Amorphous Selenium

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Amorphous, Crystalline and Crystalline/Amorphous Selenium Nanowires and Their Different (De)Lithiation Mechanisms Xiaoming Zhou,† Peng Gao,*,† Shuchao Sun,† Di Bao,† Ying Wang,† Xiaobo Li,† Tingting Wu,† Yujin Chen,*,‡ and Piaoping Yang*,† †

Key Laboratory of Superlight Materials and Surface Technology, Ministry of Education, College of Materials Science and Chemical Engineering, Harbin Engineering University, Harbin, Heilongjiang 150001, P. R. China ‡ College of Science, Harbin Engineering University, Harbin, Heilongjiang 150001, P. R. China S Supporting Information *

ABSTRACT: Selenium is applied as the cathode material of lithium/selenium (Li−Se) battery, which has high theoretical volumetric capacity density (3253 mA h·cm−3) and high output voltages. However, it has a low melting point (494 K) and a large sensitivity to thermal treatment, which often result in the phase transition between crystalline Se (trigonal phase) and amorphous Se during the charge/discharge cycles of Li−Se battery as reported in literatures. In order to clarify the different chemical (de)lithiation mechanisms between them in Li−Se battery, in this work large-area amorphous selenium (a-Se) nanowires (NWs) have been successfully prepared first through a facile high-energy ball-milling method. Subsequently the crystalline (c) and crystalline/amorphous (c/a) selenium NWs have also been prepared through annealing the above as-obtained amorphous products, respectively. The affirmative composition and morphology of the as-obtained Se nanostructures have been demonstrated by the XRD, SEM, TEM, HRTEM and Raman spectra measurements. And their specific surface area and pore size distribution have also been analyzed by BET measurements. Finally, it is proved that the as-obtained NWs used as the cathode material of Li−Se battery displayed different chemical reaction processes with Li+ and the related various storage capacities (a-Se: 755 mAh·g−1; c/a-Se: 705 mAh·g−1; c-Se: 250 mAh·g−1). This work has helped us to better understand and correlate the formation of intermediate phases with the electrochemical performance of Li−Se cells and shines new light on how to improve the cell performance by turning the phase of Se.

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INTRODUCTION High-energy and low-cost lithium-ion batteries are the most promising candidates for emerging electric vehicles and largescale renewable energy storage. However, on current technology stage, the energy density of lithium-ion batteries is mainly limited by its cathode material.1−3 Lithium/selenium battery, using selenium as cathode material, shows the following superior advantages: (a) better discharging rate and cycling performance due to its high electrical conductivity (1 × 10−3 S· m−1); (b) higher output voltages (approximately 0.5 V higher than that for Li/S) and the corresponding higher energy densities. Although Se has a lower theoretical gravimetric capacity (675 mAh·g−1), its higher theoretical volumetric capacity density (3253 mAh·cm−3) offsets the deficiency. It is known that for applications in portable devices and hybrid electric vehicles (HEVs), volumetric energy density is more important than gravimetric energy density because of the limited battery packing space.4 However, it should be noticed that Se has a low melting point (494 K, decreased to much lower value in its nanostructure) and a large sensitivity to thermal treatment,5−7 which often result in the phase transition between crystalline Se (trigonal phase) and amorphous Se (a© XXXX American Chemical Society

Se) during the charge/discharge cycles of Li−Se battery as reported in literatures.8−10 In general, selenium exists in several allotropic forms: gray (trigonal) selenium (containing Sen helical chain polymers); rhombohedral selenium (containing Se6 molecules); three deep-red monoclinic forms: α-, β-, and γselenium (containing Se8 molecules); amorphous red selenium and black vitreous selenium. The most thermodynamically stable and the densest form is gray (trigonal) selenium, which contains infinite helical chains of selenium atoms. All other forms revert to trigonal selenium on warming, which adopts a helical polymeric chain and the Se−Se distance is 2.37 Å and the Se−Se−Se angle is 103°.11 It has been found that the phase transition between crystalline Se (c-Se) and a-Se really existed in the charge/discharge process.12−14 It is known that amorphous material is in a metastable state with respect to its crystalline counterpart and they are different in the array of atoms, which means that the former is more active than the latter thanks to the “dangling bonds” and a higher surface-bulk Received: July 18, 2015 Revised: September 21, 2015

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DOI: 10.1021/acs.chemmater.5b02753 Chem. Mater. XXXX, XXX, XXX−XXX

