molecules at surfaces, in the interstitial channels and coordinative water in the vacancies.52 Within the temperature range of 300-600 oC, the weight losses are related to the decomposition of NiHCF and the liberation of C≡N units.52 The relatively lower weight loss of r-NiHCF in this stage indicates its higher stability, in accordance with its lower vacancy content. The FT-IR spectrum of both two samples (Fig. 1d) exhibit typical features of


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hexacyanoferrate compounds. The different vacancy and water contents lead to slightly different adsorption features of the water molecules. The adsorption peaks at 2100 and 2165 cm-1 of c-NiHCF, which belong to the stretch mode of the C≡N groups coordinated with the inner Fe3+ (Fe3+←C), degenerate to a singular adsorption at 2083 cm-1 of r-NiHCF (Fe2+ ←C), which also cause slight difference for the in-plane deformation of Fe-C (589 cm-1 for r-NiHCF and 594 cm-1 for c-NiHCF).53 In the Raman spectrum (Fig. S1), r-NiHCF shows two peaks at 2108 and 2142 cm-1, while c-NiHCF gives only one peak at 2185 cm-1, agreeing well with previous reports.54 Figure 1e displays the XRD patterns of the two samples, in which the sharp and strong diffraction peaks indicate well crystalline nature. The profile of c-NiHCF can be well indexed with standard Prussian blue Fe4[Fe(CN)6]3 (JCPDS No. 52-1907), indicating typical face-centered cubic phase. In contrast, the major peaks of r-NiHCF are approximate to that of c-NiHCF, despite the splitting peaks at 15, 24.6, 39.3 and 50.3°, which are similar to the reported XRD patterns of rhombohedral MnHCF.26, 35 We first performed Pawley fitting and refinement according to the XRD profile, which generates a rhombohedral lattice (a=b=7.386 Å, c=17.279 Å; α=β=90°, γ=120°, Fig. S2) with an R-3 symmetry. Due to the similarities between our r-NiHCF sample and the reported rhombohedral HCFs,35,36 we naturally carried out Rietveld refinement (Fig. 1f) using the atomic coordinates and structural parameters from ref 35. The obtained structure was employed as the model for the successive DFT calculations.



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Figure 1. SEM images of (a) r-NiHCF and (b) c-NiHCF; (c) TGA curves; (d) FT-IR spectrum; (e) XRD patterns and (f) the structure refinement result of r-NiHCF.

The CV profiles of the c- and r-NiHCF at the scan rate of 0.1 mV s-1 are presented in Fig. 2. Both two phases exhibit only one pair of well-defined and symmetric redox peaks of Fe3+/Fe2+, suggesting excellent reversibility. In contrast, nickel shows no electrochemical activity in both cases, in agreement with typical NiHCFs.23 Particularly, r-NiHCF exhibits an average charge/discharge potential of 0.66 V (0.67/0.65 V), which is 100 mV higher than that of c-NiHCF (0.57/0.55 V). To validate this phenomenon, we performed DFT+U calculations for the two phases, where the above refined structure of r-NiHCF and typical c-NiHCF structure were utilized. The discharge voltages were calculated via the following expression55:


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V = ( Echarge + nENa − Edischarge ) / n where Echarge and Edischarge represents the energy of the charged and discharged supercell, respectively, ENa denotes the energy of sodium in metal form, and n is the number of sodium atoms participating in intercalation. The calculated results show that the insertion voltage of r-NiHCF is 2.95 V, 70 mV higher than that of c-NiHCF (2.88 V). Although the measured potential values are based on those materials with a certain concentration of defects, the fact that rhombohedral NiHCF has higher working voltage than cubic phase can be qualitatively consolidated.

