Deintercalation in Metal-Rich Cu1.8S

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Unusual Na+ Ion Intercalation/Deintercalation in Metal-Rich Cu1.8S for Na-ion Batteries Hyunjung Park, Jiseok Kwon, Heechae Choi, Donghyeok Shin, Taeseup Song, and Xiong Wen (David) Lou ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b00118 • Publication Date (Web): 05 Mar 2018 Downloaded from http://pubs.acs.org on March 6, 2018

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Unusual Na+ Ion Intercalation/Deintercalation in Metal-Rich Cu1.8S for Na-ion Batteries

Hyunjung Parka, b, Jiseok Kwonb, Heechae Choic,d, Donghyeok Shinb, Taeseup Song*b, and Xiong Wen (David) Lou*a

a

School of Chemical and Biomedical Engineering, Nanyang Technological University, 62 Nanyang

Drive, Singapore 637459, Singapore b

c

Department of Energy Engineering, Hanyang University, Seoul 133-791, Korea

Center for Computational Science, Korea Institute of Science and Technology, Seoul 136-791,

Karea d

Virtual Lab Inc., Hwarangno 14-gil 5, Seongbuk-gu, Seoul, 02792, Korea

*Corresponding author E-mail: [email protected] (Taeseup Song), [email protected] (Xion g Wen (David) Lou)

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Abstract A Key issue on Na-ion batteries is to develop active materials with stable electrochemical reversibility through understanding their sodium storage mechanisms. We report a sodium storage mechanism and properties of a new anode material, digenite Cu1.8S, based on its crystallographic study. It is revealed that copper sulfides (CuxS) can have metal–rich formulas (x ≥ 1.6), due to the unique oxidation state of +1 found in group 11 elements. These phases enable the unit cell consists of all strong Cu–S bonds and no direct S–S bonds which are vulnerable to external stress/strain that could result in bond cleavage as well as decomposition. Because of its structural rigidness, the Cu1.8S shows an intercalation/deintercalation reaction mechanism even in a low potential window of 0.1–2.2 V versus Na/Na+ without irreversible phase transformation that most of metal sulfides experience through a conversion reaction mechanism. It uptakes in average 1.4 Na+ ions per unit cell (~250 mAh g-1) and exhibits ~100 % retention over 1000 cycles at 2C in a tuned voltage range of 0.5–2.2 V through an overall solid solution reaction with negligible phase separation.

Keywords: metal sulfide, digenite Cu1.8S, intercalation, anode, sodium ion batteries

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Na-ion batteries (NIBs) have recently attracted great attentions as the alternative to Li-ion batteries for the large-scale energy storage systems (ESS) because of abundance and the low cost of Na.1,2 However, Na+ ion has the larger cation radius of 1.06 Å and the heavier atomic weight of 23 g mol-1 than 0.76 Å and 6.9 g mol-1 of Li+ ion, which raises critical issues on sluggish Na+ ion kinetics as well as the large migration barrier energy that often result in poor electrochemical reversibility of the electrode materials, especially transition metal oxides and carbonaceous materials (graphite, hard/soft

carbons).3,4

Due

to

a

lack

of

promising

active

materials

with

the

intercalation/deintercalation mechanism and structural integrity, recent research on anodes for NIBs has been mainly focused on conversion and/or alloying materials. In this respect, development of new materials based on intercalation/deintercalation reaction mechanism remains challenging. Metal sulfides (MexSy, Me = Co, Cu, Fe, Mo, Ni, Sn, W, V, etc.) have been intensively explored as negative electrode materials for rechargeable batteries because of abundance, low cost, and environmental friendliness of sulfur.2, 5 According to previous reports, however, they suffer from electrochemical irreversibility due to following reasons; (1) loss of structural integrity through irreversible phase transformation and the conversion reaction, (2) incompatibility with some electrolytes, (3) decomposition of the active material owing to the dissolution of sulfur.6-9 To overcome those issues, many studies have been performed on nanostructuring, size controlling, composite preparation with carbonaceous materials such as graphene,10-13 (reduced) graphene oxide,14-16 multi-wall carbon nanotubes,17 etc..18-24 Although those strategies on metal sulfides could be effective to address issues mentioned above, a breakthrough can be also achieved by fundamental studies on material’s crystallography (i.e., unit cell, ionic channel, vacancy, bond, etc.) and an optimization of cell conditions.5 For example, pyrite FeS2 is shown to exhibit poor cycle performance due to the mixed intercalation/deintercalation and conversion reactions at the voltage window above 1.0 V versus Li/Li+.25,26 In contrast, FeS2 for NIB exhibited a capacity of ~ 200 mAh g-1 and long-term cycle life in an ether–based electrolyte in the voltage range of 0.8–3.0 V versus 3


