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Insights into the Distinct Lithiation/Sodiation of Porous Cobalt Oxide by in Operando Synchrotron X-ray Techniques and ab initio Molecular Dynamics Simulations Gui-Liang Xu, Tian Sheng, Lina Chong, Tianyuan Ma, Chengjun Sun, Xiaobing Zuo, Di-Jia Liu, Yang Ren, Xiaoyi Zhang, Yuzi Liu, Steve M. Heald, Shi-Gang Sun, Zonghai Chen, and Khalil Amine Nano Lett., Just Accepted Manuscript • DOI: 10.1021/acs.nanolett.6b04294 • Publication Date (Web): 10 Jan 2017 Downloaded from http://pubs.acs.org on January 12, 2017

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Insights into the Distinct Lithiation/Sodiation of Porous Cobalt

Oxide

by

in

Operando

Synchrotron

X-ray

Techniques and ab initio Molecular Dynamics Simulations Gui-Liang Xu,1 Tian Sheng,2 Lina Chong,1 Tianyuan Ma,3 Cheng-Jun Sun,4 Xiaobing Zuo,4 Di-Jia Liu,1 Yang Ren,4 Xiaoyi Zhang,4 Yuzi Liu,5 Steve M. Heald,4 Shi-Gang Sun,2 Zonghai Chen1,* and Khalil Amine1,* 1

Chemical Sciences and Engineering Division, Argonne National Laboratory, 9700

South Cass Avenue, Argonne, Illinois 60439, USA 2

Collaborative Innovation Center of Chemistry for Energy Materials, State Key

Laboratory Physical Chemistry of Solid Surfaces, Department of Chemistry, Xiamen University, Xiamen, 361005, China 3

Materials Science Program, University of Rochester, Rochester, NY 14627, USA

4

X-ray Science Division, Advanced Photon Source, Argonne National Laboratory,

9700 South Cass Avenue, Argonne, Illinois 60439, USA 5

Nanoscience and Technology Division, Argonne National Laboratory, 9700 South

Cass Avenue, Argonne, Illinois 60439, USA

1

ACS Paragon Plus Environment

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Abstract Sodium-ion batteries (SIBs) have been considered as one of the promising power source candidates for the stationary storage industries owing to the much lower cost of sodium than lithium. It’s well known that the electrode materials largely determine the energy density of the battery systems. However, recent discoveries on the electrode materials showed that most of them present distinct lithium and sodium storage performance, which is not yet well understood. In this work, we performed a comparative understanding on the structural changes of porous cobalt oxide during its electrochemical lithiation and sodiation process by in operando synchrotron small angel X-ray scattering, X-ray diffraction and X-ray absorption spectroscopy. It was found that compared to the lithiation process, the porous cobalt oxide undergo less pore structure changes, oxidation state and local structure changes as well as crystal structure evolution during its sodiation process, which is attributed to the intrinsic low sodiation activity of cobalt oxide as evidenced by ab initio molecular dynamics simulations. Moreover, it was indicated that the sodiation activity of metal sulfides is higher than that of metal oxides, indicating a better candidate for SIBs. Such understanding is crucial for future design and improvement of high performance electrode materials for SIBs.

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Introduction In the past decades, rechargeable lithium-ion batteries (LIBs) have been considered as the most promising power source candidate for portable electronic devices owing to their high energy density.1 However, for large scale application such as stationary storage industries, cost would be a major issue due to the limited lithium resource.2 Compared to lithium, sodium is much more abundant and distributed uniformly around the world.3 Therefore, sodium-ion batteries (SIBs) have recently received growing interest, spurred by the rapid advance in rechargeable battery technology and fast increasing demand in the market.4-7 As usual in any new chemistry, there are always setbacks during the development of high performance electrode materials. Moving from lithium-ion batteries to sodium-ion batteries brings a lot of challenge on the material development and mechanism understanding. Recent discoveries on the electrode materials showed that most of them present inferior sodium storage capability than their lithium storage capability.8-14 Wang et al. reported a mesoporous Sn/C anode, which demonstrated lower capacity and inferior rate capability in a Na-ion battery than in a Li-ion battery.15 The state of the art Si-based anode for LIBs could deliver a reversible capacity of 1000-3000 mAh g-1, but has very low capability to alloy with Na.16 Wu and co-workers reported that Fe2O3 multi-shelled core-shell microspheres could deliver a high reversible capacity of over 1000 mAh g-1 with a long cycle life (200 cycles) at 400 mA g-1 and high rate capability (>700 mAh g-1 at 6 A g-1). However, when used as an anode of SIBs, the 1st reversible capacity was only 600 mAh g-1, accompanied by a continuous capacity fading in 80 cycles even charge/discharged at a lower current density of only 40 mA g-1.17 Similar results were also reported on other metal oxide anodes,18-20 crystalline Ge21, 22 and so on. The previously reported works in the open literatures were mostly using trial-and-error approach, and the physics behind the poor performance of these materials as sodium-ion anodes were poorly discussed. These phenomena were conventionally attributed to the larger radius of Na+ than Li+, which may lead to sluggish diffusion kinetics of Na+ during the (de)sodiation process. However, Han and 4


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co-workers found that the Na ion diffusivity in NaxAl2O3 turns out to be much higher than Li ion diffusivity in LixAl2O3, a result opposite to the conventional stereotype based on the atomic radius consideration, which may be due to the weaker Na-O bond strength than the Li-O counterpart.23 Another possible factor that has been frequently discussed is the larger volumetric changes induced during the (de)sodiation process since the radius of Na+ (1.02 Å) is larger than Li+ (0.76 Å). However, using in-situ transmission electron microscopy (TEM), Shahbzian-Yassar et al. found that ZnO nanowires do not show any cracks and degradation after sodiation while the lithiated ZnO nanowire shows multiple glassy domains, which has low strength, ductility and further poor battery performance.24 Although in-situ TEM technique is a powerful tool to monitor the morphology change of electrode materials during the electrochemical reaction, it still has some limitations. An ionic liquid or Li2O (Na2O in the case of Na ion batteries) are usually used in the high-vacuum TEM chamber, which makes the electrochemical process different from real batteries based on traditional organic liquid electrolytes. Using a small all-solid-state single nanowire device, Mai and co-workers reported a result different from that reported by Shahbzian-Yassar et al., in which they found that the conductivity degradation and structure destruction for Na ions are more severe than those of Li ions during the electrochemical processes of single H2V3O8 nanowire.25 Theoretical calculations such as molecular dynamics simulations and first principles calculations were also used to study the lithium/sodium ion diffusion difference.23,