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Chemistry of Materials

ESCALab220i-XL electron spectrometer from VG Scientific using 300W Al KR radiation. The base pressure was about 3 × 10−9 mbar. The binding energies were referenced to the C 1s line at 284.6 eV from adventitious carbon. Electrochemical Measurements. The electrochemical performances of the as-obtained a-Se c/a-Se and c-Se NWs were tested as cathode materials by fabricating CR2032 coin cells. Metallic Li foil was used as the negative electrode. The positive electrode was prepared by first mixing a slurry containing 70 wt % Se material, 10 wt % carbon black, and 20 wt % polyvinylidene fluoride (PVDF) binder previously dissolved in N-methyl-2-pyrrolidinone (NMP), and the proper amount of NMP as dispersant. The slurry was then coated on aluminum foil using a doctor blade and dried in the air. A half-cell was constructed in an argon-filled glovebox. Finally, button-type model cells used porous Celgard 2400 as the separator and a standard 1.0 mol·L−1 LiPF6 solution in a mixture of EC (ethylene carbonate)-DMC (dimethyl carbonate)-DEC (diethyl carbonate) (1:1:1, v/v/v) as the electrolyte. Electrochemical performances of the samples were tested with a LAND battery test instrument (CT2001A) at different rates between 0.01 and 3.0 V vs Li+/Li. All the experiments were conducted at room temperature. Cyclic voltammogram (CV) was conducted using a CHI 660D electrochemical workstation with a scan rate of 0.1 mV·s−1.

ratio in the amorphous phase. Based on the above comprehensive analysis, it is reasonable to believe that different chemical reaction processes may occur when using a-Se or c-Se as the cathode material in Li−Se battery. However, the difference between them is never studied in the past. Therefore, we are committed to research on phase-induced different reaction of Li-ion with Se in Li−Se battery and establish foundation for the future work. Note that among the various morphologies of selenium,15−18 selenium NWs are extremely attractive since this special one-dimensional character not only can accommodate volume expansion without pulverization, but also can facilitate axial charge transport and short radial Li-ion diffusion distances, which will favor the obvious signals to be observed. In this work, first a-Se, c-Se, and c/a-Se NWs with the same size (70 nm) have been prepared through a facile ball milling process, which is novel and different from most of previous studies and the following sintering at 80 °C, respectively. Their crystallization, compositions and microstructures have been confirmed by the corresponding XRD, SEM, XPS, TEM, and HRTEM measurements. In addition, the states of Sen helical chain in them have also been demonstrated through their Raman spectra examinations. The following charge/discharge measurements and the synchronous XPS examinations in Li− Se batteries proved that a-Se reacted with Li+ through a more complex chemical path compared with c-Se due to the different stabilities of Sen helical chains in them, which resulted in their particular storage capacities: 755 mAh·g−1 (a-Se), 705 mAh·g−1 (c/a-Se), and 250 mAh·g−1 (c-Se), respectively.

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RESULTS AND DISCUSSION Compositions and Structures. First, the overall phase, crystallinity and purity of the products ball-milled for different times and raw Se powder are examined by XRD measurements as shown in Figure 1(a). It can be seen that the original power

EXPERIMENTAL SECTION

Synthesis of a-Se NWs. Selenium powders used in this work were purchased from Aladdin Industrial Corporation and their purity is 99.99%. The ZrO2 balls with 1.5 mm diameter were employed as ballmilling medium. In a typical experiment, Se powders were first blended together with ZrO2 balls, with a weight ratio of Se:ZrO2 = 1:15. After that, the mixtures were ball-milled using a planetary ball mill (pulverizette 7 plus, FRITSCH) in a ZrO2 vessel at a speed of 600 round/min (rpm) under argon atmosphere for 10 h. Finally, the asobtained powder was cooled to room temperature and then was washed with alcohol and water for several times. Synthesis of c/a-Se NWs and c-Se NWs. The above as-obtained sample was sintered at 80 °C under Ar gas for 2 and 5 h, respectively. Characterization. All the samples’ compositions were first characterized by X-ray diffraction measurements (Rigaku D/max IIIA, Cu Kα) with a scan rate of 0.05°/s in 2θ range of 10−70°. Raman spectroscopic analysis was performed by using a micro-Raman system with an Ar ion laser (488 nm) and a probing laser 50W·cm−2 was guided during the illumination. Scanning electron microscopy (SEM) images were taken with a JEOL-5600LV scanning electron microscope, using an accelerating voltage of 20 kV. Transmission electron microscopy (TEM), high resolution TEM (HRTEM) and the select area electron diffraction (SAED) images were recorded on a JEOL2010 TEM at an acceleration voltage of 200 kV. The Brunauer− Emmett−Teller (BET) specific surface area of the powders was analyzed by nitrogen adsorption in a Micromeritics ASAP 2020 nitrogen adsorption apparatus (U.S.). All of the prepared samples were degassed at 60 °C prior to nitrogen adsorption measurement. The BET surface area was determined by a multipoint BET method using the adsorption data in the relative pressure (P/P°) range of 0.05−0.3. The desorption isotherm was used to determine the pore size distribution using the Barret-Joyner-Halender (BJH) method, assuming a cylindrical pore model. UV−vis diffused reflectance spectra of the samples were obtained from a UV−vis spectrophotometer (UV2550, Shimadzu, Japan). BaSO4 was used as a reflectance standard. X-ray photoelectron spectroscopy (XPS) data were obtained by an

