Silver-Containing α-MnO2 Nanorods: Electrochemistry in Na-Based

Sep 1, 2016 - ... Structural Evolution of Ag1.2Mn8O16 under Coulombically Controlled (De)Lithiation. Jianping Huang , Xiaobing Hu , Alexander B. Brady...
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Silver-Containing α‑MnO2 Nanorods: Electrochemistry in Na-Based Battery Systems Jianping Huang,† Altug S. Poyraz,‡ Seung-Yong Lee,‡ Lijun Wu,‡ Yimei Zhu,‡ Amy C. Marschilok,*,†,§ Kenneth J. Takeuchi,*,†,§ and Esther S. Takeuchi*,†,‡,§ †

Department of Chemistry and §Department of Materials Science and Engineering, Stony Brook University, Stony Brook, New York 11794, United States ‡ Energy Sciences Directorate, Brookhaven National Laboratory, Upton New York 11973, United States S Supporting Information *

ABSTRACT: Manganese oxides are considered attractive cathode materials for rechargeable batteries due to the high abundance and environmental friendliness of manganese. In particular, cryptomelane and hollandite are desirable due to their ability to host cations within their octahedral molecular sieve (OMS-2) α-MnO2 structure. In this work, we investigate silver containing α-MnO2 structured materials (AgxMn8O16, x = 1.22, L-Ag-OMS-2 or 1.66, H-Ag-OMS-2) as host materials for Li ion and Na ion insertion/deinsertion. The results indicate a significant difference in the lithiation versus sodiation process of the OMS-2 materials. Initial reduction of Ag1.22Mn8O16 to 1.0 V delivered ∼370 mAh/g. Cycling of Ag1.22Mn8O16 between voltage ranges of 3.8−1.7 V and 3.8−1.3 V in a Na battery delivered initial capacities of 113 and 247 mAh/g, respectively. In contrast, Ag1.66Mn8O16 delivered only 15 mAh/ g, ∼ 0.5 electron equivalents, to 1.7 and 1.3 V. Study of the system by electrochemical impedance spectroscopy (EIS) showed a significant decrease in charge transfer resistance from 2029 Ω to 594 Ω after 1.5 electron equivalents per Ag1.22Mn8O16 formula unit of Na ion insertion. In contrast, both Ag1.22Mn8O16 and Ag1.66Mn8O16 exhibited gradual impedance increases during lithiation. The formation of silver metal could be detected only in the sodiated material by X-ray diffraction (XRD). Thus, the impedance of Ag-OMS-2 decreases upon sodiation coincident with the formation of silver metal during the discharge process, consistent with the more favorable formation of silver metal during the sodiation process relative to the lithation process. KEYWORDS: alpha manganese oxide, sodium battery, silver hollandite, electrochemistry, X-ray diffraction ∼80% retention at the 30th cycle.19 Moreover, the estimated energy density of this material was 520 mWh/g comparable to the ∼530 mWh/g delivered by LiFePO4 in Li ion cells.19 Additionally, several materials with tunnel structures have also been studied. For example, a 4 × 2 tunnel structure material, orthorhombic Na0.44MnO2, was considered as a suitable Na host material due to facile Na ion diffusion through its 1dimensional type tunnels.20,21 This material exhibited capacities of 140 mAh/g and a reversible intercalation/deintercalation process occurred over the Na composition range of x = 0.25− 0.65 for NaxMnO2.20,22 The results from manganese based materials motivated the research on manganese dioxides with a rigid structural framework.23,24 For example, α-MnO2, a 2 × 2 tunnel structure, showed an initial discharge capacity (278 mAh/g, vs Na/Na+), but the capacity retention was only ∼28%

1. INTRODUCTION Lithium ion batteries have been highly successful energy storage devices due to their high operating voltage, high energy density and long cycle life.1−3 However, market demands for large scale batteries including electric vehicles and hybrid electrical vehicles4,5 have raised potential concerns about battery cost and the potential scarcity of lithium resources. Therefore, numerous researchers are investigating generations of batteries beyond lithium ion batteries, including Na based,6,7 Mg based,8 Li/S9,10 and Li−air batteries.10,11 Among the candidates, Na-based batteries may offer appeal because of the high natural abundance of sodium and intercalation mechanisms reminiscent of Li.6 Notably, as the Na ion (1.02 Å)12 is larger than Li ion (0.76 Å)12 the energy storage system requires host materials with facile diffusion pathways for the larger sodium ion. Several attempts to seek suitable cathode materials for Na based batteries have focused on layered materials, such as Na x MnO 2 , 13,14 NaFeO 2 , 15 Na x CoO 2 , 16 NaCrO 2 , 17 and NaxVO2.18 A recent report proposed P2−Na2/3[Fe1/2Mn1/2]O2, and this material showed an initial capacity of 190 mAh/g with © XXXX American Chemical Society

Special Issue: New Materials and Approaches for Beyond Li-ion Batteries Received: July 13, 2016 Accepted: August 26, 2016

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ACS Applied Materials & Interfaces after 100 cycles.24 Theoretical studies suggested that a significant unit cell distortion occurs when 0.5 Na intercalates into α-MnO2, and distortion of the 2 × 2 tunnel leads to an increase in the ratio of a and b lattice parameters (a/b) from 1 to 1.28.25 Silver hollandite (AgxMn8O16), a structural analogue to αMnO2, has been previously investigated as a cathode material for Li batteries.26−28 Silver ions are located in the middle of 2 × 2 tunnels and modify the a,b dimensions of the unit cell. Due to the covalent nature of silver ion, lattice a decreases from 9.770 to 9.738 Å when the silver content increases from 1.22 to 1.66.29 Notably, the 2 × 2 tunnel dimension for α-MnO2 starts from ∼9.79 Å (α-MnO2),30 and is further expanded when occupied by potassium ions where lattice is 9.866 Å (K1.33Mn8O16)31 illustrating the significant impact of the central cation. Notably, the Ag+ cations in the tunnels of AgxMn8O16 or Ag-OMS-2, are potentially electrochemically active and may be reduced to silver metal during the reduction process. In-situ silver metal formation during the discharge process may provide an increase in material conductivity as seen in some layered materials.32−40 In this report, we investigate the electrochemistry of AgxMn8O16 with sodium as the anode material where x = 1.22 or 1.66. Galvanostatic cycling, electrochemical impedance spectroscopy, and galvanostatic intermittent titration type testing were used to characterize performance, where behavior in lithium based electrochemical cells and sodium-based electrochemical cells was compared and contrasted. The electrochemical discharge processes were probed via X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), transmission electron microscopy (TEM), and X-ray absorption spectroscopy (XAS) to elucidate the electrochemical mechanism.