26, 27

The strengths of these different

techniques allow for unique information to be gained from each but sometimes are contradictory. Therefore, a deep understanding on the (de)lithiation/(de)sodiation process of battery materials is highly desired to unravel their distinct electrochemical behaviors by a combination of advanced techniques. Rechargeable batteries are very complex systems, which involves liquid and solid, amorphous and crystalline phases, and sometimes accompanied by large volumetric changes.28 Most synchrotron X-rays have a wavelength ranging from 0.01 to 10 nm, which makes it possible to probe the atomic and molecular structures for most of the materials. Synchrotron X-ray diffraction (SXRD) has played a great role on 5


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unraveling the working mechanism of energy-storage materials by tracking their phase transformation during the charge/discharge process.29 X-ray absorption fine structure (XAFS) especially X-ray absorption near edge spectroscopy (XANES) is sensitive to the electronic and/or oxidation state changes of all the atoms of the targeted type in the samples during cycling no matter it is crystalline or amorphous, while the extended X-ray absorption fine structure (EXAFS) could reflect the local structure changes of the active elements during charge/discharge.30 Small angle X-ray scattering (SAXS) is one of the most useful analytical techniques to investigate the properties of porous materials since the SAXS technique works very well in characterizing unique nanostructures of materials with the probing range of 1-100 nm.31, 32 In this work, using porous cobalt oxide as a model material, we performed a comparative understanding on its structural changes during its electrochemical lithiation and sodiation process by in operando SAXS, SXRD and XAFS, and further unravel the mystery behind its distinct (de)lithiation and (de)sodiation capability. It was found that compared to the lithiation process, the porous cobalt oxide undergo less pore structure changes, oxidation state and local structure changes as well as crystal structure evolution during its sodiation process, which is owing to the intrinsic low sodiation activity of cobalt oxide as evidenced by ab initio molecular dynamics simulations, and thus leads to its inferior sodium storage performance than lithium storage performance. Moreover, it was found that the sodiation activities of metal sulfides are better than that of metal oxides, indicating better candidate for SIBs. Such understanding is crucial for future design and improvement of high performance electrode materials for SIBs.

Experimental Section Preparation of porous cobalt oxide. The porous cobalt oxide (PCO) was prepared via a two-step process, involving the synthesis of Co-MOF (ZIF-67) concave nanocubes and the subsequent calcination treatment at 300 °C under air flowing. In the first step, the ZIF-67 concave nanocubes precursors were prepared via a simple precipitation reaction at room temperature according to our previous work.33 Typically, 6


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50 ml methanol solution of Co(NO3)2·6H2O (1 g, 3.84 mmol) was mixed with a 80 ml methanol solution of 2-methylimizole (2.5 g, 30.4 mmol), followed by stirred for 5 min and kept at room temperature (RT) for 24 h. The resulting purple precipitates were collected by centrifugation, washed with methanol for 3 times, and then dried at 50 °C under vacuum overnight. In the second step, the as-prepared Co-MOF (ZIF-67) precursors were loaded into a ceramic crucible and then heated to 300 °C under air atmosphere with a heating rate of 5 °C min-1. After dwelling at 300 °C in air for 6 h, the purple precursor was finally converted into black powder. Materials Characterization. The morphologies and structures of the as-prepared materials were characterized by field emission scanning electron microcopy (HITACHI S-4700-II) and transmission electron microcopy (FEI Titan). The Brunauer-Emmett-Teller (BET) specific surface area (SSA) and pore structure parameters of the PCO were determined from the adsorption isotherm of nitrogen at 77K (ASAP-2020 Micromeritics Co., USA). The content of the impurity in the PCO was measured by thermo gravimetric analyses (TGA) on a STA 449 F3 instrument under air atmosphere from room temperature to 800 °C with a heating rate of 10 °C min-1. The phase structure of the PCO was measured by high-energy synchrotron XRD at the Beamline 11-ID-C of Advanced Photon Source (APS) at Argonne National Laboratory with wavelength of 0.117418 Å. During the measurement, the sample was sealed in Kapton tube and XRD data was recorded with a large-area 2D detector up to wave vectors of 25 Å-1. XRD Rietveld refinement of PCO was performed by using GSAS software. Electrochemical Characterization. The slurry was prepared by spreading a mixture of 70 wt. % active material, 20 wt. % carbon black and 10 wt. % sodium aliginate (2 wt. %) onto a copper foil current collector. The as-prepared electrodes were then dried at 80 °C in a vacuum oven for 24 h. The loading density was controlled at around 1.5 mg cm-2. The electrochemical performances of the PCO electrodes were characterized by assembling them into coin cells (type CR2032) in an argon-filled glove box under conditions that the content of moisture and oxygen were both below 0.5 ppm. For lithium cells assembly, the electrodes were separated from 7