Figure 1. (a) Time-dependent XRD patterns recorded for the samples obtained at different ball-milling durations. (b) Raman spectrum of the as-obtained product ball-milled at 600 rpm for 10 h and its sintering products at 80 °C under Ar gas for 2 and 5 h, respectively. (c), (d), (e) and (f) SEM, TEM, HRTEM and SAED images of the as-obtained sample ball-milled at 600 rpm for 10 h.

is pure trigonal phase Se, which is consistent with the result reported in literature (JCPDF-No. 73−0465). After a ballmilling treatment at a speed of 600 rpm, the samples convert from highly crystalline nature to amorphous phase, which implies a possible refinement, liquefaction (melting point of Se: 494 K) and the following solidification process under a longer duration of such high energy ball-milling. The composition of the as-obtained amorphous product is further confirmed by Raman spectrum measurements, as shown in Figure 1(b). It has been pointed out in literatures that the resonance peak of B

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Chemistry of Materials crystalline Se (trigonal) locates at ∼236.8 cm−1 and that of amorphous selenium centers at around 250 cm−1, which are attributed to the stretching vibration of Sen helical chain.19,20 Furthermore, the bond strength between Se−Se atoms in the helical chain of a-Se is higher than that of c-Se due to the deficient electronic and van der Waals forces between the Sen helical chain.21 As described in Figure 1(b), after 10 h ballmilling the Raman spectrum curve of the as-obtained product exhibits a peak at 250 cm−1, which demonstrates that the products is pure amorphous Se. This result agrees with that observed in XRD measurements. In addition, after sintering the amorphous Se at 80 °C for 2 and 5 h, respectively, c/a-Se and pure c-Se are acquired. They display a peak at 236.8 cm−1 and a peak at 250 cm−1 in their according Raman spectra, as shown in Figure 1(b). In order to study the morphology of the product, we studied the resulting product by SEM (Figure 1c), TEM (Figure 1d), HRTEM (Figure 1e) and SAED (Figure 1f) measurements. After milled at 600 rpm for 10 h, the resulting product is a large area of selenium NWs. Their diameter is about 70 nm and the length is up to several micrometers. As shown in Figure 1(e) and (f), it is found that there is no significant lattice fringe in the visible range and the electron diffraction pattern taken from a typical Se NWs also proves the amorphous nature of these Se NWs. Compared with the previous reports of template methods for Se and its related 1D structures,22,23 this work through only a ball-milling process provides a more simple one for Se 1D nanostructures. After sintering the above sample at 80 °C under Ar gas for 2 and 5 h, respectively, interesting c/a-Se NWs and c-Se NWs were obtained, which maintained the 1D morphology and the size (70 nm), as shown in Figure 2.It can be seen in the SEM,

As shown in Figure 3(a), three typical type IV isotherm curves are observed, revealing the existence of mesoporous structures.

Figure 3. (a) N2 adsorption desorption isotherms and (b) pore size distributions of the as-prepared a-Se, c/a-Se, and c-Se NWs.