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Na/Na+ through the intercalation/deintercalation reaction mechanism.22 Meanwhile, one of the obstacles on such intercalation metal sulfide compounds is the high working potential window above 0.8 V versus Na/Na+ to avoid their irreversible phase transformation and decomposition in the lower voltage range, which is assigned to their structural instability. Therefore, much effort should be dedicated to the development of intercalation/deintercalation metal sulfides with good electrochemical reversibility in a broader potential range. Although various studies on metal sulfides have been conducted for rechargeable batteries, copper sulfides (CuxS) have rarely been investigated despite abundance and low cost of Cu sources (Table S1). Their reaction mechanisms and electrochemical properties have not been clearly understood up to date.27 There exist various polymorphs of CuxS (0.5 ≤ x ≤ 2) in which the location of Cu and S atoms are changed as a function of x value. Prior to the evaluation, it is important to sort out copper sulfides with favorable physicochemical properties by fundamental studies. Among them, Cu1.8S, one of the Cu-rich sulfides (x > 1.5), is considered as a promising candidate for the following reasons. First, chalcocite (Cu2S),28 djurleite (Cu1.94S),29 digenite (Cu1.8S),30,31 and anilite (Cu1.75S)32 were experimentally proven to be stable compounds from their crystallographic studies, refinement analyses, electrochemical experiment, etc. which could well serve as the active materials for Na-ion batteries. Second, they have narrow band gaps in the range of 1.1–1.5 eV as a p-type semiconductor. Especially, the Cu1.8S phase has the smaller bandgap of 1.2 eV than others (Table S2),33 which can be a better electrical conductor.34 Third, the Cu1.8S phase has the cubic structure with the fcc close-packed array of sulfur atoms and copper atoms in interstitial sites. As such, it has the simplest XRD pattern and is suitable for efficient investigation of phase transitions upon cycling (Figure S1). All these merits have motivated us to study the digenite Cu1.8S. Here we report digenite Cu1.8S with a hollow octahedral structure and its sodium storage properties. Based on the crystallographic study of the Cu1.8S, interstitial sites for Na+ ions are determined that play an essential role on structural stability and long-term cycle life. To optimize 4


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sodium storage properties, two important parameters such as the electrolyte and the voltage window are thoroughly investigated. First, two different carbonate–based and ether–based electrolytes are tested to find a better compatibility with the Cu1.8S. Second, electrochemical performances are evaluated within a wide potential window of 0.1–2.2 V versus Na/Na+. A detailed reaction mechanism is elucidated through ex-situ x-ray Diffraction (XRD) and transmission electron microscopy (TEM) studies. Finally, we demonstrate a great potential of Cu1.8S by proposing reasonable factors on the structural stability confirmed by crystallography, ex-situ XRD/TEM, and density functional theory (DFT) calculation.

Results and Discussions Morphology and phase evolution from Cu2O octahedra to Cu1.8S hollow octahedra To synthesize the target Cu1.8S phase, we designed a top-down approach through a combination of a post-calcination and a phase transition from the Cu1.6S phase. The synthesis of Cu1.6S with hollow octadedral structure was based on a previous report.35 Figure 1a and b show schematic illustrations and scanning electron microscope (SEM) images of Cu2O as the template, Cu1.6S as the intermediate product, and Cu1.8S as the final product. 8-facet Cu2O octahedral particles with a smooth surface and uniform size of ~1 µm were synthesized by a simple precipitation and a reduction of copper hydroxides. [email protected] core-shell octahedra were then prepared with Na2S through the kirkendal effect. The Cu1.6S hollow octahedral was obtained by etching the core Cu2O in the ammonium solution. After post-calcination at 300 oC, the Cu1.6S hollow octahedra were transformed to Cu1.8S hollow octahedra via partial evaporation of sulfur, while they maintain original dimension and morpholgy. Phase and crystallinity of each sample were characterized by XRD analysis as shown in Figure 1c. Each peak pattern was indexed with the references of cubic Cu2O (International Centre for Diffraction Data (ICDD) No. 034-1354), geerite Cu1.6S (ICDD No. 033-0491), and digenite Cu1.8S (ICDD No. 056-1256). For Cu1.6S hollow octahedra, the 5


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thermogravimetry (TG) and differential thermal analysis (DTA) graph in Figure 1d shows a weight loss of ~3 % and an exothermic reaction taking place in the temperature range of 300–400 oC because of a large amount of sulfur evaporation, which cause severe recrystallization and a collapse of the original shape as shown in Figure 1e. A XRD study revealed that the particles are assumed to be a composite with two different Cu–rich phases, Cu1.92S and Cu1.94S (Figure S2). Crystallography and Microstructure characterizations of Cu1.8S hollow octahedra To investigate crystallography of as-prepared Cu1.8S hollow octahedra, a Rietveld refinement experiment was performed in Figure 2a. Refined lattice constants obtained under a good Gof value are a = b = c = 5.5657 Å (V = 172.41 Å3) that is almost identical with those of the reference (Table S2). Detailed lattice parameters, atomic coordinates and site occupancies were summarized in Table S3 and S4. The inset shows a 1 × 1 × 1 unit cell scheme of the Cu7.2S4 (= Cu1.8S) along [001] projection through the first-principles method.33 In this cubic unit cell, 7.2 Cu atoms randomly and statistically occupy different Wyckoff positions at and near tetrahedral interstices including 8c (1/4, 1/4, 1/4), 96k (x, x, z), and 192l (x, y, z). Otherwise, 4 S atoms take 4b (1/2, 1/2, 1/2) as fcc lattices. On the basis of the crystallographic information, the number of ideal intercalation sites was determined. There exist 8 tetrahedral sites (Figure 2b) and 13 octahedral sites (Figure 2c). If the sites are fully occupied by foreign atoms, the total number of atoms per unit cell is equal to 12 (Figure 2d). For Cu7.2S4, the number of 4.8 is obtained for empty sites by subtracting 7.2 (Cu atoms) from 12, which means 4.8 Na atoms can be ideally intercalated in the cubic Cu7.2S4 cell without a significant lattice parameter change (for the Cu1.8S, 1.2 Na atoms). Because Cu atoms are distributed near the tetrahedral sites, Na atoms likely favor the 4b Wyckoff octahedral sites upon sodiation. In addition, the ionic radii for Na+ and S2- are 0.106 Å and 0.182 Å, respectively. The ratio between cationic and anionic radii (rC/rA) is 0.58 that gives a coordination number 6 (0.414 < rC/rA < 0.732). This indicates that six sulfur atoms surround each Na atom, and it makes an