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lithium chip by a separator (Celgard 2325). Electrolyte used in the lithium cells was 1.2 M LiPF6 in a mixture of ethylene carbonate (EC) and ethyl methyl carbonate (EMC) (3/7, v/v) with 10 wt. % fluoroethylene carbonate (FEC) as additive. For sodium cells assembly, the electrodes were separated from sodium chip by a separator (glass fiber, Grade GF/F Glass Microfiber Filter Binder Free, circle, 125 mm). The electrolyte used in the sodium cells was 1 M NaPF6 in propylene carbonate (PC) with 2 vol.% FEC as additive. The cells were charged and discharged using a MACCOR cycler at room temperature unless specific note. In-situ electrochemical impedance spectroscopy was recorded on a Solartron Analytical 1470 System in a frequency range of 100 kHz to 0.1 Hz. In operando synchrotron X-ray measurements during battery cycling. Typical coin cells with holes at both top and bottom cases were used for in operando electrochemical cells. After cell assembly, the holes were sealed with Kapton tape to prevent the contamination of moisture.28 During the in operando measurements, the cells were continuously charge/discharged without any pause process as the data acquisition takes not more than 4 minutes (SAXS: 5s, SXRD: 5s, XAFS: 4 minutes). In operando SAXS was carried out at sector 12-ID-B of the APS. The 2D SAXS images were radically averaged to 1D profiles and presented in X-ray scattering intensities verse momentum transfer, q (q=4πsinθ/λ, where θ is one-half of the scattering angle, and λ=0.886 Å is the wave length of the X-ray at 12-ID-B), measured in the range of 0.01-1.0 Å-1. In operando SXRD was conducted at sector 11-ID-D of the APS of Argonne national laboratory (λ=0.8 Å). While in operando XAFS experiments on Co K-edge between 7526 eV and 8526 eV were carried out in transmission mode at Beamline 20-BM-B of the APS. The incident beam was monochromatized by using a Si (111) fixed-exit, double-crystal monochromator. Ab initio molecular dynamics (AIMD) simulations method. The spin-polarized electronic structure calculations were carried out using the Vienna Ab-initio Simulation Package (VASP) with Perdew-Burke-Ernzerh (PBE) functional of exchange-correlation. The projector-augmented-wave (PAW) pseudopotentials were utilized to describe the core electron interaction. The on-site Coulomb repulsion 8


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correction term (U = 2.0 eV) (PBE+U) was used in the 3d electrons of Co. The Co3O4

(110) and Co3S4 (110) surfaces were modeled as a p(1×2) periodic slab with six atomic layers with 36 Co atoms and 48 O (S) atoms. The bottom layers including 8 Co atoms and 8 O (S) atoms were fixed and other atoms were relaxed during AIMD simulations. A bulk metal of 24 Li (Na) atoms was placed above the surface and the vacuum layer was 30 Å for building three systems, i.e, Li/Co3O4, Na/Co3S4 and Na/Co3O4. The cut-off energy was 400 eV and a 2×1×1 Monkhorst-Pack k-point sampling was used. For each system, we performed the AIMD simulation within the canonical (NVT) ensemble at a constant temperature of 1000 K to accelerate the reaction rate.

Results and Discussions Cobalt oxides (Co3O4 and CoO) have been explored as a potential anode material for LIBs.18,

34-36

It was widely reported that bulk cobalt oxide suffers from low

electronic conductivity and large volumetric changes during the cycling process, leading to rapid capacity fading and poor rate capability. Therefore, in order to increase the reversible capacity and extend the cycle life, we have prepared a porous cobalt oxide (PCO) material, which was derived from thermal decomposition of Co-based MOF under air atmosphere. Figure 1a shows the synchrotron high energy X-ray diffraction (HEXRD) pattern of the as-prepared PCO material, in which strong diffraction peaks could be clearly seen. The Rietveld refinement further indicates that it is composed of two phases including cubic Co3O4 (space group: Fd-3m) and cubic CoO (space group: Fm-3m) with a weight ratio of around 4/1. The detailed refinement parameter could be seen in the Table S1. Figure 1b shows the thermo gravimetric analysis (TGA) curve of the PCO under air atmosphere from room temperature to 800 o

C, in which a weight loss of around 4.2 wt.% is observed at a temperature range of

200-400 oC. This may be owing to a small amount of residual Co-MOF, which was not decomposed into cobalt oxide during the thermal calcination process (300 oC for 6 h). Figure 1c shows the N2 adsorption-desorption isotherms curve of the PCO, which presents a typical type IV isotherm, indicating the existence of porous structures. 9


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According to the Brunauer-Emmett-Teller (BET) model, the specific surface area of the PCO was measured to be 45.02 m2 g-1. The PCO has a wide pore size distribution ranging from 0-10 nm, which may be formed due to the release of CO2 during the thermal decomposition of Co-based MOF. The pore volume of the PCO is 0.172 cm3 g-1 as calculated by the Horváth-Kawazoe method. The scanning electron microscopy (SEM) image of the PCO in Figure 1d shows that it present a porous structure. This was further confirmed by the transmission electron microscopy (TEM) images in Figure S1a and Figure S1b, in which interconnected nano-sized particles of around 10-20 nm with lots of mesopores can be observed. The high resolution TEM (HRTEM) image of the PCO (Figure S1c) shows clearly two sets of clear lattice fringes with interplanar distances of 0.466 nm and 0.281 nm, which are in perfect agreement with the d-spacing of (111) and (220) planes of Co3O4, respectively. The lattice fringe with d-spacing of 0.243 nm cannot be assigned as it is close to the d-spacing value of (311) plane for Co3O4 (0.244 nm) and (111) plane for CoO (0.246 nm). However, the selected area electron diffraction (SAED) pattern in Figure S1d indeed confirms the co-existence of polycrystalline CoO and Co3O4 in the as-prepared PCO, coinciding with the HEXRD result. The electrochemical performance of the PCO anode was evaluated by assembling them into coin cells with lithium and sodium as counter electrode, respectively. Noted that FEC was both added in the electrolytes as the FEC additive could form a homogeneous and relatively thinner surface layered on the electrode compared to additive-free electrolytes, which could prevent the electrode from electrolyte decomposition and also provide lower resistance on the electrode surface and probably enhances the efficiency and conductivity.37 Figure 2a shows the 1st, 2nd and 40th discharge/charge curves of the Li/PCO cell at a discharge/charge rate of 0.1 C (1C=890 mA g-1). At the beginning of the discharge, the voltage decreases quickly to about 1.1 V (vs. Li/Li+) with a long discharge plateau and then gradually declines to 0.01 V (vs. Li/Li+). Such variation is mainly attributed to the conversion reaction of cobalt oxide with Li to form Co and Li2O. In the initial charge process, there appears a main voltage plateau at about 2.0 V (vs. Li/Li+), corresponding to 10