Figure 3(b) depicts their corresponding pore size distributions, concentrating in 21−23 nm range. With the incorporation of crystalline Se, as shown in Table 1, the BET surface areas decrease from 47.8 m2•g−1 to 28.7 m2•g−1, which may be due to the aggregation induced by the annealing treatment of the samples. Table 1. Surface Areas, Pore Volumes and Pore Sizes of the As-Obtained a-Se NWs, c/a-Se NWs and c-Se NWs sample

surface area (m2· g−1)

total pore volume (cm3· g−1)

pore size (nm)

a-Se NWs c/a-Se NWs c-Se NWs

47.8 40.0 28.7

0.025 0.021 0.015

22.4 23.2 21.8

Growth Process of a-Se. Moreover, in order to evaluate the composition and morphology evolution of the as-obtained Se products with ball-milling duration, the time-dependent SEM and XRD measurements have been carried out, as shown in Figure 4 and Figure 1(a). After a ball-milling for 3 h, many

Figure 2. SEM, TEM, and HRTEM images of the products sintered at 80 °C under Ar gas for 2 h: (a), (b) and (c), 5 h: (d), (e), and (f), respectively.

Figure 4. (a), (b), and (c) SEM images; (d) visible and ultraviolet light images; and (e) UV−vis absorption spectra of the products obtained at different moments.

TEM, and HRTEM images of c/a-Se NWs (Figure 2a, b and c) that there is no obvious lattice fringes in a typical NW’s perspective edge, but in its interior the lattice fringes of trigonal Se (001) plane is clearly observed. Interestingly, the c-Se regions are connected by a-Se parter and form a bamboo-like structure. After a prolonged sintering duration, as shown in Figure 2(d) (e) and (f), fully crystalline Se NWs are obtained. The phenomenon of the above phase transition can be interpreted in the framework of Ostwald’s rule of stages24 that predicts the formation of phases to proceed with increasing condensation time toward increasing stability.25,26 The surface areas and porous nature of the above as-obtained a-Se NWs, c/a-Se NWs and c-Se NWs have also been studied.

Se nanoparticles (size ≤70 nm) land on bulky blocks (raw Se reactant), as shown in Figure 4(a). Its XRD pattern in Figure 1(a) also exhibits the weakened diffraction peaks compared with that of raw Se reactant. All the above results confirmed that under the ball-milling treatment a refinement effect has been realized. The following analysis proves that it is only the first step of domino effect. After prolonging the duration to 5 h and 7 h, the Se nanoparticles dispersed on the bulky particles are turned into long and thin rodlike nanoparticles (Figure 4b and c). Correspondingly, all strongest diffraction peaks disappear in XRD patterns (Figure 1a), indicating an C

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Figure 5. Initial charge−discharge curves of the (a) a-Se, (b) c/a-Se, and (c) c-Se NWs (600 rpm, 10 h) at a current density of 0.1 C between 0.01 and 3.0 V. (The red ball is Se and green ball is Li) (d) and (e) XPS spectra of a-Se NWs before and after charge−discharge in Li−Se cell.

shown in Figure 4(e). According to the electronic band structure calculated using the empirical pseudopotential method (EPM),29 the outermost excitations from the valence band correspond with two direct transitions along the H, and one along the A direction.30 Transition energies of ∼2.1, ∼ 2.6, and ∼3.2 eV, respectively, were predicted for these excitation pathways. The lowest energy transition from the bonding states (i.e., 2.1 eV) could be attributed mostly to interchain interactions by calculating the electronic charge distributions. The higher energy transitions were determined to a rise from covalent bonds within the selenium chain. Spectroscopic absorption studies by Bogomolov et al. on isolated helical chains of Se atoms further confirmed that the strong absorption at ∼2.2 eV came from the intermolecular transition between the spiraling chains of selenium. This transition was only present in the trigonal phase, and not in systems that only contained spiral chains of Se atoms isolated by a solid matrix of another material or in solid Se of any other allotropic form (e.g., the amorphous and monoclinic phases).21 Here, it is found that the absorption spectrum for a-Se NWs (10 h) exhibits prominent features at 2.2, 2.6, 3.5, and 5.77 eV, each slightly blue-shifted in comparison to the values expected for the samples ball-milled less than 10 h. Interestingly, in this work, the strong absorption at ∼2.2 eV came from the interchain transition is found for fully amorphous Se NWs, which maybe related with the interconnection among the large area a-Se NWs. At the same time, it is also found that the absorptions at 2.6, 3.5, and 5.77 eV from covalent bonds within the selenium chain are enhanced with the phase transition from c-Se to a-Se, meaning strengthened covalent bonds of Se−Se, which agrees with the difference in Raman results of c-Se and aSe (Figure 1b). In addition, it has also been proved in our comparison experiments that the ball-milling speed (determining the as-provided mechanical energy) is critical to the formation of pure a-Se. Under a lower ball-milling speed, such as 200 and 400 rpm, only bulky particles like raw Se or a-Se nanorods are observed, as shown in Figure S1(a) and (b), implying a deficient crushing and recrystal process. When the