Figure 2. Cyclic voltammetry (CV) curves of r-NiHCF and c-NiHCF using 1 M Na2SO4 aqueous electrolyte with the scan rate of 0.1 mV s-1. In the optimized discharged supercell of c-NiHCF (Fig. 3a), Na+ ions reside exactly at the center of the interstitial channels due to the high symmetry of the lattice, resulting a sharp Na-N distance peak (~3.6 Å) in the radial distribution function (RDF) shown in Fig. S3a. In contrast, for the case of r-NiHCF (Fig. 3b), Na+ ions prefer to stay asymmetrically at the N-coordinated corners with much closer Na-N distances (~2.5 Å, Fig. S3a). Such a configuration leads to slightly increased C-N and Ni-N distances by 0.03 Å and 0.1 Å, respectively (Fig. S3b-c). The charge density difference maps between charged and discharged lattices (Fig. 3a and 3b) show that, the intercalation of Na+ causes very slight change for c-NiHCF, however, leads to dramatic electron polarization particularly for the adjacent N atoms, resulting in charge redistribution



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between neighboring Ni2+ and Fe2+ ions. In fact, Bader analysis (Table 1) clearly indicates that, either at charged or discharged state, the iron ions in r-NiHCF are more positive than those in c-NiHCF, indicating the stronger Fe-C interaction in r-NiHCF (Fig. S3d) resulted from the enhanced crystal field. Moreover, the lower electron density of iron ions in the r-NiHCF should benefit gaining electrons during discharge process, leading to improvement on the discharge potential.

Table 1. The charge of Fe in r-NiHCF and c-NiHCF. Charges of Fe

c-NiHCF

r-NiHCF

charged

0.875

0.894

discharged

0.703

0.817

Figure 3. The differential charge density analysis of c-NiHCF (a) and r-NiHCF (b).

We next examined the electrochemical performances of the r- and c-NiHCF test cells. Figure 4a shows the first charge/discharge profiles of the fresh cells at a current density of 10 mA g-1, in which the corresponding chronoamperograms are embedded. Again, we observe the similar voltage difference as observed in the CV scan and the prediction from DFT calculations. In the ether electrolyte, the average charge/discharge voltage of r-NiHCF is 3.37 V (3.41/3.33 V), which is 120 mV higher than that of c-NiHCF (averages to be 3.25 V (3.30/3.21 V)) and is the highest


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among the reported NiHCFs as the cathode of SIBs (Table S3). The larger gap than in the aqueous electrolyte may be ascribed to the different solvation effect of Na+ in solvents.56

Successively, we conducted rate and cyclic tests for the as-obtained r- and c-NiHCF cathode materials. Owning to its three-dimensional open framework with large interstitial channels, r-NiHCF should allow rapid Na+ transportation kinetics for high-rate performances. Figure 4b shows that, at low current density (< 80 mA g-1), the discharge capacity of r-NiHCF almost level out around 67.0 mAh g-1 without visible drop with respect to the rate and cycles. As the current density increases to 160 mA g-1, 240 mA g-1, 320 mA g-1 and 480 mA g-1, it delivers the available discharge capacity of 65.3 mAh g-1, 63.1 mAh g-1, 61.2 mAh g-1 and 58.9 mAh g-1, respectively, and recovers to 67.0 mAh g-1 as the current density shifts back to 10 mA g-1, indicating superior rate capability and capacity reversibility, which is comparable to previously reported best performances of c-NiHCF.42 In contrast, the rate performance of our c-NiHCF is considerably poorer (Fig. S4a). As the current density increases, the available discharge capacity drops rapidly and delivers only 24.0 mAh g-1 at 480 mA g-1. In principle, the large and regular Na+ diffusion channels of c-NiHCF framework should also allow rapid Na+ diffusion for high-rate charge/discharge. The relatively poor rate capability of c-NiHCF can be rationally assigned to the larger particle sizes (10 times of r-NiHCF), which would lead to lower conductivity and longer Na+ diffusion pathways. Both of r- and c-NiHCF encounter ignorable capacity decay in 100 cycles at a low current density of 10 mA g-1 (Fig. S4b). Figure 4c shows that, at a current density of 80 mA g-1, r-NiHCF could retain 96% of its initial capacity (66.8 mAh g-1) after 200 cycles. At a high current density of 480 mA g-1, r-NiHCF hosts an initial discharge capacity of 58.9 mAh g-1 and decays only 0.05 mAh g-1 per cycle. For both rates, the Coulombic efficiencies are all above 95%.