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the process of re-oxidation of cobalt nano-domains to cobalt oxide and the decomposition of Li2O matrix to Li.34 The initial discharge capacity of the PCO anode is measured at 1330.3 mAh g-1, exceeding the theoretical capacity of Co3O4 (890 mAh g-1) and CoO (745 mAh g-1). The excess capacity is usually ascribed to the decomposition of electrolytes to form a solid electrolyte interphase (SEI) layer.34 The initial reversible capacity is measured at 870.8 mAh g-1, accounting for an initial coulombic efficiency of 65.4% that is common for transition metal oxides anode materials.20 The corresponding cycling performance of the PCO anode is shown in Figure 2b. The capacities increase slightly along with the discharge/charge process, which is a common phenomenon for the cobalt oxide anode and also other transition metal oxide anode materials such as Fe2O3,17 Fe3O420 and MnO.38 The possible reasons have been ascribed to the gradual decomposition of Li2O that generated during the 1st discharge process and the reversible growth of a polymeric gel-like film on the surface of the progressively pulverized particles resulting from the electrochemical grinding effect.17, 18, 34, 39 As can be seen, a reversible capacity as high as 839.5 mAh g-1 can be still maintained after 40 cycles at 0.1 C, leading to a high capacity retention of 96.4%. The good cycling performance of the as-prepared PCO is comparable to the previous reported nanostructured cobalt oxide anodes for LIBs.34, 40, 41 This may be owing to its porous structures, which can shorten the diffusion paths of Li+ and accmondate the volumetric changes during the repeated discharge/charge processes. Given the good lithium storge performance of the PCO anode, we further evaluate its sodium storage performance. However, it was found the PCO anode demonstrate an inferior sodium storage performance than its lithium storage performance. Figure 2c shows the 1st, 2nd and 40th discharge/charge curves of the Na/PCO cell at 0.1 C (1C= 890 mA g-1), which present a characteristic voltage profile of cobalt oxide as anode of SIBs, including a sloping curve in the potential range of 1-0.2 V (vs. Na/Na+), a flat plateau at 0.2 V (vs. Na/Na+) and a short sloping tail below 0.2 V (vs. Na/Na+) at the first discharge and two inclined plateaus in 0-1.0 V and 1.0-2.0 V ranges during the first charge process.18, 35 The initial discharge 11


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capacity is about 866.2 mAh g-1, which is much lower than its initial Li-storage discharge capacity (1330.3 mAh g-1) and even its theoretical capacity (890 mAh g-1). Its initial charge capacity is only 471.6 mAh g-1, leading to a low initial coulombic efficiency of only 54.4%. Figure 2d shows the corresponding cycling performance of the Na/PCO cell at 0.1 C, which shows continuous capacity degradation in 40 cycles. The results indicate that the reaction between the PCO and Na is not that complete and reversible as that with Li. Similar results were also reported in other battery materials between their lithium storage and sodium storage capability.17, 18, 42, 43 In-situ electrochemical impedance spectroscopy was used to monitor the impedance variation of Na/CoOx cell during cycled at a charge/discharge rate of 0.1 C. Figure S2 indicates that the electrochemical charge transfer resistance has only slight increase, which is not consistent with the rapid capacity fading of Na/CoOx cell. In order to have a better understanding and unravel the mystery behind their distinct electrochemical performances, in operando synchrotron X-ray techniques were used to monitor the structural changes of the PCO anode during its (de)lithiation/(de)sodiation processes, respectively. To understand the pore structure changes of the PCO anode during its (de)lithiation and (de)sodiation process, we utilized in operando SAXS to monitor the pore structure changes during its discharge/charge process. Unlike in-situ TEM technique that usually focused on single nanowire or nanoparticle based on Li2O (Na2O in SIBs) electrolytes, SAXS could reflect the information of the whole electrode materials in conventional liquid cells. Figure 3a shows the 2D contour plot of in operando SAXS data along with the 1st discharge/charge curve of the Li/PCO cell at 0.1 C. It can be found that the signals in the low q region (0.01-0.08) were gradually decreased and shifted to a lower q region during the discharge process, which may be owing to the uptake of Li that may fill the pores of the PCO and lead to an increase of particle size. We can also see that the signals in the high q region (0.1-0.4) were gradually increased along with the discharge process, which may be due to the formation of metallic Co nano-domains during the conversion reaction of Li with the PCO. However, it can be seen that both the signals in the high q and low q 12


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region cannot resile to its pristine electrode at the end of the 1st charge process. The irreversible porous structural change is also reported in the case of mesoporous Co3O4 anode for LIBs, which may be due to the well-known de-lithiation reaction (Co + Li2O

CoO + 2Li), where CoO phase is also formed at the end of charge process,

rather than completely reversible formation of Co3O4 phase.31 Figure 3b shows the typical SAXS data of the Li/PCO cell at different discharge/charge states. As shown, there are very few changes when the cell was discharged to above 1.07 V, where the main reaction occurred between OCV and 1.07 V may be the decomposition of the electrolytes to form SEI layer. After that, it followed the structural changes in both high and low q region as indicated in Figure. 3a. Figure 3c and Figure 3d shows the 2D contour plot of in operando SAXS data along with the 1st discharge/charge curve and typical SAXS data at different discharge/charge states of the Na/PCO cell at 0.1 C, respectively. It was surprising to see that the pore structure changes for Na/PCO cell during discharge/charge is much smaller than that in Li/PCO cell considering the larger radius of Na+ than Li+, which is opposite to the conventional perspective but in consistent with the in-situ TEM results reported by Shahbzian-Yassar.24 There are two possible reasons that may lead to the less structural changes of Na/PCO cell. One is due to there may be very few insertion of Na+ into the PCO; the other one may be due to Na/PCO cell followed a non-conversion reaction process, e.g. intercalation. However, in any of the above cases, the capacity fading of the PCO anode for SIBs should not come from the volumetric changes that were usually blamed for poor cycle stability. In order to gain more information about the structural changes of Li/PCO cell and Na/PCO cell during cycling, in operando SXRD was further used to monitor the crystal structure evolution of the PCO anode. Figure 4a shows the in operando SXRD data of the Li/PCO cell along with its initial discharge/charge curve at 0.1 C. As clearly shown, the intensities for the diffraction peaks of the PCO were significantly decreased along with the discharge process and completely transformed to amorphous phase when the cell was discharged to 0.01 V. It was found that the crystalline structure of the PCO cannot be recovered during the charge process, which should be 13