amorphous transition process. Considering the low melting point of Se (494 K), it is comprehensible that the high-energy ball-milling induces a liquefaction process and the following solidification. Additionally, with the ball-milling time gradually increasing, the length of the NWs becomes longer and longer and simultaneously the raw crystalline Se bulky particles get smaller. Finally, as the reaction time is extended to 10 h, fully amorphous Se NWs with a diameter of ∼70 nm and a length of about 2 μm are obtained, as shown in Figure 1(c). In addition, no raw Se bulky particles are observed, indicating that an absolute conversion form bulky ones to NWs. It is well-known that Se is in group VIB and has six outer electrons (s2p4). Two electrons are paired in the s-orbital, two electrons in one of the three p-orbitals, and the remaining electrons are available for covalent bonding in half-filled p-orbitals. The special electron configuration results in a spiral chain of atoms with three atoms per turn. The bonds between atoms on the same chain are covalent, whereas between chains they are thought to be a mixture of electronic and van der Waals. This difference is very large and indicates a much greater cohesion between atoms in the same spiral than those in different spirals.27 During the amorphization process, selenium tends to grow preferentially along the spiral chain because of the necessity to obtain a match between the highly anisotropic structure and uniaxial geometry of the one-dimensional species, as well as to favor the strong covalent bonds in helical chains over the relatively weak van der Waals forces among chains.28 Raman results in Figure 1(b) also confirm the existence of abundant Se helical chains in a-Se. The growth mechanism mentioned above is consistent with our visual observations, as well as spectroscopic measurements. Figure 4(d) displays the samples’ pictures of suspensions with a same concentration in ethanol under visible light and ultraviolet light (365 nm), respectively. Clearly, with the increasing of ball milling time, the color of the products change from gray to brick red under visible light and from lucency to dark under ultraviolet light, suggesting an enhanced absorbance transformation. Such a color transition could be more clearly resolved from the UV−vis diffuse reflectance absorption spectra D

DOI: 10.1021/acs.chemmater.5b02753 Chem. Mater. XXXX, XXX, XXX−XXX

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In Li−Se battery Se is reduced to the polyselenides Li2Sen (n ≥ 4), Li2Se2, and Li2Se sequentially during the lithiation and Li2Se is oxidized to Se through Li2Sen (n ≥ 4) during the (de)lithiation. So the multistep lithiation processes found in aSe NWs and c/a-Se NWs agree with the above mechanism, which is further confirmed by the following XPS characterizations of the a-Se NW cathodes. Figure 5(d) and (e) depicts the a-Se NWs’ 3d XPS spectrum, in which Se 3d5/2 and 3d3/2 peaks locate at ∼55.44 and ∼56.19 eV with a spin−orbit splitting of 0.75 eV, consistent with that of selenium element.31 After the first discharge (lithiation process), Se 3d5/2 and Se 3d3/2 peaks move toward lower binding energies: 55.38, 56.27, 54.33, 55.07, 53.11, and 53.97 eV. According to the literatures, the binding energies at 55.38 and 56.27 eV are due to the Se 3d5/2 and 3d3/2 of polyselenides (Sen2− chain (n ≥ 4));32 the binding energies at 54.33 and 55.07 eV are attributed to the Se 3d5/2 and 3d3/2 of diselenides (Se22−),33 the binding energies at 53.11 and 53.97 eV belong to the Se 3d5/2 and 3d3/2 of Se2− ions.34 Therefore, in the lithiation process the chain-like Sen molecules in a-Se NWs are reduced to the low-valence polyselenides Li2Sen (n ≥ 4), Li2Se2, and Li2Se sequentially. Comparatively, chain-like Se molecules in c-Se NWs are reduced fast to Li2Se, as shown in Figure 5(c). This may be due to its active state and the weakened covalence band in Sen helical chains coming from the sufficient electronic and van der Waals forces between the Sen helical chain,21 which has also been testified in our Raman results (Figure 1b). Later, charged to 3 V, Se 3d5/2, and Se 3d3/2 peaks of a-Se shift back to high binding energies, demonstrating the reversible oxidization of Li2Se to zero-valence chain-like Sen molecules (Figure 5e) through the above-mentioned three (de)lithiation steps, including some Se−O bonds (59.33 eV),35 In order to illustrate the c-Se electrochemical process, the XPS examinations of it have also been conducted, as shown in Figure S4. It can be seen that after the first discharge, there are still Se element, Se22− ions and Se2− ions existed, which implies that its much crystalline core remains nonlithiated. All the analysis points out that different (de)lithiation processes are realized due to the different chemical activities of Sen chains in a-Se and c-Se in the Li−Se cell. In addition, interestingly the second cycle electrochemistry curve of the c-Se/Li cell resembles that of the a-Se/Li one, as shown in Figure S5, which implies the formation of a-Se during its delithiation process. Their CV curves for the Li/Se cells are shown in Figure 6(a). In the Li/Se system of a-Se NWs, the CV starts from an open circuit voltage of 3.0 V and sweeps to 0.8 V during the initial discharge, converting elemental selenium to selenides with three reduction peaks at 1.63, 1.66, and 2.12 V, which also confirms the mechanism for the Li/Se cell lithiation process by