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octahedral arrangement. Hence, those empty octahedral sites in the Cu1.8S unit cell can be a great advantage as the anode material for Na ion uptake. A microstructure of the Cu1.8S sample was characterized by TEM. Figure 3a shows a TEM image of the individual Cu1.8S particle. It is ~ 800 nm in length and ~ 650 nm in width, and the wall thickness is ~ 100 nm. The contrast difference between the core and the shell clearly confirms the hollow characteristic. A high-resolution TEM image in Figure 3b exhibits a lattice fringe of (111) plane with a d-spacing of 3.2 Å. Selected-area electron diffraction (SAED) patterns in Figure 3c confirm polycrystalline nature of Cu1.8S octahedra. Figure 3d shows scanning transmission electron microscope (STEM) and elemental mapping images. The copper (green) and sulfur (red) are homogeneously distributed throughout the overall region. Energy-dispersive X-ray spectroscopy (EDX) analysis in TEM was performed to investigate the composition (Figure S3). The atomic ratio is determined to be 1.83 (Cu) to 1 (S) that is consistent with the formula of the Cu1.8S. The other samples (Cu2O, [email protected] and Cu1.6S) were further observed by TEM (Figure S4). As shown in the images, all the samples have the same octahedal structure but with difference in the contrast of core parts. To investigate chemical and electronic states, X-ray photoelectron spectroscopy (XPS) was explored. Figure 3e shows two strong peaks at ~932.7 and ~952.7 eV assigned to Cu(I) 2p3/2 and Cu(I) 2p1/2,36 respectively. Especially, the two main peaks have small shoulder peaks at higher binding energies of ~934.7 and 954.6 eV due to a coexistence of Cu(II) 2p3/2 and Cu(II) 2p1/2. A broad peak that consists of S 2p3/2 and S 2p1/2 was also observed (Figure 3f). It reveals that the Cu1.8S phase is capable of multiple redox couples of Cu0/+1 and Cu+1/+2 when Faradaic reaction occurs such as charging/discharging. A ratio of Cu+1 to Cu+2 is determined to be around 0.7 to 0.3 by peak splitting and integrating areas. From the XPS study, a theoretical capacity of Cu1.8S is calculated on the basis of the ratio and the following equation (1).37

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Q = . × (mAh g )

(1)

Where n is the number of electron transfer per formula unit of the electrode material, F is the Faraday constant, and Mw is the molecular weight of the electrode material. In this work, n is 2.3 (e-), Mw is 146.25 g/mol, and F is 96485 sA/mol. Hence, the theoretical capacity is ~420 mAh g-1. Reversibility of the Cu1.8S hollow octahedra anode in carbonate–/ether–based electrolytes One of the key factors governing the reversibility of the active materials is the choice of a suitable electrolyte.4,

38

In this regard, two different electrolytes were prepared for the Cu1.8S

material; 1.0 M NaCF3SO3 in ethylene carbonate and dimethyl carbonate (EC/DEC, 1:1(v/v)) and 1.0 M NaCF3SO3 in diethylene glycol dimethyl ether (Diglyme). Cyclic voltammetry (CV) was used to investigate the electrochemical reaction and the reversibility. With the 1.0 M NaCF3SO3 EC/DEC electrolyte, the Cu1.8S electrode shows extraordinary cathodic peaks in the potential window below 1.0 V versus Na/Na+ (Figure S5a) that correspond to the unstable charge curves (Figure S5b). This might be due to decomposition of the electrolyte and formation of a solid electrolyte interface (SEI) layer generally observed in carbonate–based electrolytes when they are evaluated below ~1.0 V versus Li/Li+ and Na/Na+. That is supported by the larger discharge capacity observed than the theoretical value and TEM observation (Figure S6a). To restrain from the side reactions of the electrolyte, the CV test was performed in the higher potential window of 1.0 to 2.5 V versus Na/Na+ (Figure S5c). It shows symmetric and reversible cathodic/anodic peaks, but the capacity obtained is too small which is only 32 % of the theoretical one (Figure S5d). In contrast, with the 1.0 M NaCF3SO3 in diglyme electrolyte, CV curves were reversible and symmetrical even in the potential window of 0.1–2.5 V versus Na/Na+ (Figure S5e) that results in a high coulombic efficiency of ~ 99 % and a large capacity of ~403 mAh g-1 (96 % of the theoretical capacity) (Figure S5f). This can be possible owing to the higher stability of the ether-based electrolyte at the reduction potential, which leads to prohibition of the electrolyte decomposition 8