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owing to the amorphous state of the recharged product as most of the capacity could be recovered (see its corresponding discharge/charge curve). This will be further confirmed by the in operando XAFS measurement of Li/PCO cell as discussed later. While in the case of Na/PCO cell, it shows a complete different structure evolution process. The intensities for the diffraction peaks of the PCO anode have only a slight decrease and we do not observe any peak shifts that will appear when intercalation process occurred. It means that the minor pore structure changes of Na/ PCO cell are attributed to the minor electrochemical reactions between Na and PCO. At the end of charge process, it can be seen that the intensity cannot be recovered, indicating a large irreversible capacity loss as shown in its cycle performance. To further get the structural changes especially those amorphous phases that may form during the discharge/charge process of Li/PCO and Na/PCO cells, in operando Co K-edge XANES measurement were carried out as it is useful for determining the average chemical environment of the elements in all phases involves liquid and solid, as well as crystalline and/or amorphous phases. Figure 5a shows the 2D contour plot of in operando Co K-edge XANES data of Li/PCO cell in the initial two cycles at 0.2 C. The color reflects the normalized absorption intensity, the red color represents high intensity, and the blue color means low intensity. In addition, the voltage profile in the first two cycles is also displayed and correlated with the evolution of the XANES data. The intensities of the signals belong to the pristine PCO was gradually decreased and shifted to a lower energy position along with the discharge process, which may be corresponding to the reduction of cobalt oxide by lithium to form metallic cobalt and Li2O. It can be also seen that the signals at the end of the 1st charge are different from the pristine PCO electrode, indicating that the recharged product are not the pristine PCO. After that, the apparent symmetry of the dataset emphasizes that the discharging and charging proceed through a reversible sequence of reactions in the subsequent cycle. The typical Co K-edge XANES spectra of the PCO electrode at different discharge/charge states along with references of Co foil, CoO and Co3O4 are shown in Figure 5b. The Co K-edge XANES spectrum of cobalt oxide can be generally divided into two parts namely a pre-edge and a white line, corresponding to 1s to 3d transition 14


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and 1s to 4p transition, respectively.35,

44

The pre-edge peak is normally

dipole-forbidden, but becomes allowed when the metal has a tetrahedral coordination. The XANES spectrum of CoO shows a minor pre-edge peak which signifies an octahedral configuration of O around Co, and a sharp white line at 7725 eV. On the contrary, the XANES spectrum of Co3O4 exhibits a significant pre-edge peak at 7710 eV, again due to a tetrahedral configuration of O around Co, and a white line at 7730 eV. The Co K-edge absorption energy position for the pristine PCO contains the features belong to both CoO and Co3O4 (show more clearly in Figure S3a), indicating that the pristine PCO is a mixture of CoO and Co3O4, coinciding with the results from HEXRD and TEM. The spectra of the electrodes in the 1st and 2nd discharged state exhibit a shoulder peak at around 7713 eV with a decrease for the intensity of the white line (show more clearly in Figure S3b), which are similar to that of standard Co foil, indicating that the PCO has been completely transformed into metallic Co nano-domains when the cell was discharged to 0.01 V. The shift of Co K-edge to a lower energy during lithiation may be due to the decreasing valence state of Co. While at the end of the 1st and 2nd charge, it’s obvious that the spectra do not re-assemble that of the pristine PCO (show more clearly in Figure S3c), which present a pre-edge peak clearly visible at 7710 eV and a slight rise of a white line at 7730 and 7725 eV, characteristic features of Co3O4 and CoO phases, respectively. Combined with the result obtained from in operando SXRD study, it can be known that the recharged product for Li/PCO cell is a mixture of amorphous CoO and Co3O4. Noted that the PCO is a mixture of crystalline CoO and Co3O4, the weight ratio between CoO and Co3O4 in the PCO and the recharged product should be different. Figure 5c shows the 2D contour plot of in operando XANES data of Na/PCO cell at 0.1 C in the initial two cycles. It can be also seen that the intensities were slightly decreased until the voltage reached the flat plateau, indicating that the reduction of PCO by Na occurred at around 0.2 V. However, by comparison of the whole dataset during repeated discharge and charge process of Na/PCO cell, it’s obvious that the structural changes of Na/PCO cell are much smaller than that of Li/PCO cell. Figure 5d shows typical XANES spectra of Na/PCO cell at different discharge/charge states. As can be 15


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seen, there are obviously no formation of metallic cobalt nano-domains at the end of the 1st and 2nd discharge since their spectra are significantly different from the Co foil (show more clearly in Figure S3d). Compared to the pristine PCO spectrum, it can be known that the PCO was only slightly reduced. During the 1st and 2nd charge process, the spectra cannot be recovered to the pristine PCO either. The spectra lied between CoO and Co3O4, indicating that the average valence state of Co for the charged products is between +2 and +8/3, further confirming the minor electrochemical reactions between Na and PCO. In order to get the local structures changes of Co in the PCO, in operando EXAFS measurements were also conducted. The Fourier transforms (FT) of the EXAFS spectra for Li/PCO and Na/PCO cell in the course of the 1st discharge/charge process along with CoO, Co3O4 and Co foil references were shown in Figure 5e and Figure 5f, respectively. The FT of EXAFS spectrum of the pristine PCO shows three peaks at around 1.5, 2.5 and 3.0 Å, covering the information belong to CoO and Co3O4. Upon the decrease of the voltage during discharge, the intensities of the peaks belong to CoO and Co3O4 were gradually decreased accompanied by the increase of the signals belong to Co foil. At the end of the 1st discharge, it can be seen that the FT of EXAFS spectrum is very close to that of Co foil in the reference, indicating a phase transition from PCO to Co metal. During the charge process, the signals of Co foil were gradually decreased while the signals belong to CoO and Co3O4 were increased correspondingly. At the end of the 1st charge, it can be found that the FT of the EXAFS spectrum do not resile to that of the pristine PCO, but still contains the information of both Co3O4 (1.06 Å) and CoO (1.48 Å), in agreement with the result of in operando XANES. In addition, it can be also seen that there are a weak and broad peak at about 2.11 Å, indicating an existence of Co nano-domains. This may be owing to partial Li2O was not reduced by Co at the end of 1st charge, which will be decomposed during cycling and hence lead to the capacity increasing in the initial few cycles. While in the case of Na/PCO, it can be obviously seen that the original signals have only a slight decrease accompanied by a minor increase of Co metal at around 2.0 Å and then remain almost unchanged during the whole charge process. 16