ball-milling speed is elevated higher, such as 800 rpm, c-Se phase emerges, as shown in Figure S1(c) and (d), which is due to the calefaction effect induced by higher speed ball-milling. Furthermore, the effect of sintering temperatures to the morphology of c-Se is also investigated at 100, 150, and 200 °C respectively. As shown in Figure S2, it is clear that under all the above higher temperatures c-Se are obtained. However, the initial NW’s length gets shorter and shorter, even to spherical nanoparticles. So the sintering at 80 °C temperature most favors maintaining the a-Se NW’s morphology and size. Li−Se Battery. Se, as the critical cathode material of Li−Se battery, has a high theoretical volumetric capacity density (3253 mAh·cm−3), which is more important than gravimetric energy density because of the limited battery packing space in portable devices and HEVs. However, it is also noticed that Se has low melting point (494 K) and phase transition temperature (less than 353 K for its NW in this work), and a large sensitivity to thermal treatment,5−7 which has been proved during the charge/discharge cycles of Li−Se battery as reported in literatures.8−10 In addition, it is found that different Li−Se compounds (Li2Sen, Li2Se2 and Li2Se) formed and the phase transition between crystalline Se (c-Se) and a-Se existed in the charge/discharge process.12−14 It is known that amorphous material is in a metastable state with respect to its crystalline counterpart, which means that the former is more active than the latter thanks to the “dangling bonds” and a higher surfacebulk ratio in the amorphous phase. And it is reasonable to believe that different chemical reaction processes ((de)lithiation processes) may occur when using a-Se or c-Se as the cathode material in the Li−Se battery, which is never detailed studied in the past and may establish foundation for the future work about Li−Se battery. In this research, the different (de)lithiation processes of a-Se, c/a-Se, and c-Se have been demonstrated through a serial of detailed comparison experiments. As shown in Figure 5(a), (b), and (c), the galvanostatic discharge−charge voltage performances of a-Se, c/a-Se, and c-Se NWs are first tested, in which the current density is fixed at 0.1 C (1 C = 675 mAh·g−1) and the potential is in the range 0.01−3.0 V. For amorphous and crystalline/ amorphous materials, the initial discharge involves three welldefined plateaus (at 1.55 V, 1.6 and 2.04 V) indicative of the Liion insertion (reaction of Li-ion with Se via getting electrons). While that of c-Se NWs presents only one platform (at 1.55 V). In addition, the as-prepared a-Se NWs and c/a-Se NWs exhibit the reversible capacities of 755 mAh·g−1 and 705 mAh·g−1 respectively, which are much higher than that of c-Se (250 mAh·g−1). That is obvious that the elevated capacities of a-Se NWs and c/a-Se NWs are mainly due to the two extended (de)lithiation processes (or sites) at 1.6 and 2.04 V. It has been proved in literature.10 The comparison of storage capacities between c/a-Se sample and the mixture of a-Se and c-Se (mass ratio 1:1) have been carried out, as show in Figure S3. It is obvious that the capacity of c/a-Se is much higher than that of the mixture of a-se and c-Se, which also proves the initial (de)lithiation process occur at only surface. As the (de)lithiation processes in the initial several cycles occur mainly at surface, the exterior structure and composition of the NWs determine their capacities. In a c/a-Se NW, it can be seen that its exterior part is fully composed of a-Se, as shown in Figure 2b and c, which indicates that c/a-Se NWs have a similar high capacity with a-Se. Its examination results (Figure 5b) prove the above speculation.