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and the SEI layer formation (Figure S6b).39,40 Hence, it reaches a conclusion that the ether-based diglyme electrolyte has a good compatibility with Cu1.8S. Optimization of electrochemical performances in the diglyme electrolyte The operating potential window should be also considered because the reaction mechanisms of active materials strongly depend on it. Figure 4a to c show voltage profiles of the Cu1.8S hollow octahedra anodes in three different potential windows of 0.1/0.5/0.7 to 2.2 V versus Na/Na+. They give initial discharge/charge capacities of 413/383, 311/272, and 226/230 mAh g-1 at 0.05C, respectively. The voltage profiles upon cycling differ as a function of the potential window. In the voltage range of 0.1–2.2 V, the curves continuously changed, and plateaus disappeared from 0.05 to 0.2C during 50 cycles (Figure 4a). In contrast, the profiles in potential windows of 0.5–2.2 V and 0.7–2.2 V were slightly changed, and the capacities increased in subsequent cycles (Figure 4b and c). Figure 4d shows cycle performances and Coulombic efficiencies of the three electrodes at a current rate of 0.2C over 50 cycles. The Cu1.8S hollow octahedra anode shows ~75 % retention in the potential window of 0.1–2.2 V. It should be noted that the cycle performance in that potential range is much better than most of the bare metal sulfides without carbonaceous materials (CoS2, Co9S8, MoS2, Ni3S2, SnS2, WS2, etc.) which normally show drastic capacity fading and the retention below ~30 % (< 100 mAh g-1) within 50 cycles.10,11,17,18,21,24,41,42 The better cycle life of Cu1.8S can be due to enhanced structural stability originated from the Cu–S bonds discussed in a later section. Meanwhile, cycle performances were remarkably enhanced in the higher potential window of 0.5– 2.2 V and 0.7–2.2 V. Particularly, it shows ~100 % Coulombic efficiencies and >100 % capacity retentions. From the results, long-term cycle performance was evaluated at a higher current density of 2C in the optimal potential window of 0.5–2.2 V. As shown in Figure 4e, with a slight capacity increase during the initial a few tens of cycles, the anode gives a capacity of ~250 mAh g-1 and superior cycle stability over 100 % retention after 1000 cycles. It is interesting that this superior cycle performance was achieved in the capacity range of 200–250 mAh g-1 corresponding to 1.1–1.4 9


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Na atoms in the Cu1.8S unit cell that are very close to the ideal value of 1.2 calculated from the interstices (Figure 2), which strongly implies its intercalation/deintercalation reaction mechanism. As a result, the Cu1.8S hollow octahedra anode exceeds various anode materials previously reported as intercalation materials such as metal oxides,43-50 metal sulfides,22,51,52 carbonaceous materials,40, 53-57

etc..58 (Table S5). Electrochemical impedance spectroscopy (EIS) was further used to see any

change in physicochemical properties upon cycling. With the increasing cycle number, EIS spectra in Figure 4f show a trend of decrease in semicircles and charge transfer resistance that might be correlated to defect chemistry discussed in detail in a later section. Detailed electrochemical reaction mechanism of Cu1.8S hollow octahedra To understand the sodium storage mechanism in detail, an ex-situ XRD study was carried out in terms of different state of charge (SOC) and depth of discharge (DOD) corresponding I–VII points in the 1st cycle galvanostatic profile (Figure 5a). As shown in Figure 5b, after sodiation of ~1.5 Na ions (discharging: I → III), the original phase maintained with a negligible peak shift to higher angles ((111): 27.7o → 27.8o and (220): 46.1o → 46.4o), and a small peak of a new phase was found at ~38o. After desodiation (charging: III → V), it recovers the original phase, but the peaks were slightly broadened. To further investigate the reaction mechanism, an ex-situ TEM observation was also performed. A low-magnification TEM image in Figure 5c confirms no morphological change of the Cu1.8S hollow octahedron after the first cycle. Figure 4d to f show HRTEM images and SAED patterns of I, III, and V samples. After discharging to 0.5 V (Figure 5e), the dominant (111), (220), and (220) original planes were observed in an overall area, but the new and minor phase with d-spacing of 2.3 Å was found in confined area as a shape of a scattered island, which is also confirmed by small dots, not the ring pattern in SAED image. The longer plateau during charging might be due to the phase separation between the original phase and new phases and disappearance of the new phases. At the stage V after charging up to 2.2 V (Figure 5f), the same original lattices and patterns were observed without the new phase. As a result, it was proved that 10


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the Cu1.8S phase experiences the intercalation/deintercalation reaction upon cycling, more specifically, a combination of the dominant solid solution reaction because of the sodium intercalatioin into the interstices and the negligible phase separation reaction originated from the uncertain phase evolution, which agrees with the coexistence of the overall smooth slope and the short plateaus in the voltage profile, respectively.59 It is worth noting that Na+ intercalation into the interstitial sites is less likely to cause a severe change in lattice parameters and initial XRD patterns in the ideal value of ~1.2 because the diffraction patterns of possible Na–intecalated compounds including CaF2, NaCl and Li3Bi structures (Figure 2 b to d) are all identical. In contrast, when more Na+ ions were intercalated above 1.2 by discharging down to 0.1 V (stage VI), the main (111) peak largely moved to a higher angle from 27.7o to 28.6o, and the new peaks at 39o became more predominant than the original (Figure 5b) that results in more changes of cell parameters as well as a phase transformation also confirmed by a SAED image in which the original patterns (green lines) are obscure, but the new pattern (orange line) is prominent (Figure 5g). It is hard to confirm the new phase, NaxCuySz, because of a lack of crystallography and diffraction data on sodium copper sulfides. Only two phases were found in ICDD such as NaCu2S2 (ICDD No. 42-1366) and Na3Cu4S4 (ICDD No. 33-0488). Interestingly, matching those to the VI patterns somehow shows resemblances (Figure S7). Those phases have the same primitive cell with one or two larger lattice constants than those of Cu1.8S (Table S6). This implies that more Na+ intercalation beyond the ideal value could transform the cubic cell of Cu1.8S to a primitive cell with modified lattice parameters such as axial lengths and angles like the unit cell of Na3Cu4S4 (Figure S8). In this unit cell, every Na+ ion coordinates with 6 sulfur anions (i.e., CuS6 octahedron), which agrees well with our assumption that Na+ ions rather prefer to occupy octahedral sites in the Cu1.8S cell. However, more studies are needed to clarify the crystallography of the new sodium cupper sulfide (NaxCu1.8S) in future. It is deduced from this study that the dependancy of cyclability on the operating voltages can be correlated to the unit cell change generally found on intercalation/deintercalation materials due 11