processed utilizing the Athena software packages, and E0 value (6.539 keV) from Mn foil reference was used for all the samples. 2.3. Electrochemical Characterization. A mixture of silver hollandite, graphite, carbon black and polyvinylidene fluoride (PVDF) was tape cast on aluminum foil and used as the cathode. The electrolyte for Na battery was 1 M NaPF6 in ethylene carbonate (EC)/ diethyl carbonate (DEC) at a volume ratio of 1:1. The electrolyte for the Li battery was 1 M LiPF6 in ethylene carbonate/dimethyl carbonate at a volume ratio of 3:7. Coin cells were assembled using the cathode, Na foil or Li foil, polypropylene separators and the electrolyte in an argon filled glovebox. Galvanostatic cycling tests were conducted at 30 °C under a current density of 20 mA/g. Rate capability tests were performed at 10 (0.1 C), 50, 100, 200, and 10 mA/g for 10 cycles each. Electrochemical impedance spectroscopy (EIS) was measured using a BioLogic VSP impedance analyzer with a 10 mV amplitude and a frequency range of 100 kHz to 10 mHz. For EIS measurements of Na batteries, the cells were discharged at a current density of 10 mA/g for 0.5 h followed by a rest time of 4 h, and EIS data was then recorded in the beginning of each discharge pulse. In Li batteries, the cells were discharged at a current density of 10 mA/g for 2 h followed by a rest time of 22 h. ZView software was utilized to analyze the EIS data.

3. RESULTS AND DISCUSSION 3.1. Materials Synthesis and Characterization. As prepared silver hollandite materials were characterized by XRD, and all the diffraction peaks can be assigned to the tetragonal phase of Ag1.8Mn8O16 (PDF # 01−077−1987) with no apparent impurities, Figure 1. The inset in Figure 1 shows

2. EXPERIMENTAL SECTION 2.1. Materials Synthesis. Silver hollandite samples were synthesized by an ambient pressure reflux method according to a previously reported method.26,27 Silver permanganate (AgMnO4), manganese monohydrate (MnSO4·H2O) and nitric acid were the starting materials used for the synthesis. The solution was heated to reflux, and after isolation, the precipitates were then washed with DI water several times. After synthesis, the samples were annealed at 300 °C for 6 h to remove physisorbed and structural water. 2.2. Materials Characterization. X-ray diffraction (XRD) patterns were collected using a Rigaku SmartLab X-ray diffractometer (Cu Kα radiation, λ = 1.5406 Å) with Bragg−Brentano configuration. The Ag/Mn ratio in the products was identified using inductively coupled plasma spectrometry-optical emission spectroscopy (ICPOES) with a Thermofisher iCAP 6300 series instrument. The water content and thermal stability of the materials were analyzed using thermogravimetric analysis (TGA) with a TA Instruments SDT Q600 under nitrogen gas atmosphere. X-ray photoelectron spectroscopy (XPS) experiments were performed with a UHV chamber equipped with the SPECS Phoibos 100 MCD analyzer and an X-ray source (hν = 1486.6 eV) operating with an accelerating voltage of 10 kV and a current of 30 mA. The samples were mounted onto a sample holder using double-sided copper foil or carbon tape for measurement. All the spectra were calibrated using the binding energy of C 1s peak at 284.1 eV.41,42 High-resolution TEM/STEM images were obtained from the double aberration-corrected JEOL-ARM 200CF microscope with a cold-field emission gun operated at 200 kV. The microscope is equipped with JEOL and Gatan HAADF detectors for incoherent HAADF (Z-contrast) imaging. Mn K-edge X-ray absorption spectra (XAS) were collected at APS 12-BM-B at the Argonne National Laboratory. All the samples were sealed between polyimide tape for measurement. The XAS data were

Figure 1. X-ray diffraction pattern of annealed L-Ag-OMS-2 (black) and H-Ag-OMS-2 (red). Inset: crystal structure of silver hollandite, MnO6: purple, Ag: gray, O: red.

the crystal structure of silver hollandite viewed along the c-axis where 2 × 2 tunnels are formed by the interconnection of edge sharing MnO6 octahedra, and silver ions residing in the center of the tunnels. The crystallite size for each sample was estimated from the XRD patterns using Scherer equation,43 where the (211) diffraction peak was selected for size analysis. The calculated crystallite size was 10.8 nm for L-Ag-OMS-2 and 22.2 nm for H-Ag-OMS-2. The Ag content in AgxMn8O16 samples was determined by the use of ICP-OES, and was found to be 1.22 and 1.66, close to the lower limit and upper limit of the range reported previously.27 After annealing, both materials maintained a small amount of physisorbed water as the TGA data showed weight losses of 1−1.4% below 120 °C, Figure S1. It has been reported that materials with this structure type reabsorbed water very quickly after dehydration,44 so the weight loss below 120 °C was probably due to the reabsorbed water. On the basis of the TGA results, we assign formulas of Ag1.22Mn8O16·0.9H2O and Ag1.66Mn8O16·0.8H2O to the two materials studied. A designation previously utilized for α-MnO2, B

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Figure 2. (a) Cycling performance of L-Ag-OMS-2 (black) and H-Ag-OMS-2 (red) in Na batteries. Inset: discharge/charge curves of Ag-OMS-2 at cycle 10 for two voltages windows. (b) Cycling performance of L-Ag-OMS-2 (black) and H-Ag-OMS-2 (red) in Li batteries. (c) Rate capabilities of L-Ag-OMS-2 in Na (red) and Li batteries (black). Inset: Comparison of L-Ag-OMS-2 in Na (red) and Li batteries (black). (d) Capacity retention at different C-rates in Na (red) and Li batteries (black).