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Therefore, by a careful examination on the Li/PCO and Na/PCO cell during the charge/discharge process using in operando SAXS, SXRD and XAFS techniques, it was found that there are less pore structure changes, crystal structure evolution and valence states as well as local structural changes for Na/PCO cell compared to the Li/PCO cell during cycling, which is actually attributed to the less electrochemical reactions between Na and PCO. Similar result was also reported in the crystalline Ge nanowire, which is also Na inactive. Kohandehghan and co-workers reported that activation with Li induces amorphization in Ge nanowires could reduce the barrier for nucleation of the NaxGe phase(s) and accelerates solid-state diffusion that aids the performance.22 The in operando SXRD of Li/PCO cell during charge/discharge has revealed that the recharged product became amorphous at the end of the 1st charge process. Therefore, Li/PCO cell was disassembled after a single (de)lithiation cycle at 0.1 C and then re-assembled into Na/PCO cell. However, as shown in Figure S4a, even charge/discharged at a lower rate (0.05C), the reversible capacity and the cycle stability of Na/Li-activated PCO cell was not obviously improved as that reported in the case of Ge anode for SIBs,22,

23

indicating that the crystallinity does not

significantly affect the sodiation activity of PCO. The temperaure effect was also evaluated by charging/discharging Na/PCO cell at 55 oC. However, the PCO anode still shows rapid capcity fading (Figure S4b), revealing that the low sodiation capability of PCO was also not temperature-limited. Therefore, the poor sodium storage performance of PCO may come from the intrinsic large sodiation barrier of cobalt oxide, which significantly lower its sodiation activity. Na+ has a much larger radius of 1.02 Å than Li+, which may hinder the diffusion of Na+ into the bulk structure of Co3O4 and further the conversion reaction between Na and Co3O4. Therefore, in order to facilitate the transportation of Na+ and drive the conversion reaction, the host materials may need a large anion framework such as metal sulfides owing to the larger radius of S2- (1.84 Å) than O2- (1.40 Å). In order to confirm our hypothesis, we performed ab initio molecular dynamics (MD) simulations to reveal the lithiation and sodiation activity of Co3O4 and Co3S4. Figure 6a, 6b and 6c shows the structural evolution processes of lithiated Co3O4, 17


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sodiated Co3O4 and sodiated Co3S4 as a function of ab initio MD simulations time, respectively. The result shows that compared to the lithiation process, the Co3O4 undergo less atomic distortion during the sodiation proces. In another word, less conversion reaction between Na and Co3O4 occurred, indicating a poor sodiation activity of Co3O4. While the atomic distortion of sodiated Co3S4 was found to be larger than that of sodiated Co3O4, suggesting more conversion reacton occurred. Figure 6d further compares the degree of Li/Na oxidation for lithiatied Co3O4, sodiated Co3O4 and sodiated Co3S4 as a function of ab initio MD simulations time, in which the degree of oxidation is defined as the ratio between the calculated positive charges in Li or Na and the theoretical positive charges when they were fully oxidized and can be used as indicator for lithiation/sodiation activity. The theoretical positive charges of a Li+ or Na+ ion were calculated from the Li2O, Na2S and Na2O unit cells respectively, as listed in Table S2. While the calculated positive charges in Li or Na metals are listed in Table S3. Based on these calculations, it be found that the degree of oxidation are all graudally increased along with the MD simulation time. However, the degree of oxidation for sodiated Co3O4 after 2 ps is much lower than that of lithiated Co3O4 and only interfacial 36.4% Na atoms are converted into Na+ ions while other Na atoms maintain their metallic states, corresponding to a poor sodiation activity. Instead, the the degree of oxidation for lithiated Co3O4 is as high as 90% after 2 ps, indicating that Co3O4 is highly Li-active. This is very consistent with the electrochemical results and the corresponding structural changes observed during charge/discharge by in operando synchrotron X-ray techniques. Moreover, it’s very inspiring to see that Co3S4 indeed demonstrate better sodiation activity than Co3O4 considering its higher degree of oxidation (ca. 60% after 2 ps). In addition to the perspective from MD simulations, the previous reported lithium/sodium storage performances of metal oxides and metal sulfides indeed further supported our hypothesis. As shown in Table 1, the Co3O4 and Fe2O3 could both showed very high capacity retention during the (de)lithiation process. However, their capacity retentions during the (de)sodiation process were significantly decreased to lower than 50%. While the metal sulfides such as FeS2 and MoS2 indeed demonstrated better cycle 18


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stability even at a higher charge/discharge rate. Future mechanism studies to uncover more underline physics between lithium-ion batteries and sodium-ion batteries are highly recommended so that a rational design of high-performance sodium-ion anode can be really enabled.

Conclusions In this work, we performed a comparative understanding on the pore structure changes, crystal structure evolution, oxidation state as well as local structure changes of a porous cobalt oxide anode material during its electrochemical lithiation and sodiation process by using in operando SAXS, SXRD and XAFS. It was found that there are less structural changes during the sodiation process, which is owing to the intrinsic low sodiation activity of cobalt oxide as evidenced by ab initio MD simulations, and thus lead to its inferior sodium storage performance than lithium storage performance. Moreover, it was found that metal sulfides show a better sodiation activity than metal oxides and thus are better candidates as high performance electrode materials of SIBs. Such understanding is crucial for future design and improvement of high performance electrode materials for SIBs.

Appendix A. Supporting materials Supporting Information Available: Tables for the Rietveld refinement parameters of the porous cobalt oxide and the DFT calculation parameters. TEM characterization on the porous cobalt oxide. In-situ EIS of Na/CoOx cell in the first 10 cycles. XANES spectra of Li/CoOx and Na/CoOx cells at different charge/discharge states. Cycle performance of Na/CoOx cell after a single (de)lithiation cycle and Na/CoOx cell at 55 o

C. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author *To whom correspondence should be addressed. Zonghai Chen: Email: [email protected] Khalil Amine: Email: [email protected] 19


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Notes The authors declare no competing financial interest.