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to the confined number of interstitial sites for foreign ions such as Li+ and Na+ ions. Despite the lattice constant change, Cu1.8S retained its original phase and morphology (Figure S9) even in this lower potential window of 0.1–2.2 V versus Na/Na+ also confirmed by ex-situ TEM study (Figure 5h), which indicates that it surprisingly does not experience the conversion and/or alloying reactions as do most of the metal sulfides. The reaction of Cu1.8S is quite similar with that of vanadium pentoxide with several short plateaus (reaction: V2O5 → NaV2O5 + Na2V2O5 (discharge) → NaV2O5 + V2O5 (charge)).60 To further clarify the reaction mechanism, an ex-situ XPS study was carried out (Figure S10). With discharging to 0.5V and 0.1V, the redox couple of Cu+1/+2 turns into Cu0/+1. This is another evidence for the intercalation reaction of Cu1.8S. Therefore, the overall reaction of Na/Cu1.8S cell can be described by the following equation: xNa + Cu1.8S ⇆ NaxCu1.8S (x < 2.3 within 0.1–2.2 V vs. Na/Na+)

(2)

Defect evolution and its potential effect on sodium storage properties We further carried out ex-situ XRD measurements on the samples after more prolonged cycles in the voltage window of 0.5–2.2 V versus Na/Na+ (Figure 6a). As shown in the graph, the main peaks especially those corresponding to (111), (200) and (220) planes are well maintained, but more broadened upon cycling. In crystalline materials, the broadening of the XRD patterns is related to the development of defects in materials supported by well-established theories.61 In this regard, we introduced theories to explain the peak broadening phenomenon, and microstrain of the Cu1.8S hollow octahedra in terms of cycle numbers was calculated on the basis of Willson–Stokes definitions (3) and (4) and Williamson–Hall equations (5) and (6). η = β cotθ %

ε = &

(3) (4)

where η is the apparent strain, β is the integral breadth of the reflection (A/I0, A: the peak area, I0 is the height of the line profile), ε is the strain, and θ is the Bragg angle at [h k l] reflection. 12


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β' cosθ = ε=

0123 &

)* +

+ 4εsinθ

cotθ, if

)* +

=0

(5) (6)

where βhkl is the full width at half maximum (FWHM), K is the Scherrer constant related to particle shape and the arbitrary value, and λ is the wavelength of the radiation (1.54056 Å for CuKα radiation). In general, x-ray line broadening is affected by size broadening and strain broadening. The size broadening can be developed through severe deformation by ball milling, rolling, bending, etc.. But the size broadening term can be excluded in some cases like rechargeable batteries under the condition that the active material experiences the intercalation/deintercalation reaction that is usually accompanied with no particle size and shape changes during cycling. In this respect, Willson and Stoke definition (W–S) well fit this assumption because it only deals with strain broadening. Another equation, Williamson–Hall method (W–H) consists of size term (Kλ/D) and strain term (4εsinθ), but the size term was excluded for the same reason. For determination of the microstrain of the Cu1.8S hollow octahedra, a profile fitting was performed on the most dominant (220) peaks obtained from ex-situ XRD samples (Figure 6b). Based on the fitting data, microstrains of the samples were determined as a function of cycles, and the calculated values between W–S and W–H methods are consistent (Figure 6c and Table S7). It was revealed that the microstrain increased from 0.2 % to 0.9 %, but a rate of change in the stress-strain plot decreased and saturated after 50 cycles (slopes in Figure 6c). In addition, the peaks corresponding to the (220) plane shifted to the higher 2θ degree (46.1o → 46.7o) that indicates the lattice distance decreased from 1.91 Å to 1.88 Å posibly due to the strain formed by Na+ intercalation/deintercalation. To visualize the evolution of the defects, an inverse fast Fourier transformation (IFFT) technique was used (Figure 6d to g). IFFT images clearly show the development of the defects in the cycled Cu1.8S hollow octahedra. In contrast to the pristine with well–ordered lattices and no defect (Figure 6e), two kinds of crystallographic defects such as an edge dislocation and a screw dislocation were found after one 13


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cycle (Figure 6g). It is expected that this microstructural change can make a direct impact on the sodium storage properties by providing additional sites for Na+ ions and better ionic diffusivity, which is correlated to the slight change in the voltage profile as well as the increase in the capacity of Cu1.8S upon a few tens of early cycles in accordance with previous reports on the capacity increase in graphite and metal oxide with defects deliberately formed by deformation.62,63 Discussion on phase and structural stability of Cu1.8S From these studies, one question is raised on how digenite Cu1.8S can show the better stability than the other metal sulfides. We propose some reasonable factors on it in this section. First, the oxidation states of copper and extraordinary formulas of copper sulfides; copper has common oxidation states of +1 and +2, and especially +1 state is only found in group 11 due to the possible ionization by losing one s–orbital electron on top of the completely filled d–electron shell (Table S8). These oxidation states of copper enable formation of copper sulfides with a wide range of formulas, CuxSy (1.0 ≤ x ≤ 2.0, y = 1, 2) as shown in Table 1. Especially, metal–rich phases (MexS, 1.6 ≤ x ≤ 2.0) are only found in group 11 metal sulfides highlighted in red for copper sulfides. Those compounds have isolated sulfur anions in hcp or fcc lattices and copper cations in trigonal and/or tetrahedral interstitial sites, and they consist of all Cu–S bonds without a direct S–S bond in unit cells. Figure S11 shows unit cell structures of covellite CuS and pyrite CuS2 (pyrite MeS2, Me = Fe, Co, Ni, etc.). Cell parameters were listed in Table S9. Digenite Cu1.8S consists of all covalent Cu–S bonds (inset in Figure 2a). In contrast, the total numbers of direct S–S bonds for covellite CuS and pyrite CuS2 are 4 (Figure S11a) and 13 (for CuS2, 4 bonds (#1-4) in Figure S11b, 8 same symmetry-equivalent S–S bonds in other sides of the square, and 1 bond (#13) at the center in Figure S11c. Coordinates of the sulfur atoms for each bond were summarized (Table S10). According to our investigation, the formulas and the bonds of metal sulfides could be a key for the phase and structure rigidness for the electrode material. Scaini et al. have reported the fracture of pyrite FeS2 can be developed by the rupture of the weaker direct S–S bond than the stronger Fe–S 14