Figure 6. (a) Cyclic voltammograms of the a-Se, c/a-Se, and c-Se NWs in the initial cycle vs Li at a scan rate of 0.1 mV·s−1 between 0.8 and 3.0 V. (b) Cycle performances of a-Se c/a-Se and c-Se NWs at a current density of 0.1 C between 0.01 and 3.0 V. E

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Chemistry of Materials correlating the voltage profiles. Selenium is first reduced to Li2Sen (n ≥ 4) at 2.12 V, which is then further reduced to Li2Se2 and Li2Se after discharging to 1.66 and 1.63 V, respectively. During charging, Li2Se is directly oxidized to chain-like Sen molecules.36−39 It is noticed that the overpotential of c/a-Se/Li cell is significantly larger than those of cSe/Li and a-Se/Li cells during the first discharge process. As the (de)lithiation processes in the initial several cycles occur mainly at surface, the exterior Se atoms’ state determines the overpotential during the discharge process. In a c/a-Se NW, it can be seen that its exterior part is fully composed of a-Se and interior part contains many c-Se, as shown in Figure 2b and c. The interaction existing in this special host/guest structure induces the change of c/a-Se NW’s surface tension and Se atoms’ state at the surface. This may be the reason why the overpotential of c/a-Se/Li cell is different from those of c-Se/Li and a-Se/Li cells. The specific capacities and cycling stabilities of a-Se NWs, c/ a-Se NWs, and c-Se NWs have also been studied. As shown in Figure 6(b), from the first cycle to second one the fading is more serious, but from the second cycle their capacities are almost does not decay. Hence, there is an electrochemical activation process during the first lithiation, and similar phenomena have been reported in other cathode materials.3,40 At the same time, in a-Se NWs cathode the Coulombic efficiency shows a rapid increase to 96.6% after the second cycle, which then approaches 100% over 100 cycles (the pink dots in Figure 6b). It has been found in literature21 and our Raman results (Figure 1b) that the electronic and van der Waals force between the Sen helical chains of c-Se is stronger than that of a-Se, which implies that a more stable structure composed of Sen helical chain in c-Se exists. It is also noticed that the lower capacity of c-Se (Figure 5c) compared with a-Se and c/a-Se comes from the less (de)lithiation reactions during the charge−discharge process. All the above factors result in the better recycling stability of c-Se electrode.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. *E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank the Program for NCET in University (NCET-130754), the Natural Science Foundation of China (Grant No. 51272050 and 51072038), Harbin Sci-tech innovation foundation (RC2012XK017012), Harbin Youth Fund (RC2014QN017004), the Fundamental Research funds for the Central Universities (HEUCF2015), Outstanding Youth Foundation of Heilongjiang Province (Grant No. JC201008), and Youth Fund of Heilongjiang Province (QC2014C006) for the financial support of this research.

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REFERENCES

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CONCLUSIONS In summary, different (de)lithiation mechanisms between a-Se and c-Se have been proved in this wok. It is found that in a-Se lithiation process Se helical chains are first reduced to Li2Sen (n ≥ 4) and then further reduced to Li2Se2 and Li2Se, respectively. Comparatively, Se helical chains in c-Se NWs are reduced fast to Li2Se, which is due to its active state and the weakened covalence band in Sen helical chains coming from the sufficient electronic and van der Waals forces between the Sen helical chain. In addition, the different (de)lithiation mechanisms result in their distinct storage capacities: a-Se NWs and c/a-Se NWs exhibit the reversible capacities of 755 mAh·g−1 and 705 mAh·g−1 respectively, which is much higher than that of c-Se (250 mAh·g−1). This work has promoted us to better comprehend and correlate the formation of intermediate phases with the electrochemical performance of Li−Se cells.

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experiment data (charge−discharge curves and XPS) (PDF)

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The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b02753. Additional SEM images and XRD patterns at different ball-milling speeds and sintering temperatures; contrast F

DOI: 10.1021/acs.chemmater.5b02753 Chem. Mater. XXXX, XXX, XXX−XXX

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DOI: 10.1021/acs.chemmater.5b02753 Chem. Mater. XXXX, XXX, XXX−XXX