Acknowledgement Research at the Argonne National Laboratory was funded by U.S. Department of Energy, Vehicle Technologies Office. Support from Tien Duong of the U.S. DOE‫׳‬s Office of Vehicle Technologies Program is gratefully acknowledged. Use of the Advanced Photon Source, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science by Argonne National Laboratory, was supported by the U.S. DOE under Contract No. DE-AC02-06CH11357. Sector 20 facilities at the Advanced Photon Source, and research at these facilities, are supported by the US Department of Energy Basic Energy Sciences, the Canadian Light Source and its funding partners, the University of Washington, and the Advanced Photon Source. Research at State Key lab of Xiamen University was funded by National Natural Science Foundation of China (Grants No. 21321062).

Notes and Reference 1. Xu, G. L.; Wang, Q.; Fang, J. C.; Xu, Y. F.; Huang, L.; Sun, S. G. J. Mater. Chem. A 2014, 2, 19941-19962. 2. Xu, G. L.; Chen, Z.; Zhong, G. M.; Liu, Y.; Yang, Y.; Ma, T.; Ren, Y.; Zuo, X.; Wu, X. H.; Zhang, X.; Amine, K. Nano Lett. 2016, 16, 3955-3965. 3. Zhu, Y.; Wen, Y.; Fan, X.; Gao, T.; Han, F.; Luo, C.; Liou, S.-C.; Wang, C. Acs Nano 2015, 9, 3254-3264. 4. Luo, W.; Shen, F.; Bommier, C.; Zhu, H.; Ji, X.; Hu, L. Acc. Chem. Res. 2016, 49, 231-240. 5. Fang, C.; Huang, Y.; Zhang, W.; Han, J.; Deng, Z.; Cao, Y.; Yang, H. Adv. Energy Mater. 2016, 6, 501727. 6. Fang, Y.; Xiao, L.; Qian, J.; Ai, X.; Yang, H.; Cao, Y. Nano Lett. 2014, 14, 3539-3543. 7. Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Chem. Rev. 2014, 114, 11636-11682. 8. Douglas, A.; Carter, R.; Oakes, L.; Share, K.; Cohn, A. P.; Pint, C. L. Acs Nano 2015, 9, 11156-11165. 9. Lotfabad, E. M.; Ding, J.; Cui, K.; Kohandehghan, A.; Kalisvaart, W. P.; Hazelton, M.; Mitlin, D. Acs Nano 2014, 8, 7115-7129. 10. Zhu, C. B.; Mu, X. K.; van Aken, P. A.; Yu, Y.; Maier, J. Angew. Chem. Int. Ed. 2014, 53, 2152-2156. 11. Li, C.; Yin, C.; Mu, X.; Maier, J. Chem. Mater. 2013, 25, 962-969. 12. Zhao, Y.; Manthiram, A. Chem. Mater. 2015, 27, 3096-3101. 13. Wang, H.-G.; Yuan, S.; Ma, D.-L.; Zhang, X.-B.; Yan, J.-M. Energy Environ. Sci. 2015, 8, 1660-1681. 20


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14. Liu, J.; Wen, Y. R.; van Aken, P. A.; Maier, J.; Yu, Y. Nano Lett. 2014, 14, 6387-6392. 15. Xu, Y.; Zhu, Y.; Liu, Y.; Wang, C. Adv. Energy Mater. 2013, 3, 128-133. 16. Xu, Y.; Swaans, E.; Basak, S.; Zandbergen, H. W.; Borsa, D. M.; Mulder, F. M. Adv. Energy Mater. 2015, 6, 1501436. 17. Wu, Z. G.; Zhong, Y.-J.; Liu, J.; Wu, J. H.; Guo, X. D.; Zhong, B. H.; Zhang, Z. Y. J. Mater. Chem. A 2015, 3, 10092-10099. 18. Wen, J.-W.; Zhang, D.-W.; Zang, Y.; Sun, X.; Cheng, B.; Ding, C.-X.; Yu, Y.; Chen, C.-H. Electrochim. Acta 2014, 132, 193-199. 19. Jiang, Y.; Hu, M.; Zhang, D.; Yuan, T.; Sun, W.; Xu, B.; Yan, M. Nano Energy 2014, 5, 60-66. 20. Fu, Y.; Wei, Q.; Wang, X.; Zhang, G.; Shu, H.; Yang, X.; Tavares, A. C.; Sun, S. Rsc Adv. 2016, 6, 16624-16633. 21. Lu, X.; Adkins, E. R.; He, Y.; Zhong, L.; Luo, L.; Mao, S. X.; Wang, C.-M.; Korgel, B. A. Chem. Mater. 2016, 28, 1236-1242. 22. Kohandehghan, A.; Cui, K.; Kupsta, M.; Ding, J.; Memarzadeh Lotfabad, E.; Kalisvaart, W. P.; Mitlin, D. Nano Lett. 2014, 14, 5873-5882. 23. Jung, S. C.; Kim, H. J.; Choi, J. W.; Han, Y. K. Nano Lett. 2014, 14, 6559-6563. 24. Asayesh-Ardakani, H.; Nie, A.; Mashayek, F.; Klie, R. F.; Shahbazian-Yassar, R. Microsc. Microanal. 2015, 21, 1371-1372. 25. Xu, X.; Yan, M.; Tian, X.; Yang, C.; Shi, M.; Wei, Q.; Xu, L.; Mai, L. Nano Lett. 2015, 15, 3879-3884. 26. He, K.; Lin, F.; Zhu, Y.; Yu, X.; Li, J.; Lin, R.; Nordlund, D.; Weng, T.-C.; Richards, R. M.; Yang, X.-Q.; Doeff, M. M.; Stach, E. A.; Mo, Y.; Xin, H. L.; Su, D. Nano Lett. 2015, 15, 5755-5763. 27. Mayo, M.; Griffith, K. J.; Pickard, C. J.; Morris, A. J. Chem. Mater. 2016, 28, 2011-2021. 28. Xu, G. L.; Ma, T.; Sun, C. J.; Luo, C.; Cheng, L.; Ren, Y.; Heald, S. M.; Wang, C.; Curtiss, L.; Wen, J.; Miller, D. J.; Li, T.; Zuo, X.; Petkov, V.; Chen, Z.; Amine, K. Nano Lett. 2016, 16, 2663-2673. 29. Xu, G.-L.; Li, Y.; Ma, T.; Ren, Y.; Wang, H.-H.; Wang, L.; Wen, J.; Miller, D.; Amine, K.; Chen, Z. Nano Energy 2015, 18, 253-264. 30. Cui, Y. J.; Abouimrane, A.; Lu, J.; Bolin, T.; Ren, Y.; Weng, W.; Sun, C. J.; Maroni, V. A.; Heald, S. M.; Amine, K. J. Am. Chem. Soc. 2013, 135, 8047-8056. 31. Park, G. O.; Yoon, J.; Park, E.; Park, S. B.; Kim, H.; Kim, K. H.; Jin, X.; Shin, T. J.; Kim, H.; Thon, W.-S.; Kim, J. M. Acs Nano 2015, 9, 5470-5477. 32. Shon, J. K.; Lee, H. S.; Park, G. O.; Yoon, J.; Park, E.; Park, G. S.; Kong, S. S.; Jin, M.; Choi, J. M.; Chang, H.; Doo, S.; Kim, J. M.; Yoon, W. S.; Pak, C.; Kim, H.; Stucky, G. D. Nat. Commun. 2016, 7, 11049. 33. Chong, L.; Goenaga, G. A.; Williams, K.; Barkholtz, H. M.; Grabstanowicz, L. R.; Brooksbank, J. A.; Papandrew, A. B.; Elzein, R.; Schlaf, R.; Zawodzinski, T. A. ChemElectroChem 2016, 3, 1-6. 34. Xu, G. L.; Li, J. T.; Huang, L.; Lin, W. F.; Sun, S. G. Nano Energy 2013, 2, 394-402. 35. Kim, H.; Kim, H.; Kim, H.; Kim, J.; Yoon, G.; Lim, K.; Yoon, W.-S.; Kang, K. Adv. Funct. Mater. 2016, 26, 5042-5050. 36. Poizot, P.; Laruelle, S.; Grugeon, S.; Dupont, L.; Tarascon, J. M. Nature 2000, 407, 496-499. 21