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bond according to the analysis of the bond energy.64 Liberato et al. have studied the structural and compositional transformations of covellite (CuS) and djurleite (Cu1.94S) via cation exchange reactions. The reaction in the presence of a reducing agent as an electron donor triggers the entry of foreign metal cations in the CuS lattice causing the reorganization of the anion frame due to the rupture of S–S bonds while the reaction is not possible in Cu1.94S due to a lack of S–S bonds.65 Those studies strongly imply that the existence of S–S bond can initiate a collapse of its original crystallinity and phase of the metal sulfides. Indeed, considering the major byproduct of metal sulfides through the conversion reaction is alkali metal sulfides (2Me+ + S- → Me2S, Me = Li, Na, K, etc.), it is reasonable to assume that the metal sulfide with fewer sulfur atoms ensures the better structure stability as the anode material. Hence, Cu-rich sulfides (Me/S ≥ 1.6) can have the better structural stability than others. Second, a low possibility of a reaction between copper and sodium; Cu has a very low solubility in Na and even resistant to reaction with Na up to the elevated temperature of 300–400 oC, which results in no natural sodium copper compounds.66 Third, a low valence of copper cations and shielding effect of sulfur anions; it is clear that both cations with lower valence and anions with larger electron density can decrease the migration energy according to the previous report.67 The Cu cation has +1/+2 oxidation state lower than +2/+3 of other metal sulfides, and sulfur anions have a better shielding effect between sodium cations and metal cations than that of oxygen anions which can lead to weaker electrostatic repulsion against Na+ ions during intecalation/deintercalation. DFT calculation was carried out to determine Na+ ion migration energy in the bulk Cu1.8S. It should be noted that all Cu atoms are randomly located in and/or near the tetrahedral sites due to its non–stoichiometry and position change of Cu and S atoms (Figure 7a), the migration energy can slightly vary from site to site that makes it challenging to determine the migration energy in every site. In this respect, the barrier energy was calculated through a pathway between minimum and maximum coordination numbers of Na+ ion with Cu atoms, which has the largest impact on the Na+ ion diffusivity (Figure 7b). An asymmetric energy plot in Figure 7c shows 15


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the formation energy of Na+ migration from one interstice coordinated by 3 Cu atoms (labelled Asite) to another adjacent interstice coordinated by 6 Cu atoms (labelled B-site). The calculated energy barrier of Na+ diffusion is ~0.51 eV/atom that is in a low level compared to those of promising metal oxides previously reported such as Li4Ti5O12 and Na0.66[Li0.22Ti0.78]O2, etc. in the range of 0.4–0.9 eV.45, 68 Therefore, all these results demonstrate that Cu1.8S can be an excellent Na+ ion conductor. Conclusion In summary, digenite Cu1.8S was studied as a new anode material for Na-ion batteries (NIBs). Important parameters including electrolyte and potential window were optimized on a Cu1.8S hollow octahedra anode. Unlike an unstable reaction in the EC/DEC electrolyte, the Cu1.8S hollow octahedra anode exhibited good reversibility and a capacity of 403 mAh g-1 (93 % of the theoretical one) in 1.0 M NaCF3SO3 in the dyglmer. The Cu1.8S hollow octahedra electrode shows a reversible charge capacity of ~250 mAh g-1 and a high Coulombic efficiency of ~100 % over 1000 cycle at a high rate of 2C in a potential range of 0.5–2.2 V versus Na/Na+ without irreversible phase transformation and structural collapse. Although the hollow structure of Cu1.8S can enhance the electrochemical properties due to short ionic diffusion length, the phase and structural stabilities are more likely to be originated from its intrinsic properties such as octahedral interstitial sites for Na+ ions as well as all direct and strong Cu–S bonds in the Cu1.8S phase. Transition metal sulfides can be considered attractive intercalation electrode materials from the viewpoint of better ionic conductivity and lower migration energy than those of transition metal oxides as proven by the simulations and calculations. However, this has rarely been realized up to now due to some issues on the irreversible phase transformation. In this respect, this finding of Cu1.8S can give a way to develop new intercalation materials, and enhanced sodium storage properties could be achieved by engineering size and geometry, doping foreign atoms, and developing other metal-rich sulfides such as Cu4FeS for Na-ion batteries. 16


