In Situ Electrical Conductivity of Li x MnO 2 ... - ACS Publications

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In Situ Electrical Conductivity of LixMnO2 Nanowires as a Function of x and Size Mya Le,† Yu Liu,‡ Hui Wang,‡,¶ Rajen K. Dutta,‡ Wenbo Yan,† John C. Hemminger,† Ruqian Q. Wu,‡ and Reginald M. Penner*,†,§ †

Department of Chemistry, ‡Department of Physics and Astronomy, and §Department of Chemical Engineering and Materials Science, University of California, Irvine, California 92697, United States ¶ Department of Physics, Fudan University Shanghai, 200433 China S Supporting Information *

ABSTRACT: Manganese oxide, MnO2, excels as a hybrid electrical energy storage material: The manganese centers in MnO2 are capable of undergoing a reduction from 4+ to 3+ balanced by the intercalation of lithium ions to form LixMnO2 while its conductive surfaces simultaneously store energy as an electrical double layer capacitor. The highest capacitance and power performance for MnO2 has been obtained for ensembles of nanowires that are 200 nm or less in width and many microns in length. Typically such MnO2 nanowires are attached to a current collector at just one end, and electrical conductivity of the nanowire is therefore required in order to maintain a consistent redox and charge state along its axis. The electrical conductance of the nanowire therefore plays a very important role, and yet this parameter has been measured in few previous studies. In this work, we directly measure the electrical conductance of δ-MnO2 nanowires in situ in 1 M LiClO4, acetonitrile as a function of the equilibrium Li content for nanowires with varying lateral dimensions. This measurement is accomplished using arrays of 200 MnO2 nanowires that are 40−60 nm in height and 275−870 nm in width and which span a 10 μm gap between two gold contacts. Nanowires of fully oxidized MnO2 are first prepared at +0.60 V vs MSE in acetonitrile. As the equilibrium electrode potential is decreased from 0.60 V to −0.80 V and lithium is intercalated, the electrical conductivity of MnO2 nanowires increases by up to 1 order of magnitude. The measured change in conductivity is dependent on the equilibrium potential, which in turn is related to the Li content, and also depends on the width of nanowires. After doping at −0.80 V vs MSE, the conductivity increases by 30% for a 870 nm wide nanowire array and 880% for a 275 nm wide nanowire array. TEM investigations implicate the nanowire porosity in this difference.

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INTRODUCTION Manganese oxide, MnO2, excels as a hybrid electrical energy storage material:1,2 The manganese centers in MnO2 are capable of undergoing a reduction from 4+ to 3+, producing LixMnO2, balanced by the intercalation of lithium ions. In detail, this reaction can be summarized as follows:

influence of rate-limiting Li transport within the active material by reducing the maximum diffusion path length for Li. In this paper we address two questions that strongly influence the efficiency with which a LixMnO2 nanowire is able to function as an energy storage element: First, what is the dc electrical conductance, G, of a LixMnO2 nanowire and the dc conductivity, σ, of the LixMnO2 within it as a function of x? Second, what influence, if any, do the lateral dimensions of a nanowire have on G and σ? To answer these two questions, we have used the LPNE method20−23 to prepare arrays of ≈200 MnO2 nanowires where each nanowire is 40−60 nm in height, 275−870 nm in width, and 10 μm in length, spanning two gold electrical contacts. Using these nanowire arrays, we have directly measured the dc electrical conductivity in situ in 1.0 M LiClO4, acetonitrile. To be clear, in situ refers to the

Lix (Mn 3x +Mn14−+x) O2 + δ Li+ + δ e− → Lix + δ (Mn 3x ++ δ Mn14−+(x + δ)) O2

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The conductive surfaces of the MnO2 can simultaneously store energy without intercalating Li by acting like an electrical double layer capacitor.3−10 The capacitance and the lithiation/ delithiation kinetics of MnO2 are both enhanced by reducing the critical dimension of the MnO2 either by reducing the film thickness, in the case of films, or by employing a nanowire morphology where the critical dimension is the nanowire radius.cf.11−19 Reducing the critical dimension reduces the © 2015 American Chemical Society

Received: March 10, 2015 Revised: April 13, 2015 Published: April 16, 2015 3494

DOI: 10.1021/acs.chemmater.5b00912 Chem. Mater. 2015, 27, 3494−3504

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Chemistry of Materials measurement of σ for an array of LixMnO2 nanowires as a function of their equilibrium potential, Eeq, while the nanowire array is momentarily disconnected from the potentiostat. In these measurements, we confine our attention to the δ phase of MnO2 that, upon lithiation, is transformed into the Birnessite phase of LixMnO2.3 This study is the first to our knowledge to measure the xdependence of σ for LixMnO2 nanowires, but these data have been reported for LixMnO2 films previously (vide infra).24−27 The first difference is that the maximum electrical conductivity of LixMnO2 nanowires is higher, approaching 5 × 10−3 S/cm at −0.80 V vs MSE. This is a factor of 25 higher than the in situ conductivity reported by Uchida et al. for spinel films of LixMnO2.24,25 Second, the conductivity change for nanowires as a function of x is greater in absolute terms than seen previously for films. Finally, we observe a strong dependence of σ on the nanowire width at constant nanowire height: Increasing the nanowire width reduces σ for all values of x while also dramatically reducing the degree to which σ is modulated by x. This suggests that the properties of the MnO2 evolve even as it is electrodeposited to form a nanowire. A hypothesis is advanced to account for this surprising observation.