the hollandite structure with no central cation, is ‘octahedral molecular sieve 2’ (OMS-2), 45−47 thus for simplicity, Ag1.22Mn8O16·0.9H2O is denoted here as L-Ag-OMS-2 and Ag1.66Mn8O16·0.8H2O is denoted as H-Ag-OMS-2. 3.2. Electrochemistry. L-Ag-OMS-2 materials were discharged at a current density of 5 mA/g in Na cells and showed several changes in slope for the voltage profile, Figure S2. When the voltage reached 1.0 V, ∼ 11 electron equivalents of Na ions per formula unit (about 368 mAh/g) were inserted, indicating the ability of Ag-OMS-2 to achieve high discharge capacity in Na batteries. Voltage windows of 3.8−1.7 V and 3.8−1.3 V were used to investigate the reversibility of Ag-OMS-2 in Na batteries at a current density of 20 mA/g, Figure 2a. In the first cycle, L-Ag-OMS-2 delivered a discharge capacity of 7.8 electron equivalents (247 mAh/g) between 3.8−1.3 V. The capacity in the 10th cycle faded to 3.3 electron equivalents (104 mAh/g). On the other hand, when the Na/L-Ag-OMS-2 cells were cycled between 3.8−1.7 V, the material was able to provide a capacity of ∼3.0 electron equivalents (∼95 mAh/g) in the initial 10 cycles, and it showed better cycle retention compared to the larger voltage window of 3.8−1.3 V. After 50 cycles, 1.9 electron equivalents (59 mAh/g) transfer was maintained for the cell cycled between 3.8 and 1.7 V, but it was only 1.2 electron equivalents (38 mAh/g) for the cell cycled between 3.8−1.3 V. It maintained 54% capacity in the voltage window of 3.8−1.7 V over 50 cycles while only 18% in 3.8−1.3 V. The inset in Figure 2a shows the discharge/charge curves of L-Ag-OMS-2 in the 10th cycle. The discharge and charge profiles of the two cells were similar, but below 2.0 V the discharge voltage dropped faster for the cell cycled in the larger voltage window, 3.8−1.3 V. Thus, the capacity retention of LAg-OMS-2 in Na batteries was significantly impacted by the voltage windows. The limited cycle stability in the Na battery may result from structural deformation at high sodiation levels as predicted by theory for α-MnO2.25

In comparison to L-Ag-OMS-2, H-Ag-OMS-2 delivered only 1.3 electron equivalents (42 mAh/g) transfer in the voltage window of 3.8−2.5 V and >2.1 electron equivalents (65 mAh/g) in 3.8−2.0 V in Li batteries. These capacity differences between L-Ag-OMS-2 and H-Ag-OMS-2 have been attributed to differences in surface defects which affect guest ion diffusion: the L-Ag-OMS-2 material has a greater concentration of oxygen vacancies at the nanorod surface, facilitating diffusion in the ab plane.29 It has also been reported that MoO3 with oxygen vacancies shows more facile sodiation due to increased electronic conductivity and Na ion diffusion.48 When H-Ag-OMS-2 and L-Ag-OMS-2 are lithiated under low rate, the difference in capacity is small, Figure S3. Thus, the difference in the electrochemistry between the two materials appears kinetic in origin consistent with oxygen vacancies affecting lithium ion diffusion in the electrochemistry. Capacity was tested under conditions where the molar equivalents of reduction were kept similar for the lithium and sodium cases. The voltage window for the lithium case was held between 3.8 and 2.5 V. In the sodium case, the selected voltage window was larger, (3.8 to 1.7 V) due to a higher amount of polarization. The capacity retention in Li/L-Ag-OMS-2 was 47% from the second cycle to the 50th cycle. The Na capacity retention was similar to that of the Li battery when the cells were cycled in the appropriate voltage windows (3.8−1.7 V in the Na battery and 3.8−2.5 V in the Li battery), Figure 2c, inset. The rate capability of L-Ag-OMS-2 was evaluated by cycling the cells at 0.1 C, 0.5 C, 1 and 2 C rates using the voltage windows as noted above of 3.8−1.7 V in the Na battery and 3.8−2.5 V in the Li battery, Figure 2c. The material delivered C

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Figure 3. (a) Ionic radii of Li+, Na+ and Ag+. (b) Discharge curves of L-Ag-OMS-2 (black) and H−Ag−OMS-2 (red) in Na batteries under galvanostatic intermittent titration technique (GITT) mode. Nyquist plots of (c) L-Ag-OMS-2 and (d) H-Ag-OMS-2 at different sodiation levels. Charge transfer resistance as a function of electron equivalents in a Na battery for (e) L-Ag-OMS-2 and (f) H-Ag-OMS-2. Inset: equivalent circuit used to fit AC impedance data.

similar discharge capacities in the initial 10 cycles at 0.1 C in the Li battery (2.5 electron equivalents, 78 mAh/g) and the Na battery (3.0 electron equivalents, 94 mAh/g). When the current density increased, stable capacity was maintained in both battery systems without significant fading. However, the capacity decreased to 1.5 electron equivalents (47 mAh/g) in the Na battery and 0.7 electron equivalent (22 mAh/g) in the Li battery at the 2 C rate. Both Na and Li batteries showed partial capacity recovery when the current density shifted from 2 to 0.1 C. Capacity retention values were calculated based on the discharge capacities in the 10th cycle of each rate, Figure 2d. Although both batteries retained ∼65% capacity at 0.5 C, the Na battery exhibited capacity retention of 61% at 1 C and 50% at 2 C. In contrast to the Na battery, the Li battery only maintained 29% capacity at 2 C. In this study, the better rate capability of L-Ag-OMS-2 in the Na battery may be associated with improved electrical conductivity during cycling due to the formation of silver metal as a result of discharge as noted below. To investigate the Na insertion effect in the materials, we employed electrochemical impedance spectroscopy (EIS) and galvanostatic Intermittent titration during the sodiation process, Figure 3. The sodiation process was controlled by a galvanostatic intermittent titration technique (GITT), and LAg-OMS-2 delivered 9.3 electron equivalents (295 mAh/g) when discharged to 1.3 V, whereas H-Ag-OMS-2 only transferred 1.8 electron equivalents (54 mAh/g), Figure 3b. This large capacity disparity is consistent with the results under constant current discharge. The impedance data were recorded