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Figure 4. Galvanostatic charge/discharge voltage profiles of r-NiHCF and c-NiHCF (chronoamperograms were embedded) (a); rate capability of r-NiHCF (b); cycle performance of r-NiHCF under current density of 80 mA g-1 and 480 mA g-1 (c).

It has been reported that the rhombohedral PBA will encounter a rhombohedral-cubic phase transition during electrochemical redox.28, 37 However, either the CV scans in aqueous electrolyte or the rate/cycle tests in organic electrolyte suggest that the as-prepared r-NiHCF cathode material hosts excellent electrochemical stability, implying the rhombohedral-cubic phase transition of the as prepared r-NiHCF is largely inhibited. In fact, our ex-situ XRD patterns of the r-NiHCF at different depth


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of charge/discharge (Fig. 5) show that, at either full discharge or full charge state, the cathode material exhibits typical rhombohedral diffraction features, despite negligible shift of the peaks at 24.1°, 24.9°, 38.6° and 39.7°. No clue shows the rhombohedral-cubic phase transition, suggesting visible structural stability when serving as the cathode material of SIBs. Our DFT calculations also suggest that, the volume variation before and after Na+ intercalation is only 0.72%, which is comparable to the “zero-strain” c-NiHCF as reported by Guo et al.41 Such inferior lattice

changes

bring

negligible

stress

variation

within

the

Na+

intercalation/deintercalation processes, which, of course, will minimize the structural instability and supply excellent cyclic stability.

Figure 5. Ex-situ XRD patterns of r-NiHCF (NaxNiHCF) cathodes at various charge and discharge states (1: x = 0.93, 2: 0.65, 3: 0.44, 4: 0, 5: 0.19, 6: 0.67). Conclusion Developing stable cathode materials with high energy density for room-temperature sodium-ion batteries represents one of the most focal desires for large-scale stationary energy storage technologies. Herein, we report a low-cost rhombohedral phase nickel hexacyanoferrate (NiHCF) which exhibits high operation voltage than that of cubic phase NiHCF. Using the refined structure, the DFT+U calculations suggest that the



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perfect lattice of r- and c-NiHCF gives a discharge voltage of 2.95 V and 2.88 V, respectively. The asymmetric residence of Na+ ions in the rhombohedral framework in conjugation with the low charge density at the Fe2+ sites are proven to be closely correlated with its high working voltage. In aqueous and ether-based electrolytes, the voltage of r-NiHCF is 100 mV and 120 mV higher than that of c-NiHCF, respectively. In particular, a specific capacity of 66.8 mAh g-1, approximating to the theoretical capacity of r-NiHCF, can be implemented at the current density of 80 mA g-1, which can retains 96% after 200 charge/discharge cycles. The excellent cathode stability of r-NiHCF can be assigned to the absence of rhombohedral-cubic phase transition and negligible volume variation during electrochemical redox, as demonstrated by the ex-situ XRD patterns at different depth of charge/discharge and the DFT calculations, respectively. Our results figure out the mechanism of the high operation voltage of r-NiHCF and may provide an applicable cathode candidate for developing high performance room-temperature SIBs.

Supporting Information Electronic Supplementary Information (ESI) available: ICP-OES and Elemental analysis results, Raman curves, XRD indexing curves, RDF of r- and c-NiHCF, Charge/discharge voltages of cubic NiHCF from references, Rate and Cycle performance of c-NiHCF, are included.

Acknowledgements This work was supported by the Natural Science Foundation of Hubei (2013CFB413) and the Fundamental Research Founds for National University, China University of Geosciences Wuhan (Innovative Team, Grant CUG120115; “Yaolan” plan, Grant CUGL150414 and CUGL140413). Support from the National Natural Science Foundation of China (No. 21203169) is also gratefully acknowledged.


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Na1+xMnFe(CN)6


Cathode

for


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