Figure 1. (a) Process flow for fabrication by LPNE of δ-MnO2 nanowire array with gold contacts for in situ electrical resistance measurements. (b) Scanning electron microscopy (SEM) image of δMnO2 nanowire array with gold contacts (false color). These arrays contained ≈200 nanowires.

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EXPERIMENTAL SECTION Chemicals and Materials. Gold pellets (5 N purity, Kurt J. Lesker Co.) were used for the preparation by thermal evaporation of ultrathin gold layers. Manganese perchlorate hydrate (Mn(ClO4)2·H2O, 99%), iodine (I2, 99.8%), and lithium perchlorate (LiClO4, 99.99%) were used as received from Sigma-Aldrich. Potassium iodide (KI, 99%) and acetone were used as received from Fisher (ACS Certified). Electrochemical characterization experiments were conducted in 1.0 M LiClO4 (dry, 99.99%) in dry acetonitrile. Ultrapure, dry acetonitrile was obtained by filtering reagent grade acetonitrile using a Jorg Meyer Phoenix SDS column purification system. Positive photoresist Shipley S-1808 and developer MF-319 were used as received from Microchem Corporation. Synthesis of δ-MnO2 Nanowires. The fabrication of δMnO2 nanowire arrays for in situ resistance measurements was accomplished using the LPNE process (Figure 1a).12 Starting with 1″ × 1″ squares of cleaned glass, a 60 nm layer of gold is first thermally evaporated. Gold covered slides are then spincoated with positive photoresist (PR, Shipley, S1808) and softbaked (90 °C for 30 min). Then, the PR layer is photopatterned using a chromium/quartz contact mask in conjunction with 365 nm UV light source combined with a shutter and alignment stage (Newport, 83210i-line). Next, the UV exposed photoresist pattern was developed for 18s in developer solution (Shipley, MF-319). This process, which defines the positions of the 200 nanowire to be deposited onto the glass surface, produced the starting point depicted in Figure 1a at upper left. Then the gold was removed from this patterned surface by etching with aqueous KI/I2 solution (4/2g in 80 mL of H2O) for 21 s (step 1). This step removed exposed gold and also formed an undercut at the edge of the PR forming a template for the subsequent electrodeposition of the δ-MnO2 (Figure 1a, top center). Except for one edge where an electrical contact is connected to the gold layer, the entire patterned region was immersed in an aqueous 2 mM Mn(ClO4)2, 50 mM LiClO4 plating solution for deposition. Within this trench, MnO2 nanowires were synthesized using multipulse chronoamperometry in which +0.6 V versus MSE (saturated mercurous sulfate reference electrode) growth pulses with a

duration of 0.5 s were applied at 20 s intervals. By controlling the total deposition time, the nanowire width was adjusted. After MnO2 electrodeposition was completed, the PR layer was removed using acetone, and the exposed gold was etched for 1 min, producing an array of freestanding MnO2 nanowires supported on the glass surface (step 2). Electrical contacts were prepared using a second photolithography operation in which the MnO2 nanowire arrays were then again spin-coated with PR and patterned (step 3). A 40−60 nm layer of gold was then evaporated onto this patterned surface (step 4). Finally, metal lift-off removed excess gold, leaving the two contact pads. The complete device therefore consisted of a parallel array of 200 MnO2 bracketed by two gold contacts across a distance of 10 μm (step 5). Electrochemical Characterization. Electrodeposition was conducted using a three-electrode electrochemical cell with a Princeton Applied Research 2263 potentiostat with Pt foil as counter electrode and Hg/Hg2SO4, K2SO4 (saturated) (MSE) as reference electrode. All electrochemical characterization was conducted in 1.0 M LiClO4 (dry, 99.99%) in dry acetonitrile. Prior to each measurement, the acetonitrile was presaturated with N2 gas. Structural Characterization. Scanning electron micrographs were acquired using either a Philips XL-30 (field emission gun scanning electron microscope) or a FEI Mallegan XHR SEM (extreme high-resolution scanning electron microscope), both instruments operating at an accelerating voltage of 10 kV. All samples were sputtered with a thin layer (2 nm) of Ir before imaging to prevent charging. Atomic force microscopy (AFM) images and amplitude traces were acquired using an Asylum Research, MFP-3D AFM equipped with Olympus AC160TS tips in a laboratory air ambient. Transmission electron microscopy (TEM) images and selective area electron diffraction (SAED) patterns were obtained using a Philips CM 20 TEM operating at 200 keV. TEM samples were prepared by LPNE-deposition of MnO2 nanowires onto a Si grid with 20 nm thick Si3N4 windows (Ted Pella, Inc.). Grazing-incidence 3495