after 0.16 electron equivalent (5 mAh/g) of discharge and 4h relaxation. The Nyquist plots showed a semicircle followed by a linear line for the EIS data, indicative of a similar electrochemical environment during discharge (Figure 3c, d). It is notable that there were significant changes of the semicircles in the initial discharge processes for L-Ag-OMS-2 (below 2 electron equivalents) and H−Ag−OMS-2 (below 1 electron equivalent). Therefore, an equivalent circuit model (Inset in Figure 3e) was utilized to obtain specific impedance values. This model consists of a resistor (Rs), a parallel combination of a resistor (Rct) and a constant phase element (CPE), and a Warburg element (open, Wo) in series with Rct. Before the discharge, the cell containing L-Ag-OMS-2 showed a charge transfer resistance of 1980 Ω, but the resistance decreased to 1639 Ω when one electron equivalent of Na ion was inserted (per AgxMn8O16). A more significant resistance decrease from 2029 Ω to 799 Ω occurred at ∼1.5 electron equivalents, and the cell maintained a stable resistance value of ∼350 Ω from 3 to 9 electron equivalents. H-Ag-OMS-2 also showed a resistance drop from 3743 Ω to 577 Ω after 1 electron equivalent of discharge (Figure 3f). We also conducted a similar EIS measurements for Li cells containing L-Ag-OMS-2 and H-AgOMS-2 materials. L-Ag-OMS-2 and H-Ag-OMS-2 achieved 6.1 (194 mAh/g) and 2.1 electron equivalents (63 mAh/g) transfer in Li cells, respectively, Figure S4. Contrary to the Na cells, the Li cells displayed increased semicircle diameters in the Nyquist plots during discharge suggesting decreased conductivity, Figure 4. The impedance data were fitted with the same D

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Figure 4. Nyquist plots of (a) L-Ag-OMS-2 and (b) H-Ag-OMS-2 at different lithiation levels. Insets: charge transfer resistance as a function of electron equivalents in the Li cell.

Figure 5. X-ray photoelectron spectra (XPS) of L-Ag-OMS-2, (a) Mn 2p before and after lithiation/sodiation, (b) Na 1s at two discharge states, inset: area ratio of Na 1s peak and Mn 2p1/2 peak as a function of electron equivalents.

equivalent circuit model and the charge transfer resistance increased from 44 Ω to 129 Ω for L-Ag-OMS-2 and 129 Ω to 144 Ω for H-Ag-OMS-2 during the lithiation processes. Detailed charge transfer resistance values for the Na cells and Li cells are listed in Table S1. 3.3. Structural Analysis. X-ray photoelectron spectroscopy (XPS) was used to study the surface chemical composition and Mn oxidation state for L-Ag-OMS-2 after discharge, Figure 5. Mn 2p1/2 and Mn 2p3/2 peaks of L-Ag-OMS-2 before discharge showed binding energies at 653.8 and 642.6 eV, respectively, and both are in the reported ranges of Mn 2p binding energies for MnO249,50 (Figure 5a). This indicates that Mn valence in the sample is Mn(IV). The cells containing L-Ag-OMS-2 were discharged at a current density of 10 mA/g to two different lithiation/sodiation levels, 1 and 4 electron equivalents. For the sample discharged to 1 electron equivalent in a Na battery, the Mn 2p3/2 peak shifted to a lower energy (642.0 eV) indicating a reduced Mn valence. The shift of Mn 2p3/2 for the sample sodiated to 4 electron equivalents was more prominent, and it further decreased to 641.3 eV. The decrease of Mn oxidation state with increasing sodiation level indicates a Faradaic process occurred on the cathode. The lithiated samples showed similar Mn 2p peak shifts with a shift to 641.5 eV for Mn 2p3/2 lithiated to 4 electron equivalents. We also observed an increase in the intensity of the Na 1s peak for the sample with 4 inserted Na

ions (Na4Ag1.22Mn8O16) where the area ratio of Na 1s peak to Mn 2p1/2 peak had a linear trend with electron equivalents (inset in Figure 5b). Unlike the apparent peak shifts of Mn 2p after discharge, the shifts of Ag 3d peaks were minor for the samples at different depths of discharge, Figure S5. Notably, Na4Ag1.22Mn8O16 displayed a distinctive peak broadening implying a change of silver chemical environment. To characterize structural change during discharge and evidence for silver metal formation, the partially lithiated/ sodiated samples were analyzed by XRD, Figure 6. The samples at two depths of discharge showed clear differences in the XRD patterns. After 1 electron equivalent of discharge in the Na battery (Figure 6a), the sample (Na1Ag1.22Mn8O16) maintained two intense peaks of L-Ag-OMS-2 at 29.0° (310) and 37.4° (211), indicating only minor changes in the structure after 1 Na ion insertion. Also, there were no detectable peaks showing the presence of silver metal. For the sample discharged to 4 electron equivalents (Na4Ag1.22Mn8O16), the peak at 37.4° (211) was maintained but the peak at 29.0° (310) faded. This is consistent with the unit cell expansion in the ab plane after Na ions insertion. Prior theory work proposed a large lattice distortion (a/b = 1.28) upon sodiation of α-MnO2 (αE