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Methods Synthesis. Cu2O octahedrons were prepared based on a previous report.35 Synthetic procedures for [email protected] core-shell structure were modified from the previous one. Two separate solutions were prepared for a typical synthesis. As solution A, 0.20 g of Cu2O octahedral particles was put in 5 ml of mixture solution with deionized water and ethanol (2 : 1 (v/v)). To disperse the particles, the solution was sonicated for 10 min. As solution B, 0.4 g of sodium disulfide was dissolved in 15 ml of mixture solution with deionized water and ethanol (2 : 1 (v/v)). Under stirring at 300 rpm, the solution A was added into the solution B. After 10 min, red colored Cu2O power was changed to black colored power. Those particles were washed with deionized water and ethanol in three times by centrifugation. The resultant powder was dried at 70 oC for 12 h in a vacuum oven. To etch out inner Cu2O, [email protected] octahedral particles were immersed in an ammonia solution (~28 %) for 72 h. During the time, the ammonia solution was refreshed every 24 h. After that, the resultant particles were washed with deionized water till the supernatant becomes nearly neutral in pH value. Then the powder was dried at 70 oC for 12 h in a vacuum oven. To obtain the final product of Cu1.8S hollow octahedra and to remove some residual sulfur on the surface, Cu1.6S hollow octahedral particles were annealed at 300 oC in Ar gas (200 sccm) for 30 min by using quartz tube furnace. During the annealing process, the Cu1.6S phase was transformed to the Cu1.8S phase. Characterizations. Morphological and microstructural studies on as-prepared particles were performed using field-emission SEM (JEOL JSM07600F), field-emission TEM (JEOL JEM-2100F), and HRTEM (JEOL JEM-2100F). TEM images were further analyzed by Gatan microscopy suite software version 3 for lattice parameters, SAED patterns, IFFT images, etc. XRD patterns were obtained on as-prepared Cu2O, Cu1.6S, Cu1.8S and Cu1.92S/Cu1.94S samples by X-ray diffraction analyzer (Rigaku D/MAX RINT-2000) in the range of 20 –70 o at a scan rate of 3.0 o per min using CuKα X-ray source. To investigate the phase transition, the weight change was examined by TGDTA (SDT Q 600, Auto-DSCQ20 system). X-ray photoelectron spectrometry (XPS, VG Microtech 17


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ESCA2000) was used for the analysis of the electronic and chemical states of the Cu1.8S hollow octahedra. Rietveld refinement. A XRD measurement was performed on the Cu1.8S hollow octahedra in the range of 5–120 o at a scan rate of 1.0 o per min by Bruker XRD (D80-Advanced, a1 system), and the refinement on lattice parameters including phase, lattice constants, site occupancy, etc. was carried out through the TOPAS software (Bruker). Ex-situ XRD measurement. For ex-situ XRD studies, 2032 type coin cells were assembled and cycled at different cut-off voltages (discharge to 1.0/0.5/0.1 V and charge to 1.5/2.2 V). After cycling, the cells were disassembled to take out the anodes. The samples were immersed in a diethyl ethylene carbonate (DEC, Sigma Aldrich, anhydrous, ≥ 99%) solution to remove any residual electrolyte and byproduct. The electrode films were peeled off from the copper current collector using a polyimide (Kapton) Tape and sealed to prevent them from contact with air and moisture. All those experiments for the sample preparation were carried out in an Ar filled glove box. Then, exsitu XRD patterns were collected at a slower scan rate of 0.4 o per min by X-ray diffraction analyzer (Rigaku D/MAX RINT-2000). Unit cell drawing. The unit cell structures of Cu1.8S, CuS, CuS2, FeS2 and Na3Cu4S4 were drawn by

the 3D visualization program for structural models (VESTA 3). Electrochemical evaluation. For half coin cell tests, anodes were prepared with Cu1.8S hollow

octahedral particles as an active material, acetylene black as a conducting agent, and carboxylic methylcellulose/emulsified styrene-butadiene rubber as binders in a weight ratio of 80 : 15 : 5 (2/3). All the materials were put in a 20 mL vial and thoroughly mixed using a Thinky mixer at a high rotation speed of 2000 rpm for 5 min. As-prepared slurry was casted on a copper foil (T = ~18 µm) using a doctor blade. Before cell fabrication, the anodes were dried in a vacuum oven at 120 oC for 2 h to thoroughly remove residual moisture in the electrode. Coin-type cells (2032 R type) that consist of the anodes as a working electrode, sodium foil as a counter electrode, 1.0 M NaCF3SO3 in 18


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ethylene carbonate and dimethyl carbonate and 1.0 M NaCF3SO3 in diethylene glycol dimethyl ether (≥99 %, Sigma-Aldrich) as two different types of electrolytes, and a polypropylene (PP) film as a separator, were fabricated. Electrochemical performances were obtained by using a battery cycle tester (TOSCAT 3000, Toyo System, Tokyo, Japan). Impedance spectra were measured using impedance analyzer (PARSTAT 2273, Princeton Applied Research). DFT calculation. Density functional theory (DFT) calculations were performed with the generalized gradient approximation (GGA) and Perdew-Burke-Ernzerhof (PBE) parameterization. We used the Vienna ab initio simulation package (VASP) program. Kohn-Sham orbitals were expanded to a cutoff energy of 400.0 eV. 2×2×2 equally spaced k-point grids were used for the Brillouin zone sampling for the unit cell consisting with 90 atoms (Cu58S32), which is used as the supercell for Na-diffusion. To calculate the diffusion energy barrier of Na atom in Cu9S5 crystal, we used nudged elastic band (NEB) method.