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Figure 2. (a) Flowchart for the acquisition of in situ resistance versus equilibrium electrochemical potential data. (b) Schematic diagram of the electrochemical cell for equilibrating the nanowire array at various potentials. (c) Schematic diagram showing configuration for in situ resistance measurement. (d) Current versus time transients acquired for an array of 685 nm (w) × 47 nm (h) δ-MnO2 nanowires. The initial potential in each case was +0.60 V vs MSE, and the equilibrium potential, Eeq, is as indicated. (e) Current versus potential plots for this δ-MnO2 nanowire array after equilibration at five different Eeq values, as indicated. The “as-synthesized” plot corresponds to the Eeq = +0.60 vs MSE.

Figure 3. Scanning electron micrographs (SEM) of δ-MnO2 nanowires: (a) low-magnification image of a nanowire array, (b−e). Higher magnification images of individual δ-MnO2 nanowires prepared using 20, 25, 30, and 35 voltages pulses, respectively, as demonstrated in Figure 3. (f) Nanowire width versus number of voltage pulses. Error bars represent ±1σ.

reference28 and further calibrated with Au (4f7/2) at 84.0 eV from gold foil physically attached to the sample surface. Deconvolution and spectral line fitting were carried out using XPS Peak 4.1. Density Function Theory. Density functional simulations were performed with the Vienna ab initio simulation package (VASP).29 The spin-polarized generalized gradient approximation (GGA)30 was used for the description of the exchangecorrelation interaction among electrons. We treated Mn-3d4s, O-2s2p, H-1s, and Li-2s as valence states and adopted the projector-augmented wave (PAW) pseudopotentials to represent the effect of their ionic cores.31 To properly describe the strong correlation among electrons in the partially filled d-shell, we used the GGA+U method introduced by Liechtenstein et al.,32 with U = 6.7 eV and J = 1.2 eV for Mn atoms.33 Furthermore, we included the van der Waals correction in the self-consistent loop, using the nonlocal optB86b-vdW func-

X-ray diffraction (GIXRD) and in-plane X-ray diffraction were both obtained using a Rigaku Smartlab X-ray diffractometer employing parallel beam optics with a fixed incident angle of 0.9°. The X-ray generator was operated at 40 kV and 44 mA with Cu Kα irradiation. The PDXL (Materials Data, Inc.) X-ray pattern data processing software was used to analyze acquired patterns. XPS (X-ray Photoelectron Spectroscopy). XPS measurements were performed with an ESCALAB MKII (VG Scientific) surface analysis instrument. The ultrahigh vacuum multichamber system is equipped with a twin anode X-ray source (Mg/Al) and a 150 mm hemispherical electron energy analyzer. Spectra presented here were collected using Al Kα Xrays (1486.6 eV) in constant energy mode with pass energy of 20 eV. During acquisition the base pressure of the spectroscopy chamber was 5 × 10−10 Torr. Binding energies were calibrated using the C(1s) peak of adventitious carbon set at 284.8 eV as a 3496


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Chemistry of Materials tional that was implemented in VASP34,35 for the structural optimization. The energy cutoff for the plane-wave expansion was 700 eV, sufficient for transition-metal oxide systems according to our test calculations. Ab initio molecular dynamics (AIMD) simulations were performed to elucidate the formation and evolvement of MnO2 + H2O + Li bonds at finite temperature. At the end of AIMD steps, we further optimize the atomic structures for MnO2 with the presence of H2O and Li, with a criterion that the atomic force on each atom becomes weaker than 0.01 eV/Å, and the energy convergence is better than 10−5 eV. We sampled the Brillouin zone with a single gamma point in the AIMD simulations and with a 5 × 5 × 9 Monkhorst−Pack36 k-mesh for other calculations.