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showed structural distortion as a result of lithiation or sodiation, clear evidence of silver metal by XRD only appeared in the sample reduced by sodiation. These results are consistent with the impedance data of L-Ag-OMS-2. In the Li battery, the impedance stays high and increases on discharge. In contrast, the impedance of L-Ag-OMS-2 decreases upon sodiation consistent with the formation of silver metal during the discharge process. This indicates that Na ion insertion results in more favorable formation of silver metal than Li ion insertion. The morphology and crystallinity of L-Ag-OMS-2 at different sodiation levels were investigated by high-resolution TEM, Figure 7. The nanorod-like morphology of L-Ag-OMS-2 was maintained upon sodiation up to 8 electron equivalents. However, a gradual amorphization process was observed after discharge. The as-prepared L-Ag-OMS-2 nanorods showed clear lattice fringes, and distinct spots were identified in the fast Fourier transformed (FFT) pattern indicating good crystallinity of pristine nanorods, Figure 7a. When the material was sodiated to Na2Ag1.22Mn8O16 and Na4Ag1.22Mn8O16, some level of crystallinity still remained while some defects were frequently observed in the tunnel direction and the nanorod surface was not as flat as the pristine nanorod. Spots from FFT patterns broadened and blurred. All of them imply that nanorods lose crystallinity slightly during sodiation. Small particles were also observed after sodiation, which are identified as silver metal in Figure 8. Upon sodiation to 8 electron equivalents, Na8Ag1.22Mn8O16, the tunnel structure was partially broken with lattice fringes not observed at certain areas, as indicated with yellow circles in Figure 7d. In addition, spots from the FFT patterns of these sodiated materials showed broadened and blurred, implying slightly decreased crystallinity. Na8Ag1.22Mn8O16 nanorods presented more amorphous regions without showing continuous lattice fringes. The large amorphization at high sodiation levels may result from damage to the 2 × 2 tunnels in L-Ag-OMS-2. Also, small particles were observed on the surfaces of all the sodiated L-Ag-OMS-2 nanorods. In order to determine the formation of silver metal nanoparticles, STEM high angle dark field (HAADF) images of Ag1.22Mn8O16 and Na2Ag1.22Mn8O16 were compared, Figure 8. Pristine nanorods (Ag1.22Mn8O16) had neat surface without any precipitates (Figure 8a). After 2 electron equivalents of Na ion insertion, some tiny nanoparticles (∼2 nm) were dispersed all over the nanorods (Figure 8b). Magnified STEM-HAADF image showed an apparent periodicity of atomic arrangement, which could be assigned to the cubic Ag0 phase (Figure 8c). The FFT pattern of the small particle region matched the simulated electron diffraction pattern of silver metal, verifying silver metal formation on the surface of L-Ag-OMS-2 nanorods upon sodiation. As noted above, these silver nanoparticles may contribute to enhanced electrical conductivity, consistent with the resistance decrease observed as a result of the initial sodiation process via EIS. The Mn oxidation state and local structural changes of L-AgOMS-2 during discharge were further analyzed by X-ray absorption spectroscopy and analysis of the X-ray near edge absorption structure (XANES), Figure 9. Normalized Mn Kedge XANES data showed a clear edge shifts for the samples at different discharge levels. After discharge, the Mn K-edge energies of sodiated and lithiated materials all shifted to lower energies, indicating a decrease of average Mn oxidation state in the discharged samples consistent with the XPS observation. Na1Ag1.22Mn8O16 and Li1Ag1.22Mn8O16 showed almost identical XANES spectra with edge energies of 6557.6 eV. The edge

Figure 6. X-ray diffraction patterns of 0 electron (black), 1 electron (blue), and 4 electron (red) equivalents discharged L-Ag-OMS-2 in (a) Na and (b) Li batteries.

Na0.5MnO2), which leads to the structural degradation.25 It is worth noting that two additional peaks could be clearly identified for Na4Ag1.22Mn8O16, which are assigned to (111) and (311) planes of silver metal (PDF # 01−071−4612). This result confirms that silver ions were reduced to silver metal during discharge. Notably, silver metal could be detected by XRD near the surface of the cathode after cycling, and the LAg-OMS-2 material became amorphous upon repeated lithiation/delithiation, Figure S6. When the sample was discharged in Li battery (Figure 6b), 1 electron equivalent lithiated sample (Li1Ag1.22Mn8O16) displayed similar XRD pattern as Na1Ag1.22Mn8O16. The more highly lithiated sample (Li4Ag1.22Mn8O16) also showed peak intensity fade of the (310) plane. However, no distinct silver metal peak could be identified from XRD patterns of Li1Ag1.22Mn8O16 and Li4Ag1.22Mn8O16. The ionic radii of Li+, Na+ and Ag+ are compared in Figure 3a. The radius of Na+ (1.02 Å) is larger than that of Li+ (0.76 Å), but slightly smaller than that of Ag+ (1.15 Å). OMS-2 materials have tunnel dimensions of 4.6 × 4.6 Å (2 × 2 tunnel) and 1.9 × 1.9 Å (1× 1 tunnel).51,52 For α-MnO2 with no large cation in the 2 × 2 tunnel, the 8h site is regarded as the favorable insertion site for both Li+ and Na+ in the initial insertion process.25 Interestingly, even with K+ as the central ion, the 8h site in the 2 × 2 tunnel is the preferred site of lithiation.53 Thus, with ion insertion to that site, the larger Na ion may experience an interaction with the Ag ions located in the tunnel, and facilitate the reduction-displacement of the Ag ion as silver metal. These results indicate a significant difference in the lithiation versus sodiation process of L-Ag-OMS-2. Although the material F

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Figure 7. High-resolution TEM (HRTEM) images showing crystallinity of Ag1.22Mn8O16 nanorods with different sodiation levels. Each image corresponds to the sodiation levels of (a) 0 Na (pristine), (b) 2 Na, (c) 4 Na, and (d) 8 Na per unit structure, Ag1.22Mn8O16. The insets are fast Fourier transformed (FFT) patterns of each HRTEM image. Yellow circles in d indicate regions where the lattice fringes were not observed.