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Figure 1. (a) Rietveld refinement of HEXRD, (b) TGA curve at air atmosphere, (c) N2 adsorption/de-sorption isotherm curve and (d) SEM image of the as-prepared PCO.

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a 3.5 +

)

3.0

2.0 1.5 1.0 0.5

100 80

1500

-1

2.5

Capacity, mAh g

Voltage, V (vs. Li/Li

b 2000

1st 2nd 40th

60 1000

0.1C

40

Charge Discharge Coulombic efficiency

500

0.0

20

0 0

200

400

600

800

1000 1200 1400

0 0

-1

10

Capacity, mAh g 3.0

+

1.5 1.0 0.5

40

100

800

2.0

30

Cycle number

1000

80

Charge Discharge Coulombic efficiency

-1

2.5

Capacity, mAh g

Voltage, V (vs. Na/Na

d

1st 2nd 40th

20

600 400

60 40

0.1C

200

20

0.0

0 0

200

400

600

800

1000 1200 1400

0 0

-1

Capacity, mAh g

Coulombic efficiency, %

)

c

3.5

Coulombic efficiency, %

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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10

20

30

40

Cycle number

Figure 2. (a) Discharge/charge curves and (b) cycle performance of Li/PCO cell at 0.1 C (1C=890 mA g-1); (c) Discharge/charge curves and (d) cycle performance of Na/PCO cell at 0.1 C (1C=890 mA g-1).

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Figure 3. (a) 2D contour plot of in operando SAXS data along with the 1st discharge/charge curve at 0.1 C and (b) typical SAXS data at different charge/discharge states of Li/PCO cell; (c) 2D contour plot of in operando SAXS data along with the 1st discharge/charge curve at 0.1 C and (d) typical SAXS data at different charge/discharge states of Na/PCO cell. The color scale bar means the intensity of the SAXS spectra, red means high, blue means low.

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Figure 4. In operando SXRD patterns along with the 1st discharge/charge curve at 0.1 C of (a) Li/PCO cell and (b) Na/PCO cell.

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Figure 5 (a) 2D contour plot of in operando XANES data along with its discharge/charge curves at 0.2 C and (b) typical XANES spectrum at different charge/discharge states of Li/PCO cell; (c) 2D contour plot of in operando XANES data along with its discharge/charge curves at 0.2 C and (d) typical XANES spectrum at different charge/discharge states of Na/PCO cell; (e) Fourier Transform of in operando EXAFS data along with its discharge/charge curves of (e) Li/PCO cell and (f) Na/PCO cell.

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Figure 6. The structure evolution of (a) lithiated Co3O4, (b) sodiated Co3O4 and (c) sodiated Co3S4 as a function of MD simulations time; d) the degree of oxidation of Li/Na as a function of MD simulation time for lithiated Co3O4, sodiated Co3O4 and sodiated Co3S4. The grey and purple atoms represent Li and Na, respectively. The red atom means Oxygen, the blue atom means Cobalt and the yellow atom means Sulfur.

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Table 1. Summarizes on the lithium and sodium storage performances of metal oxides and metal sulfides Initial, mAh g-1

Last, mAh g-1

Rate, A g-1

Capacity retention

Li

877

1441.1

0.2

164%@50 cycles

Na

611.2

290

0.2

47.4%@50 cycles

Li

1050

1440.9

0.4

137%@200 cycles

Na

671.9

288

0.04

42.8%@70 cycles

Li

816

672

0.5

82.3%@200 cycles

Na

463

329

0.5

71%@50 cycles

Li

800

630

0.2

78.7%@100 cycles

Na

820

600

1

73.1%@100 cycles

Li

886

660

0.1

75%@100 cycles

Na

640

512

0.25

80%@200 cycles

Co3O4

Fe2O3

Co3S4

FeS2

MoS2

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46

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Sodiation of ... - ACS Publications

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