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RESULTS AND DISCUSSION Synthesis and Characterization of δ-MnO2 Nanowires. The electrochemical synthesis of manganese oxide, MnO2, involved the electrochemical oxidation of Mn2+:7,12,22 Mn 2 + + 2H 2O + x Li+ → LixMn 3x +Mn14−+xO2 + 4H+ + (2 − x)e−

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Reaction 2 was carried out in aq. 2 mM Mn(ClO4)2, 50 mM LiClO4 at lithographically patterned gold nanoband electrodes that were ≈47 nm in height, prepared by LPNE as described above (Figure 1a). The value of x, estimated from XPS analysis (vide infra), is 0.1. From an initial potential of +0.30 V, LixMnO2 was electrodeposited by applying multiple potential pulses of +0.60 V vs MSE for 0.50 s. Consecutive +0.60 V × 0.50 s growth pulses were separated by 20 s rest periods at +0.30 V. As previously described,12 this pulsed deposition growth procedure provided the means for depositing LixMnO2 nanowire with a compact morphology, avoiding dendritic growth. Arrays of 200 linear MnO2 nanowires (Figure 3a), patterned at an interwire pitch of ≈4 μm on glass, were employed for in situ electrical conductivity measurements. The nanowire width was directly proportional to the number of growth pulses (Figure 3b−e), and 20−35 pulses were applied to produce MnO2 nanowire ranging in width from 275 to 870 nm (Figure 3f). The characterization of MnO2 nanowires using SEM, TEM, AFM, and XRD (shown in Figure S1 of the Supporting Information) was carried out on arrays of MnO2 nanowires on glass like that shown in Figure 3a, before electrical contacts were deposited for electrical measurements (Figure 1, steps 3−5). A transmission electron microscope (TEM) image (Figure 4) shows a single, as-prepared MnO2

Figure 5. Mn 3s XPS data. (a) Spectra of the Mn 3s region for LixMnO2 films (thickness ≈120 nm) on Pt as a function of Eeq in 1.0 M LiClO4 in dry acetonitrile. These data were acquired ex situ for freshly deposited and equilibrated films. (b) Plots of νMn versus Eeq (red trace) and of ΔEMn 3s versus Eeq (blue trace). The quantities ΔEMn 3s and νMn are linearly related through eq 3.

nanowire before the removal of the gold edge. At high magnification (Figure 4, right), a wrinkled morphology is seen for the MnO2, and considerable porosity is evident as well. We discuss the morphology of these nanowires in greater detail below. XPS measurements (Figure 5), on the other hand, were conducted on MnO2 films (thickness ≈120 nm) on platinum because the signal-to-noise ratio afforded by nanowire arrays was insufficient. XPS allowed the mean manganese oxidation state, νMn, to be determined at every equilibrium potential, Eeq. This is because, as previously demonstrated,37,38 νMn is linearly correlated with the exchange splitting of the Mn 3s photoelectron peak, ΔEMn3s:38 νMn = 12.5 − 1.82 ΔEMn3s (eV)

Figure 4. TEM image of an MnO2 nanowire. (a) Low-magnification TEM image of a MnO2 nanowire after deposition using 35 pulses of +0.60 V. This image shows the nanowire at step 1, Figure 1a after the removal of the photoresist layer, but not the gold electrode which is shown at upper right. (b) Higher magnification TEM image of this nanowire showing the wrinkled texture of the MnO2 sheets present within the nanowire.

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As seen in Figure 5a, ΔEMn 3s has a minimum value of 4.73 eV at Eeq = +0.60 V, corresponding to νMn = 3.9, close to the value of 4.0 expected for the fully delithiated state of LixMnO2 with x = 0. ΔEMn3s increases as Eeq is made more negative, reaching 5.23 eV at −0.80 V, corresponding to νMn = 3.0 (Figure 5a), as expected for the fully lithiated state, x = 1.0. A 3497


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Figure 6. Cyclic voltammetery and specific capacitance, Csp. (a) Cyclic voltammograms (CVs) at 5 mV/s for an array 685 nm × 47 nm LixMnO2 showing three voltage windows: −0.9 V to +0.60 V (1.6 V, red trace), −0.8 V to +0.60 V (1.4 V, blue trace), and −0.8 V to +0.40 V (black trace). Reversible capacitance is maximized in the 1.4 V voltage window. (b) CVs for five scan rates, as indicated, for the same nanowire array studied in (a). (c) Plots of Csp versus scan rate for four nanowires investigated below.

Figure 7. In situ electrical conductivity, σ, of LixMnO2 nanowires with thickness ≈47 nm and width as indicated: (a) σ versus equilibrium electrochemical potential, Eeq. (b) σ versus lithium content, x, calculated from the doping charge according to reaction 1. All data were acquired in random order of Eeq (or x) in 1.0 M LiClO4 in dry acetonitrile. Also shown are conductivity data reported previously for LixMnO2 films by Chen et al.,26 Uchida et al.,24,25 and Kanoh et al.,27 who reported ionic, not electronic, conductivity values.

plot of ΔEMn3s and νMn versus Eeq (Figure 5b) demonstrates that the 1.4 V range from Eeq = +0.60 V to −0.80 V vs MSE provides access to the full range of lithiation states of LixMnO2 in acetonitrile. Within the voltage range from +0.60 V to −0.80 V vs MSE in 1.0 M LiClO4, acetonitrile, an array of 10 μm length nanowires exhibit state-of-the-art energy storage properties. Extending the voltage range to more negative potentials of −0.90 V causes an irreversible reduction seen in the cyclic voltammograms (CVs) of Figure 6a. This reduction has been assigned to the reduction of Mn3+ centers to Mn2+a destabilizing reaction that is associated with the dissolution of the MnO2. Repeated excursions of the potential to +0.60 V (Figure 6b), on the other hand, do not destabilize MnO2 nanowires. From scan rate-dependent CVs (Figure 6b), the specific capacitance, Csp, of an array of nanowires can be calculated using eq 4 Csp =

Q ΔE × mMnO2

in which Q is the total charge associated with scanning between −0.80 V and +0.60 V in either direction, ΔE is 1.4 V, and mMnO2 is the dry mass of the MnO2, calculated from the deposition charge, Qdep, as previously described (mMnO2 = 0.68 μg/mC).12 Plots of Csp versus potential scan rate (Figure 6c) show decreasing Csp with increasing scan rate from 1 to 100 mV/s, as previously reported for MnO2.11,12 This behavior is attributed to the influence of rate-limiting Li+ insertion and the solid-state diffusion of Li on charge storage. Also apparent (Figure 6c) is the nanowire size dependence of Csp which increases as the nanowire width is reduced from 872 to 277 nm at a constant 47 nm height, also as previously described. We discuss this observation in greater detail below. The limiting value for Csp is ≈1100 F/g (at 1 mV/s), which approximates the highest Csp values measured in our prior work.11,12,39 In spite of the similarities between the results seen in Figure 6b,c and our prior work,11,12,39 one must recognize an important difference in the system being studied: Here, gold current collectors are located at the ends of an array of 10 μm MnO2 nanowires. This means that MnO2 nanowires must

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Figure 8. TEM analysis of LixMnO2 nanowire density. (a) Four TEM images (top) of a 780 (w) × 47 nm (h) nanowire showing relationship to the gold nanoelectrode. Thresholded images (bottom) are used to derive an electron dense fraction that can be compared for nanowire of various width. (b) Plot of electron dense fraction as a function of distance from the gold edge. Wide nanowires are less porous than narrower nanowires, and for a given nanowire, porosity increases with distance from the gold edge.

conduct electrons up to 5 μm in order for charge storage and delivery to occur. In contrast, in our previous work this was not the case: We have either measured Csp without removing the gold edge in the LPNE process,12 enforcing intimate contact between the gold current collector and the nanowire edge along its entire length, or we have prepared and studied core/shell nanowires11,39 in which a gold nanowire core is surrounded by a MnO2 shell where, again, proximity (within 100 nm) between the current collector and MnO2 is enforced by the geometry. The difference between 100 nm maximum electrical path length through MnO2 (in prior work) and 5 μm (in this work) is significant. We nevertheless demonstrate with the data of Figure 6c that 10 μm lengths of these nanowires appear to be fully electrochemically addressed. In Situ Conductivity Measurements. In situ investigations of the electronic conductivity of MnO2, σ have been reported in just two prior publications to our knowledge. In 1998, Uchida et al.24 performed in situ measurements of spinel films with a thickness of 0.20 μm of LixMnO2 prepared by magnetron sputtering followed by thermal annealing to 700 °C. In that study, an interdigitated electrode array (IDA) was employed in dry 1.0 M LiClO4− propylene carbonate electrolyte. σ increased with decreasing potential by a factor of ≈2, from 10−5 S/cm to 2 × 10−5 S/cm, as fully delithiated MnO2 (at Eeq = 4.15 V vs Li/Li+) was equilibrated at Eeq = 3.55 V. These conductivity end points are indicated in Figure 7b. Subsequently the same group used electrostatic spray deposition to deposit spinel films of LixMn2O4 with a thickness of 0.70 μm.25 In this case, reduction of such films from Eeq = 4.3 V vs Li/Li+ to 3.3 V resulted in a transient increase in σ from 10−5 S/cm to 1.5 × 10−4 S/cm at 4.15 V, followed by a reduction in σ to a minimum value of 0.5 × 10−4 S/cm at 3.6 V. The three measured extrema in σ are indicated in Figure 7b.

Ionic conductivities have also been measured for LixMnO2 films as a function of x,26,27 but we are not able to measure ionic conduction using the experimental arrangement employed here. The results of our measurements are summarized in Figure 7. σ was measured for seven equilibrium potentials from +0.60 V vs MSE to −0.8 V which, as already demonstrated, span the full range of lithiation for LixMnO2. Nanowire arrays were equilibrated at each Eeq value for 60 s which, as shown in Figure 2d, provides ample time for equilibration to be achieved, and these seven Eeq values were measured in random order. The electrical conductivity, σ, is the reciprocal of electrical resistivity, ρ: σ = 1/ρ = l/RS, where S is the cross-sectional area of the nanowire (w × h), l is the length of the wire between the 2 contacts (10 μm), and R is the nanowire resistance. Two nuances in this calculation are the following: (1) The mean R value for a single nanowire is measured from the slope of the IV curve acquired for an array of 200 nanowires (Figure 2e) multiplied by the number of nanowires in the array (200). (2) The measured σ value is background-corrected by subtracting the conductivity of the solution, measured using the identical device architecture used for nanowire measurements except that no nanowires are present. The nanowire w and h values were measured using SEM and AFM, respectively. Four observations are the following: First, for all four sets of nanowires, σ increases monotonically with increasing x. This trend is the same as reported by Uchida et al. in their 1998 paper that probed spinel phase material,24 but we observe no peak in σ as reported in their 2000 paper.25 Second, σ depends on the nanowire width, w (the height of all nanowires is 47 ± 13 nmAFM data shown in Figure S2 of the Supporting Information). For example, the narrowest nanowire with a width of 277 nm has a conductivity that is 30 times higher than 3499


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Figure 9. Proposed growth mechanism for a δ-MnO2 nanowire prepared by pulsed electrodeposition using LPNE. (a) Voltage pulse program used for nanowire growth. (b) Current versus time plots for a sequence of 12 0.5 s growth pulses showing evolution of the growth rate as a function of time. (c) First current−time transient attained a time-invariant current limit corresponding to the diffusion-controlled limit characterizing the trench height and length. (d) Within the trench, diffusion is linear to the gold edge. (e, f). As conductive MnO2 sheets (black lines in f) grow from the gold electrode into the trench, Mn2+ within the trench is rapidly electrolyzed to form MnO2, resulting in a peaked current transient (e) that decays to the same trench transport limit seen for earlier pulses.

the widest, 870 nm nanowire at −0.80 V. Third, the ratio between σ for the fully lithiated and fully delithiated states, σLiMnO2/σMnO2, is also larger for narrower versus wider nanowires. σLiMnO2/σMnO2 is 1.3 for w = 870 nm whereas for w = 277 nm, it is 9.0. Since all of these LixMnO2 nanowires were produced by the same process, the disparities in σ and σLiMnO2/σMnO2 between nanowires of various widths are very surprising. Finally, these σ values, across all values of x and all nanowire widths, are higher than reported in previous work.24,25 What is the origin of the unexpected nanowire width dependence of σ and σLiMnO2/σMnO2? To answer this question, we examined the morphology of individual nanowires using TEM. Figure 8 (top) shows a cross section of w = 870 nm nanowire (x = 0) in which the gold edge, patterned by lithography (Figure 1a, step 1), and the solution edge of the nanowire are indicated. A ubiquitous feature of these nanowires when examined by TEM is a gradient in their electron density that we interpret as a gradient in nanowire porosity. Specifically, the nanowire porosity increases (the electron dense fraction decreases) as a function of distance from the gold edge of the nanowire. Plotted in Figure 8 (bottom) is the electron dense

fraction versus distance from the gold edge. These data are culled from many (>5) images of nanowires prepared in a single batch, and error bars indicate the density variability at various positions on a single nanowire and between nanowires. The decrease in dense fraction is 40% for the w = 780 nm nanowire across its width. This effect is also evident, but less pronounced, for the w = 460 nm nanowire. No significant gradient is seen for the w = 250 nm nanowire. But plots of the electron dense fraction versus distance are shifted to higher porosity (lower electron dense fraction) as the nanowire width is reduced. If attention is focused on the porosity of the nanowire near the gold electrode, for example, the electron dense fraction for the w = 780 nm nanowire is twice that of the 250 nm nanowire (Figure 8, bottom). What process can account for the existence of a porosity gradient within each nanowire as well as the systematic differences in porosity seen for nanowires of differing width? We propose that the mechanism shown in Figure 9 is responsible for both effects. A pulsed growth program (Figure 9a) is used to obtain MnO2 nanowires with a compact growth habit, avoiding diffusion-controlled deposition that causes dendritic growth.12 This pulsed growth scheme eliminates dendritic growth because between each 0.5 s growth pulse, the 20 s rest period provides time for the trench to refill with fresh 3500


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Figure 10. Description of a MnO2 system with water and periodic boundaries (dashed line). (a) Model of δ-MnO2 + H2O, (b) interlayer distance dependence of the binding energy of δ-MnO2 with and without a H2O molecule, (c) electronic band structure of δ-MnO2 + H2O, showing spin up (black) and spin down (red) channels, and (d) total and partial density of states for δ-MnO2 + H2O.

Mn2+ from the bulk solution. This program involves the application of a train of 0.50 s voltage pulses to +0.60 V vs MSEa potential at which MnO2 is electrodeposited. Current versus time transients for this process evolve as shown in Figure 9b: The initial current versus time transient is (Figure 9c) shows an increasing deposition current that plateaus at a welldefined current which we believe represents the cylindrical mass transport limit for Mn2+ entering the LPNE-fabricated trench. Within the trench, however, linear diffusion of Mn2+ is occurring to the surface of the gold electrode (Figure 9d). Subsequent voltage pulses show a peaked response (Figure 9e), and the amplitude of this current peak increases as a function of the number of pulses (Figure 9b). This peak is attributed to efficient, radial transport of Mn2+ within the trench to nascent MnO2 sheets growing from the gold electrode (Figure 9f). That is, the extra current comprising the peak is attributed to rapid, nonlinear transport of Mn2+ to the surfaces of the porous, growing MnO2 layer. This peak current then decays to the same trench-limited current seen in pulse #1, also as seen in Figure 9 b and e. This pulsed growth process causes a porosity gradient to be imprinted within each MnO2 nanowire because the number of growth pulses contributing to the lateral growth of MnO2 layers varies directly with the distance of the MnO2 from the gold electrode. Our data do not permit us to conclude how the density gradient within these nanowires influences their charge storage performance, but we believe the reduced range of electrical conductivities that are accessible as the nanowire width increases (Figure 7a) may be a symptom of the density gradient. Dense (nonporous) nanowire regions may “freeze” the Mn redox state by isolating it from Li+ charge compensating cations. This mechanism explains why the strongest disparity between the conductivity of nanowires of different widths is seen at negative potentials where Li+ intercalation is required for charge compensation of new Mn3+ centers. This hypothesis also provides an explanation for the variation in Csp seen for MnO2 nanowires as a function of width (Figure 6c): Slow Li+

intercalation is the rate-limiting step in accessing the Faradaic capacitance of this material, and this process is impeded in dense, low-porosity MnO2 present in wide nanowires relative to the higher porosity present in narrower nanowires. More work will be required to understand the fundamental origin of these disparities in Csp, σ, and σLiMnO2/σMnO2. Density Functional Theory (DFT) Calculations. DFT calculations were performed in an attempt to better understand the strong increase in σ seen with lithiation in Figure 7. The interaction between adjacent MnO2 layers is mainly via the van der Waals forces, and the calculated total energies are shown in Figure 10b as a function of the interlayer distance, d. Clearly, the size of binding energy, Eb, is very weak (less than 0.1 eV) when d is larger than 7 Å. As d decreases, Eb increases exponentially and reaches its maximum of 0.65 eV at deq = 4.45 Å. Further decrease of d leads to an increase of the Coulomb repulsion between MnO2 layers and henceforth a decrease of Eb. Note that recent studies40,41 reported the interlayer distance of about 7 Å, for MnO2 samples about 2.5 Å larger than our present data. This controversy is mainly due to the fact that δMnO2 is typically synthesized in solutions, and H2O molecules may exist in the space between MnO2 layers. To investigate the effect of water insertion, we put one H2O molecule into the layers in a (2 × 2 × 1) supercell as shown in Figure 10a. Similarly, we obtained the d-dependent binding energy of MnO2 + H2O, and it is intriguing that the optimized interlayer distance increases to 7.1 Å, and the magnitude of Eb reaches its maximum of 0.25 eV. The excellent agreement with experimental result of deq40 demonstrates a key role of H2O in the significant lattice expansion of MnO2. Further increasing the concentration of water (two H2O molecules in the same supercell) leads to similar results, indicating that the interlay distance is not sensitive to the concentration of water molecules. Moreover, the binding energy per H2O molecule with the MnO2 lattice is about 0.5 eV so water molecules should be rather stable in our δ-MnO2 samples. The electronic band structure of the bulk δ-MnO2 with the insertion of water 3501


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Figure 11. (a) Atomic structure of MnO2 + H2O + 1Li; (b) atomic structure of MnO2 + H2O + 3Li green isosurface shows the wave function features of the gap states near Fermi; (c) and (d) band dispersions corresponding to the structures of a and b; (e) and (f) total and partial density of states for (a) and (b). (g) Charge redistribution of MnO2 + H2O + Li at 0.1 eV/ Å, electron accumulation and depletion are represented by green and black, respectively. (h) Li atom concentration dependence of electron effective mass, m*.

is shown in Figure 10c. Bulk δ-MnO2 without H2O is an indirect band gap semiconductor with its valence band minimum (VBM) between K → Γ and its conduction band maximum (CBM) between K → Γ in the Brillouin zone. Interestingly, water molecules seldom change the bands near VBM and CBM but induce gap states right below the Fermi level in both spin up and down channel as also evidenced by the partial density of states (PDOS) of H2O molecules in Figure 10d. Without considering these gap states, the band gap of δ-MnO2 is about 2.4 eV, in agreement with previous studies.41−43 To explore the origin of the increasing conductivity by inserting lithium atoms into δ-MnO2, we investigated three models with different Li concentrations. Since typical energy minimization procedures may end up with the local minima in the structural optimization, we conducted AIMD simulations for MnO2 + 2 H2O + nLi (n = 1, 2, 3) at room temperature

(300 K) and maintained its thermal equilibrium for 5 ps. The atomic structures at the end of AIMD procedure were fully relaxed at 0 K for the studies of other physical properties. The optimized atomic configurations of MnO2 + 2H2O + nLi are shown in Figure 11a (n = 1) and b (n = 3), where it is clearly demonstrated that a single Li atom in the supercell prefers to stay between the two H2O molecules and additional Li atoms bind with both H2O and the MnO2 layer. The electronic band structures of these two models show that the presence of the Li atoms induces new gap states right below the Fermi level. The band dispersion corresponding to these structures is given in Figure 11c and d. As demonstrated by the PDOS in Figure 11e and f, the induced new gap states mainly come from Mn atoms; meanwhile, the old gap states induced by H2O molecule shift down to deeper energy level with the increasing Li concentration due to the hybridization between Li atoms, H2O molecules, and substrates. In particular, 3502



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the number of gap states is proportional to the number of Li atoms in the 3 × 3 × 1 supercell. The wave function features of these gap states are displayed by the green isosurfaces in Figure 11a and b. Clearly, they are rather localized in the Mn sites of substrates, even though they originate from the conduction states of MnO2. Moreover, the Li atoms transfer their electrons to the substrates and also cause strong rehybridization among conduction states. This is more clearly manifested in the plot of charge difference [ρ(MnO2 + 2H2O + 3Li) − ρ(MnO2) − ρ(2H2O + 3Li)] in Figure 11g, where one can see significant charge depletion near Li atoms as depicted by the large black isosurface. Most of these electrons are transferred to the Mn sites in the substrates, as shown by the green isosurfaces. Although the Li-induced gap states appear to be fully occupied, they should provide the conductive channels for electrons to flow under the influence of inhomogeneous defects or finite boundaries that cause a downward shift of the Fermi level. Furthermore, the gap states also become more dispersive, and their effective masses decrease with the Li concentration, as depicted in Figure 11c and d. We also found that the effective mass of the holes in these gap states decreases significantly with increasing Li concentration. Therefore, a high Li concentration not only increases the number of conduction channels but also enhances the mobility of effective carriers in the gap states. Collectively, these two effects explain our observation that σ increases with x for our LixMnO2 nanowires.

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Article

ASSOCIATED CONTENT

S Supporting Information *

XRD and AFM data as described. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b00912.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS All aspects of this work were supported as part of the Nanostructures for Electrical Energy Storage (NEES II), an Energy Frontier Research Centre (EFRC) funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under award number DESC0001160 with the exception of the following: XPS data was acquired by Y.L. and J.C.H. who are supported by the U.S. Department of Energy, Office of Science, Basic Energy Sciences under award DE-FG02-96ER45576. Density function theory was carried out by H.W. and R.Q.W. who are supported by the National Science Foundation Center for Chemical Innovation on Chemistry at the Space−Time Limit (CaSTL) under NSF grant CHE-0802913, and by XSEDE for computing time. Valuable discussions with Professor Phil Collins and Jung Yun Kim are gratefully acknowledged. SEM, TEM, and XRD data were acquired using instrumentation of the LEXI facility (lexi.eng.uci.edu/) at UCI.

CONCLUSIONS

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LixMnO2 nanowires show relatively high conductivities that are enhanced further by reduction of Mn4+ centers to Mn3+ accompanied by the intercalation of Li+. The σ values measured here are significantly higher than measured previously for MnO2 thin films prepared using either of two methods,24,25 but the origin for this enhancementwhich is as high as a factor of 30 for small nanowires that are reducedis not apparent from our data. Enhancement of the conductivity with Li insertion is consistent with DFT calculations that show an increase in the number of conduction channels available in the lithiated material coupled with increases in the mobility of effective carriers. Two surprising results observations reported here are the pronounced disparities in the electrical conductivity, σ, and the conductivity ratio, σLiMnO2/σMnO2, seen between LixMnO2 nanowires as a function of width. TEM measurements reveal the existence of a gradient in the porosity of these nanowires that may be caused by the pulsed electrodeposition process using to prepared them. The porosity of the MnO2 in a nanowire decreases near the gold current collector as the nanowire becomes wider. We hypothesize that access for Li+ to Mn centers may be blocked in the dense MnO2 found in the widest nanowires, with widths of 400 nm or more, blocking the reduction of Mn4+ centers in the material. Depressed charge storage capacities, Csp, are also observed as the nanowire width increases and the mean porosity decreases. Collectively, our data show that an extremely high degree of dispersion for MnO2 nanosheets within the nanowire, associated with high porosity as measured by TEM, is required in order to achieve state-of-the-art charge storage performance for this material.

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