Figure 8. High-resolution STEM images showing formation of silver particles after sodiation. HAADF STEM images of (a) pristine (Ag1.22Mn8O16) and (b) 2 equiv. of Na-inserted nanorods. (a) Pristine nanorods have clean surface without precipitates. (b) Tiny particles with ∼2 nm size spread out all over nanorods after sodiation. (c) Atomic-resolution HAADF STEM image of a 2 equiv Na-inserted nanowire. Several atoms forming apparent periodicity have brighter intensity than surrounding regions, which implies that these atoms should be mainly composed of the silver element. The inset shows the FFT pattern of yellow boxed area, which matches well with [001] zone axis of silver metal. Atomic model structure of [001] silver metal is described in the yellow boxed area. G

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ACKNOWLEDGMENTS The authors acknowledge the Center for Mesoscale Transport Properties, an Energy Frontier Research Center supported by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, under award #DE-SC0012673 for financial support. The XPS experiments were carried out at the Center for Functional Nanomaterials at Brookhaven National Laboratory, which are supported by the Department of Energy, Office of Basic Energy Sciences (DE-AC02-98CH10886). TEM work was supported by the U.S. Department of Energy, Office of Basic Energy Science, Division of Materials Science and Engineering, under Contract DE-SC0012704. The X-ray absorption spectroscopy measurements were performed at Beamline 12BM-B of the Advanced Photon Source, a U.S. Department of Energy (DOE) Office of Science User Facility operated for the DOE Office of Science by Argonne National Laboratory under Contract DE-AC02-06CH11357. The authors thank Christopher J. Pelliccione for helpful discussions regarding XAS.

Figure 9. Mn K-edge XANES of L-Ag-OMS-2 before and after sodiation/lithiation.

energy further decreased to 6556.9 eV in Na4Ag1.22Mn8O16, and 6556.0 eV in Li4Ag1.22Mn8O16. The slightly higher edge energy of Na4Ag1.22Mn8O16 is consistent with partial reduction of the Ag+ sites in the discharged Na battery as also indicated by TEM and XRD.

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4. CONCLUSION Pure Ag-OMS-2 (AgxMn8O16, x = 1.22, L-Ag-OMS-2 or 1.66, H-Ag-OMS-2) materials were synthesized by a low temperature reflux-based method. L-Ag-OMS-2 exhibited good cycle stability with a reversible capacity of ∼2 electron equivalents (60 mAh/g) for 50 cycles in a Na battery, with similar cycling performance in the Li battery. However, the sample showed capacity fade when cycled in a large voltage window (3.8−1.3 V) in the Na battery, suggestive of structural instability under high sodiation levels. With cycling, when the current density increased from 10 mA/g to 200 mA/g, L-Ag-OMS-2 showed good rate capability with 50% capacity retention in the Na battery while showing only 29% capacity retention at high rate in the Li battery. EIS measurements indicated a significant drop in the charge transfer resistance as a result of the sodiation process whereas with lithiation, a gradual resistance increase was observed. The formation of silver metal particles was confirmed by XRD and TEM of sodiated samples consistent with the observed decrease in resistance. This study demonstrates that tunnel structured Ag-OMS-2 can host reversible Na ion insertion/deinsertion, where the in situ formation of silver metal through reduction-displacement of Ag+ to Ag0 during the electrochemical reduction process was concomitant with a significant decrease of the cell impedance.

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REFERENCES

(1) Goodenough, J. B. Evolution of Strategies for Modern Rechargeable Batteries. Acc. Chem. Res. 2013, 46, 1053−1061. (2) Goodenough, J. B.; Park, K.-S. The Li-Ion Rechargeable Battery: A Perspective. J. Am. Chem. Soc. 2013, 135, 1167−1176. (3) Armand, M.; Tarascon, J. M. Building Better Batteries. Nature 2008, 451, 652−657. (4) Liu, J.; Zhang, J.-G.; Yang, Z.; Lemmon, J. P.; Imhoff, C.; Graff, G. L.; Li, L.; Hu, J.; Wang, C.; Xiao, J.; Xia, G.; Viswanathan, V. V.; Baskaran, S.; Sprenkle, V.; Li, X.; Shao, Y.; Schwenzer, B. Materials Science and Materials Chemistry for Large Scale Electrochemical Energy Storage: From Transportation to Electrical Grid. Adv. Funct. Mater. 2013, 23, 929−946. (5) Thackeray, M. M.; Wolverton, C.; Isaacs, E. D. Electrical Energy Storage for Transportation-Approaching the Limits of, and Going Beyond, Lithium-Ion Batteries. Energy Environ. Sci. 2012, 5, 7854− 7863. (6) Kundu, D.; Talaie, E.; Duffort, V.; Nazar, L. F. The Emerging Chemistry of Sodium Ion Batteries for Electrochemical Energy Storage. Angew. Chem., Int. Ed. 2015, 54, 3431−3448. (7) Pan, H.; Hu, Y.-S.; Chen, L. Room-Temperature Stationary Sodium-Ion Batteries for Large-Scale Electric Energy Storage. Energy Environ. Sci. 2013, 6, 2338−2360. (8) Huie, M. M.; Bock, D. C.; Takeuchi, E. S.; Marschilok, A. C.; Takeuchi, K. J. Cathode Materials for Magnesium and Magnesium-Ion Based Batteries. Coord. Chem. Rev. 2015, 287, 15−27. (9) Manthiram, A.; Fu, Y.; Su, Y.-S. Challenges and Prospects of Lithium−Sulfur Batteries. Acc. Chem. Res. 2013, 46, 1125−1134. (10) Bruce, P. G.; Freunberger, S. A.; Hardwick, L. J.; Tarascon, J.-M. Li-O2 and Li-S Batteries with High Energy Storage. Nat. Mater. 2012, 11, 19−29. (11) Li, F.; Zhang, T.; Zhou, H. Challenges of Non-Aqueous Li-O2 Batteries: Electrolytes, Catalysts, and Anodes. Energy Environ. Sci. 2013, 6, 1125−1141. (12) Shannon, R. Revised Effective Ionic Radii and Systematic Studies of Interatomic Distances in Halides and Chalcogenides. Acta Crystallogr., Sect. A: Cryst. Phys., Diffr., Theor. Gen. Crystallogr. 1976, 32, 751−767. (13) Mendiboure, A.; Delmas, C.; Hagenmuller, P. Electrochemical Intercalation and Deintercalation of NaxMnO2 Bronzes. J. Solid State Chem. 1985, 57, 323−331. (14) Caballero, A.; Hernan, L.; Morales, J.; Sanchez, L.; Santos Pena, J.; Aranda, M. A. G. Synthesis and Characterization of Hightemperature Hexagonal P2-Na0.6MnO2 and Its Electrochemical Behaviour as Cathode in Sodium Cells. J. Mater. Chem. 2002, 12, 1142−1147.