ACKNOWLEDEMENTS This work was supported by the Korea Institute of Energy Technology Evaluation and Planning (KETEP) and the Ministry of Trade, Industry, & Energy (MOTIE) of the Republic of Korea through the reseach on Li-ion batteries (No. 20168510050080), Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT (2016R1C1B2007299) and the research fund of Hanyang University (HY-2017).

ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website at http://pubs.acs.org 19


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Prices of pure elements and bulk compounds of metal for metal sulfides (Table S1), Crystallographic data and their band gaps of copper sulfide polymorphs (Table S2), X-ray diffraction patterns of CuxS (Figure S1), SEM image and XRD patterns of CuxS hollow octahedra after post-annealing 400 oC (Figure S2), Refined unit cell parameters of the Cu1.8S (Table S3), Refined atomic coordinate and site occupancy (Table S4), TEM/EDX analysis of the Cu1.8S (Figure S3), TEM study of Cu2O octahedron, [email protected] octahedron, and Cu1.6S hollow octahedron (Figure S4), Cyclic voltammetry spectra and charge/discharge curves in two types of electrolytes (Figure S5), TEM study on the Cu1.8S hollow octahedral anode after one cycle (Figure S6), Electrochemical performances of various types of the anode materials (Table S5), Ex-situ X-ray diffraction study on the sodiated Cu1.8S hollow octahedra anode (Figure S7), Lattice parameters of digenite Cu1.8S and sodium copper sulfides (Table S6), Unit cell structures of Na3Cu4S4 (Figure S8), TEM image of the Cu1.8S hollow octahedral anode after 1 cycle (Figure S9), Ex-situ XPS spectra of Cu 2p (Figure S10), Summary of parameters and calculated microstrain (Table S7), common oxidation states and electron configuration of metals (Table S8), Unit cell parameters of the degenite Cu1.8S, covellite CuS, and pyrite CuS2 (Table S9), Unit cell structures of covellite CuS and pyrite CuS2 with the cubic structure (Figure S11), Coordinate of each sulfur atom for S–S bonds in covellite CuS and pyrite CuS2 (Table S10). AUTHOR INFORMATION Corresponding Author E-mail: [email protected] E-mail: [email protected]

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Figure 1. (a) Schematic illustration of the formation of Cu1.8S hollow octahedron. (b) Corresponding SEM images of as-prepared samples. (c) XRD patterns of as-prepared Cu2O octahedral particles (black line), Cu1.6S hollow octahedral particles (red line), and Cu1.8S hollow octahedral particles (blue line). References are indexed below each pattern. (d) TG-DTA curve of Cu1.6S hollow octahedral particles in N2 gas within the temperature range of 100–600 oC. (e) SEM image of as-prepared Cu1.92S/Cu1.94S composite.

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Figure 2. (a) Rietveld refinement of observed XRD patterns. The inset shows a 1 × 1 × 1 unit cell scheme through [001] projection. Ball and stick models of (b) CaF2 structure with tetrahedral interstitial sites and (c) NaCl structure with octahedral interstitial sites. (d) Space filling model of Li3Bi structure for [Na, Cu]12S4 (= [Na, Cu]3S). All structures have the same Fm-3m space group.

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Figure 3. (a) Low-magnification and (b) High-magnification TEM images. (c) SAED patterns. (d) STEM image and EDX map for Cu element (green) and S element (red). XPS spectra of (e) Cu 2p and (f) S 2p.

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Figure 4. Voltage profiles of the 1st cycle at a current rate of 0.05C and 1st/50th cycles at a current rate of 0.2C in the different voltage windows of (a) 0.1–2.2 V, (b) 0.5–2.2 V, and (c) 0.7–2.2 V versus Na/Na+ (1 C = 420 mA g-1). (d) Cycle performance over 50 cycles at a current rate of 0.2C. (e) Long-term cycle performance at a current rate of 2C in the voltage window of 0.5–2.2 V versus Na/Na+. (f) EIS spectra before and after cycling over 1st/20th/50th. The inset shows an equivalent circuit of the Na/Cu1.8S electrode cell.

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Figure 5. (a) Galvanostatic profile at a current rate of 0.1C in the potential window of 0.5/0.1–2.2 V versus Na/Na+. (b) Ex-situ XRD patterns obtained from the anode at each stage of I (pristine), II (discharge 1.0 V), III (discharge 0.5 V), IV (charge 1.5 V), V (charge 2.2V), VI (discharge 0.1V), and VII (charge 2.2 V). (c) Low-magnification TEM image of a Cu1.8S hollow octahedron after one cycle. (d to h) High-magnification TEM images and corresponding SAED patterns observed from the samples at I, III, V, VI and VII stages.

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Figure 6. (a) Ex-situ XRD patterns of the samples after cycling in the potential range of 0.5–2.2 V versus Na/Na+. (b) Gaussian profile fitting of (220) peaks normalized with minimum values. (c) Cycle number–Strain plot. HRTEM and inverse fast Fourier transformation (IFFT) images of (d, e) pristine and (f, g) after one cycle.

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Table 1 Polymorphs of metal sulfides with different formulas. Note that Metal sulfides, pyrite MeS2 are chemical compounds of Me2+ and S22-. MeS compounds consist of Me2+ and S2.

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Figure 7. (a) 2 × 2 × 2 unit cell scheme of digenite Cu1.8S. (b) Scheme of Na+ ion migration path inside the Cu1.8S unit cell. (c) Calculated energy barrier of Na+ ion diffusion. The grey, yellow and deep blue spheres indicate Cu, S and Na atoms, respectively.

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TOC graphic

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Deintercalation in Metal-Rich Cu1.8S

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