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Conductive Metallic Silver Nanoparticles. Chem. Mater. 2009, 21, 4934−4939. (35) Takeuchi, E. S.; Marschilok, A. C.; Takeuchi, K. J.; Ignatov, A.; Zhong, Z.; Croft, M. Energy Dispersive X-ray Diffraction of LithiumSilver Vanadium Phosphorous Oxide Cells: In Situ Cathode Depth Profiling of an Electrochemical Reduction-Displacement Reaction. Energy Environ. Sci. 2013, 6, 1465−1470. (36) Marschilok, A. C.; Kozarsky, E. S.; Tanzil, K.; Zhu, S.; Takeuchi, K. J.; Takeuchi, E. S. Electrochemical Reduction of Silver Vanadium Phosphorous Oxide, Ag2VO2PO4: Silver Metal Deposition and Associated Increase in Electrical Conductivity. J. Power Sources 2010, 195, 6839−6846. (37) Huang, J.; Marschilok, A. C.; Takeuchi, E. S.; Takeuchi, K. J. Microwave-Assisted Synthesis of Silver Vanadium Phosphorus Oxide, Ag2VO2PO4: Crystallite Size Control and Impact on Electrochemistry. Chem. Mater. 2016, 28, 2191−2199. (38) Kirshenbaum, K. C.; Menard, M. C.; Kim, Y. J.; Marschilok, A. C.; Takeuchi, K. J.; Takeuchi, E. S. Electrochemical Reduction of Ag0.48VOPO4: A Mechanistic Study Employing X-Ray Absorption Spectroscopy and X-Ray Powder Diffraction. J. Electrochem. Soc. 2015, 162, A1537−A1543. (39) Kirshenbaum, K.; Bock, D. C.; Lee, C.-Y.; Zhong, Z.; Takeuchi, K. J.; Marschilok, A. C.; Takeuchi, E. S. In Situ Visualization of Li/ Ag2VP2O8 Batteries Revealing Rate-Dependent Discharge Mechanism. Science 2015, 347, 149−154. (40) Kirshenbaum, K. C.; Bock, D. C.; Zhong, Z.; Marschilok, A. C.; Takeuchi, K. J.; Takeuchi, E. S. Electrochemical Reduction of Ag2VP2O8 Composite Electrodes Visualized via In Situ Energy Dispersive X-ray Diffraction (EDXRD): Unexpected Conductive Additive Effects. J. Mater. Chem. A 2015, 3, 18027−18035. (41) Chen, C.; Liang, B.; Ogino, A.; Wang, X.; Nagatsu, M. Oxygen Functionalization of Multiwall Carbon Nanotubes by MicrowaveExcited Surface-Wave Plasma Treatment. J. Phys. Chem. C 2009, 113, 7659−7665. (42) Liu, L.; Ma, D.; Zheng, H.; Li, X.; Cheng, M.; Bao, X. Synthesis and Characterization of Microporous Carbon Nitride. Microporous Mesoporous Mater. 2008, 110, 216−222. (43) Scherrer, P. Estimation of the Size and Internal Structure of Colloidal Particles by Means of Röntgen. Nachr. Ges. Wiss. Göttingen 1918, 2, 96−100. (44) Shao-Horn, Y.; Hackney, S. A.; Johnson, C. S.; Thackeray, M. M. Microstructural Features of α - MnO2 Electrodes for Lithium Batteries. J. Electrochem. Soc. 1998, 145, 582−589. (45) Nicolas-Tolentino, E.; Tian, Z.-R.; Zhou, H.; Xia, G.; Suib, S. L. Effects of Cu2+ Ions on the Structure and Reactivity of Todorokiteand Cryptomelane-Type Manganese Oxide Octahedral Molecular Sieves. Chem. Mater. 1999, 11, 1733−1741. (46) Huang, H.; Sithambaram, S.; Chen, C.-H.; King’ondu Kithongo, C.; Xu, L.; Iyer, A.; Garces, H. F.; Suib, S. L. Microwave-Assisted Hydrothermal Synthesis of Cryptomelane-Type Octahedral Molecular Sieves (OMS-2) and Their Catalytic Studies. Chem. Mater. 2010, 22, 3664−3669. (47) Ching, S.; Suib, S. L. Synthetic Routes to Microporous Manganese Oxides. Comments Inorg. Chem. 1997, 19, 263−282. (48) Xu, Y.; Zhou, M.; Wang, X.; Wang, C.; Liang, L.; Grote, F.; Wu, M.; Mi, Y.; Lei, Y. Enhancement of Sodium Ion Battery Performance Enabled by Oxygen Vacancies. Angew. Chem., Int. Ed. 2015, 54, 8768− 8771. (49) Yan, W.; Ayvazian, T.; Kim, J.; Liu, Y.; Donavan, K. C.; Xing, W.; Yang, Y.; Hemminger, J. C.; Penner, R. M. Mesoporous Manganese Oxide Nanowires for High-Capacity, High-Rate, Hybrid Electrical Energy Storage. ACS Nano 2011, 5, 8275−8287. (50) Yu, Z.; Duong, B.; Abbitt, D.; Thomas, J. Highly Ordered MnO2 Nanopillars for Enhanced Supercapacitor Performance. Adv. Mater. 2013, 25, 3302−3306. (51) Suib, S. L. Porous Manganese Oxide Octahedral Molecular Sieves and Octahedral Layered Materials. Acc. Chem. Res. 2008, 41, 479−487.

(15) Zhao, J.; Zhao, L.; Dimov, N.; Okada, S.; Nishida, T. Electrochemical and Thermal Properties of α-NaFeO2 Cathode for Na-Ion Batteries. J. Electrochem. Soc. 2013, 160, A3077−A3081. (16) Berthelot, R.; Carlier, D.; Delmas, C. Electrochemical Investigation of the P2−NaxCoO2 Phase Diagram. Nat. Mater. 2011, 10, 74−80. (17) Yu, C.-Y.; Park, J.-S.; Jung, H.-G.; Chung, K.-Y.; Aurbach, D.; Sun, Y.-K.; Myung, S.-T. NaCrO2 Cathode for High-Rate Sodium-Ion Batteries. Energy Environ. Sci. 2015, 8, 2019−2026. (18) Guignard, M.; Didier, C.; Darriet, J.; Bordet, P.; Elkaïm, E.; Delmas, C. P2-NaxVO2 System as Electrodes for Batteries and Electron-Correlated Materials. Nat. Mater. 2013, 12, 74−80. (19) Yabuuchi, N.; Kajiyama, M.; Iwatate, J.; Nishikawa, H.; Hitomi, S.; Okuyama, R.; Usui, R.; Yamada, Y.; Komaba, S. P2-type Nax[Fe1/ 2Mn1/2]O2Made from Earth-Abundant Elements for Rechargeable Na Batteries. Nat. Mater. 2012, 11, 512−517. (20) Sauvage, F.; Laffont, L.; Tarascon, J. M.; Baudrin, E. Study of the Insertion/Deinsertion Mechanism of Sodium into Na0.44MnO2. Inorg. Chem. 2007, 46, 3289−3294. (21) Kim, H.; Kim, D. J.; Seo, D.-H.; Yeom, M. S.; Kang, K.; Kim, D. K.; Jung, Y. Ab Initio Study of the Sodium Intercalation and Intermediate Phases in Na0.44MnO2 for Sodium-Ion Battery. Chem. Mater. 2012, 24, 1205−1211. (22) Zhan, P.; Wang, S.; Yuan, Y.; Jiao, K.; Jiao, S. Facile Synthesis of Nanorod-like Single Crystalline Na0.44MnO2 for High Performance Sodium-Ion Batteries. J. Electrochem. Soc. 2015, 162, A1028−A1032. (23) Su, D.; Ahn, H.-J.; Wang, G. Hydrothermal Synthesis of αMnO2 and β-MnO2 Nanorods as High Capacity Cathode Materials for Sodium Ion Batteries. J. Mater. Chem. A 2013, 1, 4845−4850. (24) Su, D.; Ahn, H.-J.; Wang, G. β-MnO2 Nanorods with Exposed Tunnel Structures as High-Performance Cathode Materials for Sodium-Ion Batteries. NPG Asia Mater. 2013, 5, e70. (25) Tompsett, D. A.; Islam, M. S. Electrochemistry of Hollandite αMnO2: Li-Ion and Na-Ion Insertion and Li2O Incorporation. Chem. Mater. 2013, 25, 2515−2526. (26) Zhu, S.; Marschilok, A. C.; Lee, C.-Y.; Takeuchi, E. S.; Takeuchi, K. J. Synthesis and Electrochemistry of Silver Hollandite. Electrochem. Solid-State Lett. 2010, 13, A98−A100. (27) Takeuchi, K. J.; Yau, S. Z.; Menard, M. C.; Marschilok, A. C.; Takeuchi, E. S. Synthetic Control of Composition and Crystallite Size of Silver Hollandite, AgxMn8O16: Impact on Electrochemistry. ACS Appl. Mater. Interfaces 2012, 4, 5547−5554. (28) Takeuchi, K. J.; Yau, S. Z.; Subramanian, A.; Marschilok, A. C.; Takeuchi, E. S. The Electrochemistry of Silver Hollandite Nanorods, AgxMn8O16: Enhancement of Electrochemical Battery Performance via Dimensional and Compositional Control. J. Electrochem. Soc. 2013, 160, A3090−A3094. (29) Wu, L.; Xu, F.; Zhu, Y.; Brady, A. B.; Huang, J.; Durham, J. L.; Dooryhee, E.; Marschilok, A. C.; Takeuchi, E. S.; Takeuchi, K. J. Structural Defects of Silver Hollandite, AgxMn8Oy, Nanorods: Dramatic Impact on Electrochemistry. ACS Nano 2015, 9, 8430− 8439. (30) Yang, Z.; Trahey, L.; Ren, Y.; Chan, M. K. Y.; Lin, C.; Okasinski, J.; Thackeray, M. M. In Situ High-Energy Synchrotron X-ray Diffraction Studies and First Principles Modeling of α-MnO2 Electrodes in Li-O2 and Li-Ion Coin Cells. J. Mater. Chem. A 2015, 3, 7389−7398. (31) Vicat, J.; Fanchon, E.; Strobel, P.; Tran Qui, D. The Structure of K1.33Mn8O16 and Cation Ordering in Hollandite-Type Structures. Acta Crystallogr., Sect. B: Struct. Sci. 1986, 42, 162−167. (32) Takeuchi, E. S.; Thiebolt, W. C. The Reduction of Silver Vanadium Oxide in Lithium/Silver Vanadium Oxide Cells. J. Electrochem. Soc. 1988, 135, 2691−2694. (33) Leising, R. A.; Thiebolt, W. C.; Takeuchi, E. S. Solid-State Characterization of Reduced Silver Vanadium Oxide from the Li/SVO Discharge Reaction. Inorg. Chem. 1994, 33, 5733−5740. (34) Takeuchi, E. S.; Marschilok, A. C.; Tanzil, K.; Kozarsky, E. S.; Zhu, S.; Takeuchi, K. J. Electrochemical Reduction of Silver Vanadium Phosphorus Oxide, Ag2VO2PO4: The Formation of Electrically I

DOI: 10.1021/acsami.6b08549 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Forum Article

ACS Applied Materials & Interfaces (52) Suib, S. L. Structure, Porosity, and Redox in Porous Manganese Oxide Octahedral Layer and Molecular Sieve Materials. J. Mater. Chem. 2008, 18, 1623−1631. (53) Yuan, Y.; Nie, A.; Odegard, G. M.; Xu, R.; Zhou, D.; Santhanagopalan, S.; He, K.; Asayesh-Ardakani, H.; Meng, D. D.; Klie, R. F.; Johnson, C.; Lu, J.; Shahbazian-Yassar, R. Asynchronous Crystal Cell Expansion during Lithiation of K+-Stabilized α-MnO2. Nano Lett. 2015, 15, 2998−3007.

J

DOI: 10.1021/acsami.6b08549 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX