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Modifying Current Collectors to Produce High Volumetric Energy Density and Power Density Storage Devices Hadi Khani, Timothy J Dowell, and David Wipf ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b03606 • Publication Date (Web): 04 Jun 2018 Downloaded from http://pubs.acs.org on June 4, 2018

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ACS Applied Materials & Interfaces

Modifying Current Collectors to Produce High Volumetric Energy Density and Power Density Storage Devices

Hadi Khani†, Timothy J. Dowell‡, David O. Wipf ‡* †

Materials Science and Engineering Program and Texas Materials Institute, The University of

Texas at Austin, Austin, TX 78712, USA ‡

Department of Chemistry, Mississippi State University, Mississippi State, MS 39762, USA

*Corresponding author: tel: +1 662 325 7608; email: [email protected]. Abstract Text: We develop zirconium-templated NiO/NiOOH nanosheets on nickel foam and polypyrrole-embedded in exfoliated carbon fiber cloth as complementary electrodes for an asymmetric supercapacitor device. We achieve high volumetric energy and power density by the modification of commercially available current collectors (CCs). The modified current collectors provide the source of active material, actively participate in the charge storage process, provide a larger surface area for active material loading, need no additional binders or conductive additives, and retain the ability to act as the current collector. Nickel foam (NF) current collectors are modified by use of a soft-templating/solvothermal treatment to generate NiO/NiOOH nanosheets, where the NF is the source of Ni for the synthesis. Carbon-fiber cloth (CFC) current collectors are modified by an electrochemical oxidation/reduction process to generate exfoliated core-shell structures (ECFC). Electropolymerization of pyrrole into the shell structure produces polypyrrole embedded in exfoliated core shell material (PPy@rECFC). Battery-type supercapacitor devices are produced with NiO/NiOOH@NF and PPy@rECFC as positive and negative electrodes, respectively, to demonstrate the utility of this approach.

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Volumetric energy densities for the full-cell device are in the range of 2.60–4.12 mWh cm−3 with corresponding power densities in the range of 9.17–425.58 mW cm−3. This is comparable to thinfilm lithium-ion batteries (0.3–10 mWh cm−3) and better than some commercial supercapacitors (< 1 mWh cm−3).1 The energy and power density is impressive considering that it was calculated using the entire cell volume (active materials, separator, and both current collectors). The fullcell device is highly stable, retaining 96 % and 88 % of capacity after 2000 and 5000 cycles, respectively. These results demonstrate the utility of directly modifying the current collectors and suggest a new method to produce high volumetric energy density and power density storage devices. KEYWORDS: supercapacitor, polypyrrole, nickel oxide, zirconium, exfoliated carbon fiber, volumetric energy density 1. Introduction The rapid growth of portable electronics, electrified transportation, and renewable energy conversion over the last decade has been both enabled and limited by energy-storage devices. Foreseeable demands in these markets continues to motivate researchers to develop high volumetric power and high energy density technologies that can accommodate the enormous range of energy and power requirements. Among the various electrochemical energy-storage devices, supercapacitors (SCs) and rechargeable batteries are the most promising as power-type and energy-type candidates. This work aims to combine the performance metrics of SCs and batteries by proposing a battery-type supercapacitor device that enables both high energy and power densities. Supercapacitors (SCs) are receiving increased attention as components of power systems due to their simple energy storage mechanism, high power density (10–1000 mW cm−3), long 2 ACS Paragon Plus Environment

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cycle-life (> 10,000 cycles), and favorable safety considerations. However, SC materials are often limited by low specific volumetric energy densities (EV) (< 1 mWh cm−3).1-2 An additional limitation is that SCs are typically produced in a similar way as batteries: requiring binders, electrolyte, separators, and current collectors (CCs). This inactive material can decrease gravimetric and volumetric energy density fivefold when fabricated into full-cell devices.3 The impact that inactive material will have on the performance of supercapacitor material is often overlooked.3 This is easily seen for three commonly encountered current-collector materials. Metal foil (e.g. 25–50 μm thick Al, Ti, Ni, or stainless steel) is employed as a CC. This foil, with density from 2.7 to 8.9 g cm−3,4-5 reduces the gravimetric and volumetric performance of the device; even when commercial loadings (e.g. carbon-based SC: ≈ 10 mg cm−2, 0.3–0.8 g cm−3, and 100–200 μm thick) of active materials are used.3-4, 6-7 Alternatively, metal foam and carbon-fiber cloth are less dense (< 1 g cm−3); resulting in lower volumetric performances (F cm−3 or mA h cm−3) if the entire electrode volume is not utilized (such as when thin films of active materials are grown directly onto the CCs) since the large void space reduces the gravimetric capacitance when filed with excess electrolyte. Enabling the current collectors to directly participate in the charge storage process reduces the typical dead-space/dead-weight problem. We show that high-performance positive and negative electrode materials are produced by the modification of commercially available current collectors, which are then the active electrode materials. These materials demonstrate exceptional areal and volumetric energy densities when fabricated into a full-cell device and retain extended cycle life and high power densities. A great majority of recent reports have focused on increasing the energy and power densities of just one electrode without considering the complementary electrode. Combining

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positive and negative electrodes with dissimilar performance leads to dramatically reduced rate performance of the full-cell device. Simultaneous development of both positive and negative electrode materials with high-performance metrics and with complementary properties can minimize performance reductions. Further, the energy density (E cm−3) of supercapacitor devices can be enhanced by increasing the specific capacitance (C cm−3) or the output voltage (volt) according to E = ½CV 2.8 Asymmetric electrochemical supercapacitors (ASC) use active materials with different potential windows. The resulting wider potential window (ca. 2 V) increases energy density when compared to traditional electric-double-layer capacitor (EDLC) materials.9 Many high-performance positive-electrode materials have been developed, including conductive polymers (polyaniline, poly(phenylenevinylene), polypyrrole, etc.),10 metal oxides/hydroxides (RuO2, MnO2, etc.),11-12 or metal sulfides (NiS, CoS, MoS, etc.).13 Nickelbased materials (NiO, NiOOH, and Ni(OH)2) have a layered structure and large interlayer spacing, which enable large pseudocapacitance.14-15 However, practically achievable specific capacitances and rate capabilities of such materials are often significantly lower than theoretical predictions due to poor electrical conductivities and surface-limited Faradaic charge storage (i.e. low utilization of the active material). To alleviate these problems, we have grown ultrathin NiO and NiOOH nanosheets directly from conductive nickel foam substrates. To the best of our knowledge, this is the first report on the direct growth of both NiO and NiOOH nanosheets on the NF without requiring an additional nickel source. This method has several advantages: the porous structure of the NF current collector allows ready electrolyte access to the NiO and NiOOH nanosheets, directly growing materials from the NF improves electrical contact, the resulting ultrathin and mesoporous NiO and NiOOH nanosheets facilitates ionic transport into


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the bulk of the material structure, the hierarchical mesoporous structure of NiO/NiOOH nanosheets accommodates volume changes from Faradaic reactions, improves the cycling stability, and avoids the use of inactive components. The last can be compared to slurry-coating fabrication methods, where a large portion of the active material surface can become blocked by binders and conductive additives. The intrinsically low operating potential window (∼ 0.5 V) of our positive NiO/NiOOH nanosheets material lead us to adopt an ASC device configuration, which produces a larger potential window and increases the energy density of the full-cell device. Thus, it became necessary to develop a suitable negative electrode material. Far fewer negative (anode) electrode materials have been studied compared to positive materials. The most common negative electrode materials include high surface area activated carbons, carbon nanotubes, graphene, carbon onions, and their hybrid composites.16 These carbon-based materials store charge in the electric double layer (EDL) and therefore have specific capacitances up to an order of magnitude lower than that of their pseudocapacitive positive electrode counterparts.17 As a result, when employed as anodes in asymmetric cell configurations, excess carbon-based active materials are used, which lowers the overall device performance. Recent efforts have shifted to pseudocapacitive negative electrode materials, such as iron and vanadium oxide/hydroxides (Fe2O3, Fe3O4, FeOOH, and V2O5) and other metal oxides and nitride compounds (SnO2, MoO3, Bi2O3, and VN).18-22 However, these materials typically suffer from low specific capacitance (≤ 300 F g−1) and poor cycling performance (≤ 2000 cycles).18 Polypyrrole (PPy) has attracted a renewed interest as a negative electrode material in alkaline electrolytes due to its sufficiently negative potential window (−0.6 to −0.8 V) and good cycling performance as well as a high theoretical specific capacitance (620 F g−1).19, 23-24 Furthermore, the conductivity is greatest when



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the material is in a fully doped state (i.e. full charge under negative polarization), which is an advantage for anode materials.23 However, reported specific capacitance values for PPy are often far below those predicted by theory.23 This is ascribed to the high packing density of pyrrole, which prevent the access of dopant to inner-sites (especially in thick coatings), and low utilization of active material.23, 25 Likewise, the loss of electrical conductivity in the undoped (discharged) state results in poor roundtrip Coulombic efficiency and lower specific capacitances, requiring conductive additives. Also, the polymer swells and shrinks upon charge and discharge (doping and undoping), resulting in mechanical separation from the current collector and low cycling performances.25 To alleviate these performance limitations and to design an anode material with charge storage capabilities on par with positive pseudocapacitive materials, we demonstrate a multidimensional electrode design based on permanently doped PPy embedded in an exfoliated core-shell carbon fiber cloth electrode (PPy@rECFC). This method enables the carbon-fiber cloth to participate in charge storage, provides an open framework to introduce additional active material (i.e. PPy), generates interconnecting carbon filaments in the exfoliated shell enhancing electron mobility, accommodates structural changes of PPy during cycling, and avoids thick PPy films, improving ion transport. In this report, an asymmetric charge storage device is fabricated by direct modification of commercially available current collectors by (a) growth of NiO/NiOOH nanosheets from nickel foam (NiO/NiOOH@NF) via a zirconium hydroxide soft-templated thermal activation and (b) electrochemical exfoliation-reduction-polymerization generating a PPy embedded in an exfoliated core-shell carbon-fiber cloth (PPy@rECFC). These materials both demonstrate highly accessible surface areas, low electrical resistivity, and redox-rich active sites making them excellent positive and negative high performance supercapacitor electrode materials. The high


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retention of specific capacities for NiO/NiOOH@NF (0.613 mAh cm–2 and 0.326 mAh cm–2 at 1 mA cm–2 and 40 mA cm−2, respectively) and PPy@rECFC (0.594 mAh cm–2 and 0.467 mAh cm– 2

at 1 mA cm–2 and 40 mA cm−2, respectively) demonstrate their high-rate performance. The full-

cell ASC was reversibly cycled in the potential window of 0.0 to 1.4 V and demonstrates an excellent volumetric specific capacity of 4.99 mAh cm–3. The full-cell device achieves a superb specific energy of 4.12 mWh cm−3 at a specific power of 9.17 mW cm−3 and a retention of 2.60 mWh cm−3 when the power is increased to 425.58 mW cm−3. 2. Experimental section 2.1. Preparation of Positive Electrode All chemicals were reagent grade and used without further purification. NiO/NiOOH nanosheets were grown on nickel foam through a two-step electrodeposition and heat treatment method. Ni foam substrates (1 cm × 1 cm × 0.05 cm, Shanghai Winfay Metal & Plastic Manufacturing Co) were sequentially degreased in hot acetone, cleaned in hydrochloric acid solution (3 M) under mild sonication for 10 min, washed in DDI water, rinsed in absolute ethanol, and dried under vacuum at ambient temperature. A zirconium hydroxide film was cathodically deposited onto the NF substrate from a 10 mM ZrOCl2 and 150 μL (7.5 mL/L) poly(diallyldimethylammonium chloride) (PDDA, MW = 400,000–500,000, Aldrich) methanol– water solution (5 vol % water). The electrodeposition was accomplished in 30 mL of solution at a NF working electrode, a platinum gauze auxiliary electrode, and an Ag/AgCl reference electrode, optimally at a cathodic current density of 5 mA cm−2 for 10 min. After electrodeposition, the samples were dried at room temperature for 10 h and then under vacuum at 60 °C for 5 h. After drying, samples were soaked in 6 M KOH for 10 min, heated in air to 250 °C (10 °C min–1 heating rate) for 2 h, cooled to 100 °C, and thoroughly rinsed with DDI water. 7 ACS Paragon Plus Environment


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Samples prepared under these conditions are named Zr-NF-250, where the numeric suffix refers to the final heating temperature. 2.2. Preparation of Negative Electrode Exfoliated carbon fiber cloth (ECFC) was prepared by a two-step electrochemical oxidation-reduction route. A 1 cm × 1 cm, 0.35 mm thickness, areal mass: 23 mg cm−2 piece of carbon fiber cloth (CFC) (CC4 Plain, Fuel Cell Earth, Woburn, MA) was degreased/wetted by mild sonication in acetone for 10 min. Electrochemical oxidative exfoliation was performed in 50 mL of aqueous 0.05 M (NH4)2S2O8. A DC voltage of 5 V was applied to a CFC anode and graphite plate cathode for 10 min while vigorously stirring the solution. The resulting oxidized/exfoliated core-shell CFC (ECFC) electrode was rinsed with DDI water several times. Electrochemical reduction of the ECFC was carried out under potentiostatic conditions in 4 M KOH electrolyte for 30 min at −1.2 V (vs. Ag/AgCl) and with a Pt mesh auxiliary electrode. The resulting reduced ECFC electrode (rECFC) was soaked overnight in DDI water to remove residual electrolyte. Polypyrrole was embedded into the rECFC electrode by electropolymerization in an aqueous solution of 0.15 M pyrrole and 0.075 M p-toluenesulfonic acid (p-TSA) with a Pt auxiliary and Ag/AgCl reference electrode. The electropolymerization used a square-wave potential consisting of equal length (100 ms) steps between 0.00 and 0.90 V. Best results were obtained after 15 min of electropolymerization. The resulting (PPy@rECFC) electrodes were rinsed and dried in a vacuum oven at 50 °C for 24 h. The mass loading of PPy in PPy@rECFC was determined to be ≈ 1.5 mg cm−2. Exact mass loadings were determined by the mass difference between rECFC electrodes before and after electropolymerization via a microbalance with an accuracy of 0.01 mg.


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2.3. Electrochemical Tests Both three-electrode (half-cell) and two-electrode (full-cell) configurations were used in electrochemical measurements. The electrochemical performance of the cells was characterized by cyclic voltammetry (CV), galvanostatic charge–discharge (CC) tests, and electrochemical impedance spectroscopy (EIS). Three-electrode cell measurements were made using NiO/NiOOH@NF or PPy@rECFC as the working electrode, Pt mesh as the auxiliary electrode, and a double-junction Ag/AgCl reference electrode. The electrolyte was aqueous 4 M KOH. For two-electrode cell measurements, the full-cell device was assembled into a coin-type cell with NiO/NiOOH@NF as the positive electrode and PPy@rECFC as the negative electrode, cellulose filter paper (150 μm thickness) as the separator, and 4 M KOH aqueous solution as the electrolyte (Figure S1). Cycling tests were carried out at room temperature for 5000 cycles at a current density of 30 mA cm−2 over the potential range of 0 to 1.4 V using an electrochemical workstation (Solartron Analytical 1470E Cell Test System). EIS measurements employed a frequency response analyzer (Solartron Model SI 1250) and used an AC perturbation of 5 mV at frequencies of 10 kHz to 0.01 Hz while polarizing the electrode at the midpoint of the charging plateau potential. The areal specific capacities (Cs) were calculated from the charge-discharge curves according to the following equation: ∆

(1)

where Cs is the areal specific capacity (mAh cm−2) of a single electrode, I is the charge-discharge current (mA), ∆t is the discharge time (h), and S represents the geometric surface area of a single electrode. For the assembled full-cell device, the volumetric specific capacities (CV, mAh cm−3)



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is reported, where the geometric surface area in eq 1 is replaced by the device volume (V°). To balance the charge in the full-cell device (Q+ = Q−), the electrode size was balanced as follows: (2) where S is the surface area of a single electrode (cm2) and Cs (mAh cm−2) is the areal specific capacity of the negative or positive electrode. The total volume of the full-cell device was considered for the full-cell performance calculations. The calculation of volumetric specific energy and power is based on the following equations: (3) ∆

(4)

where E is the volumetric specific energy (J cm−3), I is the discharge current density (A cm−3), V is the discharge voltage (V), P is the volumetric specific power (W cm−3), and ∆t refers to the total discharge time (s). The volumetric specific capacitance (Cp) of the full cell device is calculated from the following equation by substituting the specific energy obtained from eq 3. ∆

(5)

where Cp is the volumetric specific capacitance (F cm−3), E is the volumetric specific energy (J cm−3), and ∆V is the working potential of the cell (V). 2.4. Material Characterization The morphology, structure, and composition of the samples were characterized by fieldemission scanning electron microscopy (FESEM, Zeiss SUPRA 40), transmission electron microscopy (TEM, JEOL 2100 200 kV), optical microscopy (Olympus BH-2), energy dispersive spectroscopy (EDX), X-ray diffraction (XRD, Rigaku SmartLab, Cu Kα radiation, λ = 1.5406 Å), and X-ray photoelectron spectroscopy (XPS, Thermo Scientific K-Alpha). Raman spectra were 10 ACS Paragon Plus Environment

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obtained with a Horiba LabRam HR800 confocal Raman microscope system and a 633 nm Raman excitation laser, and IR spectra were recorded with a Thermo Scientific Nicolet 6700 FTIR. 3. Results and Discussion 3.1. Structural Characterization of the Positive Electrode The synthetic zirconium redox mediated soft-templating method for producing NiO/NiOOH nanoflakes on a NF substrate is illustrated in Figure S2. The initial step is the production of a uniform film of zirconium hydroxide on the NF surface by cathodic electrodeposition. The precursor ZrOCl2·8H2O undergoes hydrolysis in aqueous solutions to form the hydrated tetramer-hydroxo-complex [Zr4(OH)8(H2O)16]Cl8.26-27 The coordinated water can dissociate (eq 6), forming acidic solutions: [Zr4(OH)8(H2O)16]8+ → [Zr(OH)2+x·(4-x)H2O)]4(8-4x) + 4xH+

(6)

A cathodic current at the NF working electrode reduces water through the following reactions: 2H2O + 2e− → H2 + 2OH− and O2 + 2H2O + 4e− → 4OH−. The resulting increase in pH at the electrode surface leads to formation of negatively charged colloidal hydrated zirconium hydroxide, Zr(OH)4·xH2O (isoelectric point 6.7).28-29 The electrostatic interaction between negatively charged particles of zirconium hydroxide and the negatively polarized NF electrode will prevent uniformly coating zirconium hydroxide onto the NF surface.30 Better adhesion and uniformity of the zirconium hydroxide coating is achieved by adding a cationic surfactant, PDDA, to the zirconium deposition solution. It has been reported that formation of a uniform film is driven by heterocoagulation due to Coulombic attraction between PDDA and the negatively charged colloidal particles at the electrode surface.31 Figure S3 shows the SEM images of bare and zirconium hydroxide-coated NF. As the SEM images reveal, the NF is 11 ACS Paragon Plus Environment


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uniformly coated by a 2-3 µm thick film by the galvanostatic electrodeposition method. The XRD pattern of zirconium hydroxide coated-NF sample (Figure 1a) shows two very broad peaks centered at 2θ = 32° and 55°, which agree with literature reports for hydrated zirconium hydroxide.31-32 Formation of NiO and NiOOH nanosheets proceeded by immersing the dried zirconium hydroxide coated-NF electrodes into a 6 M KOH solution for 10 min followed by heating in a furnace. Figure 2 shows SEM images of Zr-NF-250 produced by this method. As seen, two unique morphological regions are present on the NF surface: rose-like (cf. Figure 2i) and flakelike (cf. Figure 2h) micro-domains, each of which is composed of smaller bundles of nanosheets with thickness of ≈ 100 nm. Grazing-angle-XRD analysis was used to determine the chemical composition of the ZrNF-250 nanosheets (Figure 1b). The diffraction peaks are well indexed to the standard peaks for γ-NiOOH (JCPDS card JCPDF: 06-0075) and NiO (JCPDS card JCPDF: 089-7101), indicating both NiO and NiOOH are present. SEM-EDX line scans shows that the composition of the roselike and flake-like regions is consistent with NiOOH and NiO, respectively, based on O/Ni atomic ratios (Figure S4). Figure 3 shows the HRTEM images of Zr-NF-250. As indicated by red outlined areas (Figure 3c: a magnified region of Figure 3b), many pores are formed in the NiO/NiOOH nanosheets. This suggests that, in addition to the microporous nature of the rose-like and flakelike domains, individual nanosheets contain mesoporous cavities (≈ 1–3 nm) throughout their structure. HRTEM images show that the nanosheets are polycrystalline and consist of ≈ 4–5 nm cubic grains (see yellow outlined-areas in Figure 3b), producing vast numbers of exposed edge and surface sites (Figure 3b and c). The disordered packing of these randomly oriented cubic


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grains is the likely cause of the observed mesoporous cavities throughout the nanosheet structure. HRTEM images taken from different nanosheets (either rose-like or flake-like origin) (Figure 3d) display lattice fringes with spacing of 0.343 and 0.237 nm, which matches well with the (006) and (102) planes of NiOOH. Lattice fringes with spacing of 0.241 and 0.208 nm are also observed, which correspond to the (003) and (202) planes of NiO. The Ni 2p XPS spectra of the Zr-NF-250 sample can be fit, via a Gaussian fitting method, to two spin−orbit doublets (i.e. 2p1/2 and 2p3/2) and two shake-up satellites (Figure S5a). The Ni 2p3/2 components observed at 854.2 and 855.7 eV correspond to Ni2+ (NiO) and Ni3+ (NiOOH) species, respectively, while the corresponding Ni 2p1/2 bands for Ni2+ and Ni3+ appear at 871.8 and 873.2 eV, respectively. The overall conclusion of XPS analysis is that the near-surface region of the as-prepared Zr-NF-250 sample contains both NiOOH and NiO. Moreover, the ratio of XPS peaks for Ni2+ 2p (NiO) and Ni3+ 2p (NiOOH) shows that the ratio of NiOOH (Ni3+) to NiO (Ni2+) is 1.25. The growth process of densely packed NiO/NiOOH nanosheets was examined by varying the heating time from 30 to 120 min on the zirconium hydroxide-coated NF electrodes. SEM images of electrodes heated at 250 °C for 30 min (Figure S6) show numerous ≈ 1 μm size cracks in the zirconium hydroxide gel coating as it shrinks due to water loss. Initial growth of NiO flakes within the cracks is also observed. We assume that the cracks provide a route for atmospheric oxygen and KOH solution to reach the surface of NF and begin the growth of NiO flakes. Electrodes prepared by heating for 60 min at 250 °C show the formation of the rose-like (NiOOH) micro-regions (≈ 10×10 μm2) within the walls of NiO flakes (Figure S7). Formation of new cracks (< 400 nm) in the zirconium hydroxide gel is likely caused by additional water loss and mechanical pressure from the growth of NiOOH beneath the film. The zirconium hydroxide



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gel is in intimate contact with the rose-like structures, presumably capable of acting as both physical template and oxidizing agent. SEM images of electrodes prepared via a prolonged heating period (120 min at 250 °C) show continued growth of material in flake-like (≈ 5–10 μm across) regions and larger, more densely packed rose-like structures when compared to 30 and 60 min heating. However, heating times beyond 120 min at 250 °C resulted in electrodes with inferior electrochemical performances and were not further examined. Oxidation of the NF substrate generates the NiO and NiOOH nanosheets and nanoflakes by some or all of the mechanisms in eqs 7-11. Figure S6 shows that NiO nanosheets begin forming in-between cracks in the zirconium hydroxide coating. Additionally, EDX data (Figure S4), show that metallic nickel underneath the zirconium hydroxide gel is converted into NiOOH nanoflakes that then push the zirconium hydroxide gel away from the electrode surface. Our interpretation of the process of oxidizing Ni to form NiO and NiOOH relies on locally distinct oxidizing conditions. The nickel substrate is oxidized via a redox reaction between metallic nickel and zirconium(IV) hydroxide to generate NiO/Ni(OH)2 and zirconium(III) hydroxide (eqs 8 and 9). Cracks in the drying zirconium hydroxide gel allows NiO/Ni(OH)2 to form rapidly in this region. However, the NiO/Ni(OH)2 formed beneath the zirconium hydroxide gel micro platelets is subject to a slower, spatially impeded growth mechanism. The predominance of NiO flakes between the cracks is proposed to be due to dehydration of nickel hydroxides by exposure to the heated alkali solution (eqs 7-9). However, the more oxidized, predominately nickel(III) oxyhydroxide (NiOOH) rose-like regions, are proposed to form due to the presence of Zr(IV) (from zirconium hydroxide gel, eqs 10 and 11). Following the heating period (120 min), any residual zirconium hydroxide platelets are completely removed by rinsing with DDI water. Ni + ½O2 → NiO

(7)


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Ni + 2Zr(OH)4 → Ni(OH)2 +2Zr(OH)3

(8)

Ni(OH)2 → NiO + H2O

(9)

Ni(OH)2 + Zr(OH)4 → NiOOH + Zr(OH)3 + H2O

(10)

NiO + Zr(OH)4 → NiOOH + Zr(OH)3

(11)

XPS measurements confirm that Zr(OH)4 (Zr(IV)) is reduced during the oxidation of nickel foam. Figure S5b shows an XPS spectrum of the Zr-NF-250 sample prior to removing the residual dried zirconium hydroxide platelets (by rinsing). The photoelectron peaks associated with the Zr 3d5/2 components are at 181.7 and 182.6 eV and correspond to Zr(III) (as Zr(OH)3) and Zr(IV) (as Zr(OH)4) species, respectively. The Zr 3d3/2 bands for Zr(III) and Zr(IV) appear at 184.1 and 184.7 eV, respectively, consistent with reported literature values.33-34 The area of the Zr(III) 3d fitted peaks is three times larger than the Zr(IV) 3d peaks, which is consistent with our proposed redox mechanism for oxidation of nickel foam to NiO/NiOOH nanoflakes by Zr(IV). The positive impact of the zirconium hydroxide gel coating on forming of NiO/NiOOH nanoflakes is evident by subjecting a bare NF electrode to the same KOH-heat treatment (250 °C for 120 min). SEM images (Figure S8) of this sample show sparsely packed NiO microsheets with thick individual layers. These may have originated from the oxidation of nickel foam by dissolved oxygen (eq 7). Additionally, the NF surface is extensively pitted, consistent with a corrosion process. We assume that, without the zirconium hydroxide gel, NiO microsheets form by precipitation of dissolved nickel ions (from the concentrated KOH solution) generated via an oxidative corrosion process. The zirconium hydroxide gel coating is thus providing the benefit of dynamic physical templating and oxidation to form NiO/NiOOH nanoflakes. The effect of heating temperature on the surface morphology of the final product was examined after treatment in KOH and heating at 200 °C (Zr-NF-200) and 300 °C (Zr-NF-300)



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for 120 min. SEM images of Zr-NF-200 (Figure S9a-c) reveal that the redox reaction between Zr(OH)4 and nickel substrate to form the rose-like regions (NiOOH nanoflakes) has not taken place and the formation of flake-like regions (NiO) has been significantly retarded. SEM images of Zr-NF-300 (Figure S9d-f) reveal that the final morphology of the sample evolves into a layer of spherical NiO aggregates lacking rose-like or flake-like regions. These observations demonstrate that both temperatures of 200 °C and 300 °C do not achieve high surface area NiO/NiOOH nanoflakes. Treatment temperatures of 225 °C and 275 °C also result in inferior electrochemical performances compared to Zr-NF-250 samples. 3.2. Structural Characterization of the Negative Electrode 3.2.1. Preparation of Exfoliated Core-Shell Carbon Fiber Cloth Several chemical, electrochemical, or thermal based processes for preparing EDLC-type supercapacitors by modification/activation of carbon-fiber-cloth electrodes have been reported.3544

Chemical and thermal oxidation/exfoliation routes are typically used to modify the entire

electrode (i.e. complete oxidation of entire carbon fiber), resulting in a singular type of morphology. In contrast, electrochemical routes can offer more selectivity in the produced morphology (i.e. core-shell or complete exfoliation, respectively) over the oxidized morphology. However, anodically generated exfoliated layers on the carbon fiber surfaces are typically less than 100 nm thick, which is small compared to the carbon fiber diameter of 7-8 μm.37-38, 42 This suboptimal utilization can be attributed to electrical insulation by the exfoliated/oxidized-shell. Chemical, thermal, or electrochemical reduction of the oxidized carbon shell is often used to reestablish the electrical conductivity of the material. However, published results show that the active material thickness still remains less than 100 nm after reduction of the oxidized/exfoliated shell, likely because of the poor conductivity of thicker shells.36 16 ACS Paragon Plus Environment

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As discussed above, asymmetric cells are limited by the lower specific capacity of EDLC materials as compared to pseudocapacitive (Faradaic) material. Therefore, to successfully employ CFCs in an asymmetric cell using our proposed direct modification of current collectors, greater utilization of the carbon fiber as well as the incorporation of additional active material are required. With this in mind, we employ a three-step oxidation-reduction-polymerization route (Figure S10) to generate polypyrrole embedded in a carbon fiber core-shell electrode (PPy@rECFC). The core-shell structure is produced by an oxidative exfoliation followed by a reduction step. Electrochemical exfoliation of carbon materials in dilute aqueous electrolytes is known to proceed via several simultaneous processes. During exfoliation (at 5 V vs. OCP), a large amount of gas is evolved from water splitting and electrolyte decomposition, which can mechanically exfoliate individual CFs and lead to the observed cracks between carbon filaments bundles (Figure 4).45 In addition, electrochemically generated species, such as hydroxide and sulfate radicals, can attack the graphitic regions of the CFs and introduce oxygen containing functional groups (OCFGs) on graphene layers.46-49 Furthermore, electrochemically generated persulfate ion is strongly oxidizing (E° = 2.01 V) as well as its peroxide decomposition product (E° = 1.80 V) and these are capable of introducing OCFGs during the exfoliation process.48-49 Optical microscopy of the ECFCs produced by the process above reveal that individual CFs have undergone structural changes to form a core-shell structure composed of an ≈ 3–4 μm thick translucent red-brown graphitic-hydrogel shell surrounding an intact carbon fiber core of ≈ 4–5 μm diam, Figure 4a and c. The overall diameter of the fibers has, likewise, increased from ≈ 7 μm (pristine CFC) to ≈ 10–13 μm (Figure 4a, inset). The distal ends of individual CFs show a completely translucent red-brown hydrogel indicating a complete loss of the carbon fiber core 17 ACS Paragon Plus Environment


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(Figure 4b). This is likely due to the geometrically more electrolyte-accessible tip, which allows both axial and perpendicular diffusion of ions into the tip, compared to solely perpendicular diffusion at the shaft of the fiber (Figure 4b, inset). A second type of ECFC is also observed, more swollen with large cracks traveling along the fiber axis (Figure 4c, yellow dashed lines). The later may be due to the strong oxidizing nature of the electrolyte (i.e. persulfate anions and in situ generated: hydrogen peroxide, hydroxyl radicals, and sulfate radicals), which continues to oxidize the exfoliated shell after electrochemical oxidation has ceased (because the shell has grown thick enough to become insulating). The chemical oxidation further expands the shell, producing fractures. The ECFC material changes appearance from the grey luster of CFC to a matte black with a red-brown haze apparent under diffuse illumination. Pulling apart the bundles of ECFC fibers clearly reveals their inner core and outer shell structure (Figure 4d). Following the severe oxidative step, we use electrochemical reduction to remove the OCFGs and regain the electrical conductivity of the ≈ 3–4 μm thick oxidized/exfoliated shell of ECFGs. Previous work has shown that electrochemical reduction of oxidized materials (i.e. graphene oxide and oxygenated carbons) can selectively remove OCFGs and reestablish the electrical conductivity of the material.38, 50-51 Electrochemical reduction of ECFCs was carried out potentiostatically at −1.2 V vs Ag/AgCl for 30 min in a 4 M KOH solution. After reduction, the translucent core-shell structure could no longer be seen under an optical microscope, likewise, the red-brown color for ECFCs turned to matte black for rECFCs. SEM images of the ECFCs and pristine CFC are shown in Figure 5. The surface of pristine CFCs has a smooth appearance with a cylinder like morphology comprised of axially oriented carbon filaments (Figure 5a and b). The significant structural changes of ECFCs observed with optical microscopy (Figure 4) are more clearly seen in SEM images (Figure 5d-f).


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The ECFC surface is roughened compared to CFC, with larger spacing between the individual carbon filaments of the fiber (Figure 5d). Images of the cross section of the fibers clearly show a thick shell (Figure 5d-f) surrounding a compact carbon fiber core. These structural changes are considerably different from previous work that employed electrochemical oxidation with nonoxidizing electrolytes. The most obvious difference is the formation of the several micrometerthick shells compared to 50–100 nm thicknesses reported by other electrochemical routes. After reduction of the ECFCs, SEM images show that the rECFCs have intermediate diameters and that the separation between filaments is partially retained (Figure 5g and h). FT-IR spectra (Figure S11) of ECFCs reveal the presence of alcoholic (O-H ≈ 3200 cm−1), carbonyl (C=O ≈ 1730 cm−1), and epoxy (C-O ≈ 1030 cm−1) oxygen-containing functional groups. Following electrochemical reduction, rECFCs show an almost complete removal of the broad O-H and sharp C-O peaks, as well as a significant reduction in the carbonyl peak, while the aromatic C=C stretch at ≈ 1600 cm−1 is retained from ECFC. These changes evidence a significant removal of oxygen containing functional groups after reduction. The presence of some OCFGs after reduction can increase the hydrophilicity of the exfoliated graphitic layers and prevent their restacking; allowing penetration of ions and solvent molecules between layers.52 This is potentially beneficial for improving the specific capacitance of the EDLC-type charge-storage mechanism of the rECFC by incorporating Faradaic materials having high specific capacitance (i.e. redox polymers, metal oxides, and metal sulfides) into the shell structure. Raman spectroscopy was employed to better understand the chemical and structural composition of the CFs following the two-step oxidation-reduction procedure (Figure S12). A Lorentzian fitting method 53 of the D (1350 cm−1) and G (1620 cm−1) bands (Figure S12) was 19 ACS Paragon Plus Environment


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used to determine the relative composition of sp2 and defective sp3 domains. ECFCs show a characteristic broadening of the G band, corresponding to a disruption in the graphitic structure. The ID/IG ratio for ECFC is 2.12, which is considerably larger compared the pristine CFC ratio of 1.32. This confirms that electrochemical oxidation has disrupted the graphitic domains; presumably due to the introduction of OCFGs into the structure of ECFCs. The electrochemically reduced rECFCs show that the ID/IG ratio of 1.76 is lower compared to ECFCs; suggesting graphitic sp2 domains have returned to some extent. However, the ratio is higher than pristine CFC indicating an increased number of edge defects and partial retention of OCFGs. 3.2.2. Polymerization of Pyrrole on rECFC Electrodes A PPy@rECFC electrode was prepared by oxidative electropolymerization of pyrrole on rECFC. The electropolymerization used pulses stepping between 0.00 and 0.90 V, both of 100 ms duration. Using pulsed electrodeposition of PPy allows electrolyte (e.g. p-TSA) and monomer to diffuse into the porous shell during the “off” potential (0.00 V), which results in coating the entire shell structure and avoids deposition of thick PPy layers only on the outer surface. Optimal results were obtained using a total of 15 min of pulsed electrodeposition using p-toluene sulfonic acid (p-TSA) as the “permanent” dopant. The FT-IR spectra (Figure S13) of PPy@rECFC is consistent with the presence of PPy. The peaks at ≈ 1545 cm−1 and ≈ 1463 cm−1 match the anti-symmetric and symmetric ringstretching modes of PPy, respectively.54-55 The peaks at ≈ 1298 cm−1 and ≈ 1175 cm−1 are associated with the stretching vibration of C–N of PPy while those at ≈ 1040 cm−1 and ≈ 907 cm−1 are from the =C−H bending vibrations of both PPy and rECFC.54-55 The peak at ≈ 1021 cm−1 is commonly observed for polymers containing sulfonate dopants, and in PPy@rECFC this peak is associated with interaction between PPy polarons and the sulfonate group of p-TSA.56 20 ACS Paragon Plus Environment

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Overall, the analysis of PPy@rECFC’s FT-IR spectrum indicates that electrochemical polymerization has successfully produced p-TSA doped PPy at the rECFC electrode. SEM images of PPy@rECFC (Figure 6a and b) show that the horizontal axis of the fibers appears swollen with an increased interlayer spacing between carbon filaments bundles and a deformation of the cylindrical carbon fiber shape. The images also reveal distinct morphological changes after polymerization. The well aligned filament bundles in rECFC (Figure 5g) have become non-uniform and swollen in appearance with larger interlayer spacing (Figure 5i and 6a and b) for PPy@rECFC. This change likely indicates successful embedment of PPy within the rECFC shell. To understand how the structure of rECFC improves the incorporation of PPy, we prepared samples of PPy on pristine carbon fiber cloth electrodes (PPy@CFC) under the same experimental conditions. SEM images of PPy@CFC (Figure 5c) reveal a thin conformal film coating of PPy onto the CF. The image shows large regions where PPy appears to have flaked off the CF surface (compare the false-colored CF (orange) and PPy (grey) regions), which could indicate poor adhesion of PPy to the CF surface. Note also that the cylindrical shape of individual CFs remains unchanged after polymerization to form PPy@CFC electrodes. SEMEDX analysis of PPy@rECFC show extensive infiltration of the PPy into the fiber structure. After completely charging the PPy@rECFC electrode in 4 M KOH, it was cut with a razor blade to expose a cross section. SEM elemental (C, S, and K) mapping images of the PPy@rECFC cross section are shown in (Figure 6). The distribution of carbon (Figure 6f) reveals a dense core (≈ 4 μm diam) surrounded by a lower density outer-layer shell (≈ 4–5 μm diam), which matches well with the core-shell structure observed in optical microscopy (Figure 4) and SEM images (Figure 5 d-f). Sulfur and potassium are found to be present in the PPy@rECFC shell. This is consistent with PPy incorporation into the rECFC shell (Figure 6g) since p-TSA (sulfur



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containing) is known to reside within the polymerized PPy matrix. After treatment in alkali, potassium ions are intercalated into PPy to charge balance the free p-TSA exchanged by hydroxide ions (see eq 14). Figure 6h indicates that the outer shell of PPy@rECFC is freely accessible to dopant and electrolyte ions (Figure 6h and i). An XPS survey spectrum of PPy@rECFC shows well-defined signals of S 2p (180–190 eV) and N 1s (397–404 eV), indicating the presence of p-TSA doped PPy in the composite electrode, and O 1s (530–540 eV) from OCFGs on rECFC (Figure S14a). The N 1s core level spectrum (Figure S14b) has been deconvoluted into three Gaussian peaks assigned to benzenoid amine (–NH–) at 399.8 eV and protonated and nonprotonated nitrogen cations (–N+H– and – N+=) at 400.8 eV and 402.2 eV, respectively.57 A peak at lower binding energies (< 398 eV) for the quinonoid imine (=N–) is almost non-existent for our PPy@rECFC sample, indicating few defects in the electropolymerized PPy chains.58 The calculated doping level, determined as the ratio of the combined area (fitted peaks) of the polaron signals (–N+H– and –N+= peaks) to the total nitrogen signal (–N+H–, –N+=, and –NH–), was 33.2 % for our material, which is at the theoretical maximum for polypyrrole (i.e. 33 % doping).23 The high doping level of our material could be due to the p-TSA permanent dopant, which can act to permanently charge balance the polarons in PPy, and to an additional doping contribution from the negatively charged OCFGs in the exfoliated shell of the rECFC.55, 57, 59 3.3. Electrochemical Characterization of Positive and Negative Electrode The electrochemical performance of the NiO/NiOOH@NF electrodes was evaluated by cyclic voltammetry (CV) in a three-electrode cell with 4 M KOH aqueous electrolyte. CV curves (Figure 7a) over a range of scan rates show a pair of redox peaks, which are assigned to the well-


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accepted Faradaic reaction between NiO (Ni2+) and NiOOH (Ni3+) via the insertion/extraction of OH─ in alkaline electrolyte, according to eq 12: NiOOH + e– ⇌ NiO + OH–

(12)

Figure 7b depicts the XRD pattern of the NiO/NiOOH@NF electrode after a complete discharge. The XRD pattern is well indexed with the standard peaks for NiO (JCPDS card JCPDF: 06-0075), indicating NiO is the main discharge product and confirming the redox reaction in eq 12. Figure 7c shows CVs for the Zr-NF-200, Zr-NF-250, Zr-NF-300 samples compared to bare NF. The CVs show that the Zr-NF-250 sample has superior electrochemical redox performance compared to Zr-NF-200, Zr-NF-300, and bare NF. This result is correlated to the different morphologies of these materials observed in SEM images (Figure 2 and S9), where the high surface area of Zr-NF-250 electrode provides enhanced charge storage capacity. Galvanostatic charge−discharge measurements were carried out in a three-electrode cell at current densities from 5 to 40 mA cm−2 (Figure 7d) and 1 mA cm−2 (Figure S15a). A calculated areal specific capacity of 0.613, 0.525, 0.431, 0.378, 0.347, and 0.326 mAh cm–2 was found at current densities of 1, 5, 10, 20, 30, and 40 mA cm−2, respectively. Consistent with the CV results (Figure 7a), the plateaus in the charge/discharge curves indicate the existence of battery-type Faradaic processes. In addition, the time-symmetrical charge/discharge curves indicate that Zr-NF-250 has high Coulombic efficiency (100 %) because of the superior reversibility of redox reactions at the NiO/NiOOH nanosheets. Furthermore, when the Zr-NF250 electrode is subjected to a high charge/discharge current density of 40 mA cm−2, 53 % of the charge storage capacity (0.326 mAh cm–2) is retained compared to 1 mA cm−2 (0.613 mAh cm–2), which is an impressive rate capability compared to traditional battery-type storage materials.



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The excellent electrochemical energy storage performance of Zr-NF-250 can be attributed to several factors: First, the microporous separation between individual nanosheets (see Figure 2) can act as an ion-buffering reservoir to sustain concentrations of OH− at the surface, and the mesoporous cavities (see Figure 3) allows electrolyte ions to diffuse three dimensionally throughout the entirety of the stacked nanosheets. Furthermore, the polycrystalline nature of the nanosheets (see Figure 3), which are composed of numerous nano-sized (≈ 4–5 nm) cubic grains, provides numerous exposed edges, surface sites, and mesoporous cavities that maximize utilization of the material. These unique structural features enable rapid electrolyte intercalation within the bulk of NiO/NiOOH nanoflakes. In addition, the good intrinsic electrical conductivity and adhesion of NiO/NiOOH grown directly from the NF substrate improves electrical contact and charge-transfer properties at the interface of the active material and current collector. The electrochemical properties of the PPy@rECFC employed as a negative electrode was investigated via cyclic voltammetry in a three-electrode cell with 4 M KOH aqueous electrolyte (Figure 8a). Initially, when PPy@rECFC is placed in alkali solution, p-TSA dopant ions are rapidly exchanged with hydroxide ion dopants (potassium ions are also inserted to charge compensate free p-TSA) via an anion exchange mechanism (eq 13).60-63 The reduction and oxidation peaks in the CV are assigned to the Faradic redox reaction of PPy accompanied by the insertion/expulsion of OH‒ ions according to the mechanism in eq 14:63-66 PPy+(p-TSA─) + K+ + OH─ → PPy+(OH─) + (p-TSA─ )K+

(13)

PPy+(OH─) + e─ ⇌ PPy0 + OH─

(14)

Figure 8b shows the CV of rECFC, PPy@rECFC, and PPy@CFC at a scan rate of 10 mV s−1. Integration of the CV curves shows that the total amount of charge stored in PPy@rECFC is


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3.4 and 9.4 times higher than that of PPy@CFC and rECFC, respectively. Furthermore, the mass loadings of PPy onto PPy@CFC and PPy@rECFC were both determined to be the same (i.e. 1.5 mg). The combined charge stored in individual rECFC and PPy@CFC electrodes is only 74 % of that stored in the PPy@rECFC electrode. The analysis of CV curves for rECFC and PPy@rECFC electrodes shows that the charge storage contribution of EDLC (originated from rECFC) and pseudocapacitive-type (Faradaic) (originated from grafted PPy) are 24% and 76%, respectively. This enhanced charge storage capability of PPy@rECFC demonstrates the synergetic effect of the core-shell architecture between PPy and the rECFC materials. This result clearly shows that redox reactions of grafted PPy adds an impressive Faradaic contribution to the double layer capacitance of rECFC and subsequently increases the total capacity of the negative electrode for a full-device fabrication. The charge storage performance of the PPy@rECFC electrode was quantified with galvanostatic charge−discharge measurements. Representative plots of the charge−discharge curves in the potential window of –1.0 to –0.2 V vs Ag/AgCl at current densities from 5 to 40 mA cm−2 (Figure 8c) and 1 mA cm−2 (Figure S15b) show a distorted linear shape, which indicates the presence of both pseudocapacitive-type (Faradaic) and EDLC-type (a linear dependence of the charge with potential) charge storage contributions in hybrid materials. This is in good agreement with the CV curves that show a pair of redox peaks on semi-rectangular shape CV. The areal specific capacity during the galvanostatic charge−discharge process was calculated (eq 1) and found to be 0.594, 0.560, 0.547, 0.507, 0.480, and 0.467 mAh cm–2 at discharge current densities of 1, 5, 10, 20, 30, and 40 mA cm–2, respectively. The material also shows an excellent capacity retention of 94.2, 92.0, 85.3, 80.8, and 78.5 % at current densities of 5, 10, 20, 30, and 40 mA cm–2, respectively, when compared with the 1 mA cm–2 results. The



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excellent electrochemical performance of the PPy@rECFC electrode can be attributed the structural characteristics of the core-shell architecture. The exfoliated shell of rECFCs contains long carbon filaments along the fiber axes with enlarged spacing that provides easy access for ions (e.g. p-TSA) and monomers, facilitating incorporation of PPy into the outer shell during electrochemical polymerization. These long filaments also provide an extended electron transport network for PPy after its incorporation, enhancing the excellent rate capability of PPy@rECFC. Note that the EDLC electrochemical response of rECFC in the region of −0.8 and −1.0 V (see Figure 8b) has been retained in PPy@rECFC electrodes. Since PPy does not contribute any charge in this potential window (compare to the PPy@CFC CV in Figure 8b), a conclusion is that ions have ready access to the surface of rECFC. This indicates that embedding PPy into the shell does not inhibit the EDLC response, likely due to the porous structure and hydrophilicity of PPy. Further, it has been shown that OCFGs possess the ability to act as permanent dopants and nucleation sites for conductive polymers leading to enhanced rate performance capabilities.55, 67-68 The combined contribution of OCFG dopants likely facilitates a permanently doped PPy state and retaining the conductivity of doped PPy.23, 55 Figure 8d shows Nyquist plots of positive and negative electrodes in a three-electrode system. A DC bias potential was set at the midpoint of the charging plateau during the EIS scan– 0.23 and –0.55 V for the positive and negative material, respectively. The real axis (Z') intercept at high frequency is modelled as a combined resistance (Rs), which includes the ionic resistance of electrolyte, intrinsic resistance of active material, and contact resistance at the interface of active material/current collector. These data were fit to equivalent circuits to determine Rs and Rct values (see Figure S16). The Rs values for the positive and negative electrodes were determined to be 0.79 and 1.94 Ω, respectively. The different Rs values for the negative and positive


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electrode are mostly due to the higher resistance of carbon fiber (negative substrate) compared to the nickel foam (positive substrate). The diameter of the semicircle observed at high frequency (lower left) is related to the charge-transfer resistance (Rct) at the electrode/electrolyte interface. The small Rct values for the negative electrode (0.42 Ω) indicates large charge-transfer rates at the active material-electrolyte interface. However, the much smaller semicircle (Rct = 0.06 Ω) observed for the positive electrode (Figure 8d, inset) suggests that much more rapid chargetransfer is taking place at the NiO/NiOOH nanosheet interface. 3.4. Charge Storage Contribution The charge storage contribution from the outer surface and the inner bulk of the active materials, at both positive and negative electrodes, was examined by a method initially proposed by Trasatti et al.69-70 The total charge (QT) is separated into two components: QT = Qouter + Qinner

(15)

QT is the total charge stored during the half oxidative or reductive cycle and is calculated from cyclic voltammograms based on the equation QT=∫I dV/2v; where I is the current density (A cm−2), ν is the potential scan rate (V/s), and V is the applied potential (in volts, V). Qouter is the charge contribution from surfaces that are freely accessible to electrolyte and thus yields a capacitive response through physical adsorption or surface redox reactions, and Qinner is the charge contribution from pores, grain boundaries, crevices, and cracks and/or insertion/intercalation-based processes in the bulk of the material. If charge storage to the inner region is diffusion limited, Qinner can be replaced by a function of scan rate (v) and the charge is analyzed according to the following equation (assuming planar diffusion):69-70 Q(v) = Qouter + constant v−1/2

(16)



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Based on eq 16, Qouter can be obtained from the plot of Q(v) versus v−1/2 by extrapolating the voltammetric charge, Q(v), to v = ∞ (Figure S17a and b). As v → ∞ (i.e. v−1/2 → 0) ions in the electrolyte solution have insufficient time to diffuse into the bulk of the material, and Qinner has no charge storage contribution. In addition, plotting 1/Q(v) versus v1/2 and extrapolating to v = 0, results in a linear portion with a y-intercept equal to the 1/QT (Figure S17a and b). At low scan rates (i.e. v → 0), diffusion-limited processes contribute, and the total calculated charge is both Qinner and Qouter.71 Consequently, Qouter is found by subtracting the Qinner from QT according to eq 15. As shown in Figure S17a and b, NiO/NiOOH@NF has a QT, Qinner, and Qouter calculated to be 0.50, 0.29, and 0.21 mAh cm−2. These results indicate that more than half of the total charge storage (58 %) of the positive electrode comes from the inner bulk of NiO/NiOOH nanosheets compared to the outer surface (42 %). The large surface charge contribution can be attributed to the high surface area of the interconnected NiO/NiOOH nanosheets. Charge storage from the inner bulk of NiO/NiOOH is likely due to fast intercalation of OH− (diameter of 0.274 nm) into the crystal lattice, enabled by the large d-spacing values of 0.69 (003) and 0.343 (006) nm of the hexagonal crystal structure of γ-NiOOH. Moreover, as confirmed by TEM imaging (Figure 3), the ultrathin and mesoporous structure of the NiO/NiOOH nanosheets can provide ion reservoirs inside the nanosheets themselves, which can shorten the diffusion path upon intercalation/expulsion. The charge storage contribution of the PPy@rECFC negative electrode was calculated from the plots shown in Figure S18a and b and determined to be QT = 0.824, Qinner = 0.456, and Qouter = 0.368 mAh cm–2. This indicates that the more than half of the total charge comes from the inner bulk (55 %) of PPy@rECFC electrode. This charge contribution of PPy@rECFC is likely due to diffusion of hydroxide ions into the bulk of PPy, which is characteristic for such


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polymer based pseudocapacitive materials. The high outer surface charge contribution (45 %) further validates our claim that the shell has retained an “open” framework, which allows rapid diffusion of electrolyte ions into a micrometer thick shell after PPy embedment. As shown in the SEM-EDX maps of PPy@rECFC (Figure 6 h and i), ions are located throughout the entire micrometer thick shell of the PPy@rECFC core-shell structure, which agrees well with the charge storage contribution analysis. 3.5. Electrochemical Characterization of NiO/NiOOH || PPy@rECFC Full-Cell Device The complementary potential windows and good electrochemical performance of the NiO/NiOOH and PPy@rECFC electrodes suggest that an asymmetric supercapacitor (ASC) composed of these electrodes would show superior behavior. To this end, a two-electrode system using NiO/NiOOH@NF nanosheets as the positive electrode and PPy@rECFC as the negative electrode was constructed and investigated within a potential window of 0.0 to 1.4 V in 4 M KOH aqueous electrolyte. To maximize device performance, the anode and cathode area was initially adjusted to provide a charge balance of Q− = Q+ according to eq 2. However, the ASC with equal charge balance had poor cycling stability, only 25 % of its initial capacitance was retained after 2000 charge/discharge cycles (Figure 9a). To gain insight into the capacity loss observed for the cell, we disassembled the cycled NiO/NiOOH||PPy@rECFC cell and conducted CV scans on both the positive and negative electrodes in a three-electrode configuration (Figure 9b and c). After 5000 cycles, the electroactivity of the NiO/NiOOH electrode decreased slightly, however, that of the PPy@rECFC electrode reduced significantly. Therefore, the capacity decay seen after cycling the full-cell device is primarily caused by the negative electrode instability. Figure S15d shows the capacitance retention of positive and negative electrodes after 1000 cycles in a three29 ACS Paragon Plus Environment


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electrode cell configuration. This result shows that the performance decay during cycling is faster at the PPy@rECFC than the NiO/NiOOH@NF electrode. Therefore, when the potential window is fixed to 1.4 V, consecutively lower quantities of charge can be stored in the negative electrode compared to the positive electrode after each charge/discharge cycle of the ASC. As a result, the upper potential limit of PPy@rECFC during the discharge process shifts to more positive potentials in an attempt to match the charge of the positive electrode. As seen in Figure 9d, when a newly prepared PPy@rECFC electrode is cycled within a potential window more positive than −0.2 V, overoxidation of PPy leads to an electrochemically irreversible product.59,72 To overcome this mismatch in decay for NiO/NiOOH@NF and PPy@rECFC, a galvanic chargedischarge experiment at a current density of 30 mA cm−2 was performed on a full-cell device with an area ratio of 1:1 for the negative and positive electrode, corresponding to a charge ratio of Q−/Q+=1.34. Figure 9a shows that 47 % and 96 % of the capacitance is retained after 2000 cycles for the NiO/NiOOH@NF||PPy@rECFC full-cell device with Q−/Q+ ratios of 1 and 1.34, respectively. After 5000 cycles, the full-cell device with Q−/Q+=1.34 retains 88% of its initial capacity. Thus, increasing the negative electrode charge storage prevents the rapid cycling decay while simultaneously balancing the areal charge. Figure 10a exhibits CV curves for the full-cell device in 4 M KOH aqueous electrolyte at a range of scan rates. As shown in Figure 10b, CC curves consist of two charge and discharge voltage semi-plateaus corresponding to the pair of peaks observed in the CV curves. As discussed in section 3.3, the positive electrode based on NiO/NiOOH@NF has a battery-type charge storage mechanism with a flat discharge plateau during charge/discharge tests under constant current (see Figure 7d). The PPy@rECFC negative electrode exhibits a partially distorted linear CC curve which is from the combined pseudocapacitive and EDLC charge


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storage mechanisms of PPy and rECFC, respectively. CC curves of the full cell device indicate that the charge storage is based on both EDLC-type and Faradaic-type behavior. The Faradic behavior originates from the battery-type and pseudocapacitive-type reversible redox couples of NiO/NiOOH and PPy+/PPy while the EDLC contribution comes mainly from rECFC. Considering the complementary electrochemical redox regions of the positive and negative electrode materials, a general charge/discharge mechanism can be described for the full asymmetric device. Figure 10c, illustrates this concept by superimposing the three-electrode cyclic voltammograms of the positive NiO/NiOOH@NF and negative PPy@rECFC electrodes. During a charging cycle, the positive electrode material is oxidized per eq 12 (section 3.3) while the negative electrode material is reduced per eq 14 (section 3.3). Discharge of the full cell reoxidizes the negative electrode and re-reduces the positive electrode. Fitting the Nyquist plot (Figure 8d and S16) gives a small charge-transfer resistance (Rct = 0.51 Ω) for the full-cell device, indicating that the rapid charge-transfer performance of the individual positive (0.06 Ω) and negative (0.42 Ω) electrodes are still maintained and well matched for a high-performance ASC supercapacitor fabrication. The volumetric specific capacity (Cs,cell) based on the total volume of the full-cell device (0.09 cm3) is calculated from the CC curves (Figures 10b and S15c) according to eq 1. The Cs,cell values of the fabricated cell are calculated to be 4.99, 4.57, 4.54, 3.98, 3.66, 3.58, and 3.55 mAh cm–3 at current densities 1, 5, 10, 20, 30, 40, and 50 mA cm–2, respectively. The corresponding volumetric specific capacitances at these current densities were calculated from eq 5 and found to be 2.10, 1.89, 1.76, 1.55, 1.43, 1.38 and 1.33 F cm–3, respectively. The device demonstrates an excellent 71 % retention of the capacity at a current density of 50 mA cm–2 compared to the capacity at 1 mA cm–2, which demonstrates the full-cell device is capable of high-rate energy



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storage applications. The observed decrease in specific capacity at a high discharge current density may be explained by limits in diffusion and intercalation of hydroxide ions within the electrode inner bulk. Additionally, we find that the full-cell device maintains a good electrochemical reversibility, having 98 % Coulombic efficiency after 5000 cycles at the high current density of 30 mA cm−2 (Figure 10d). The SEM images of negative (Figure S19 a and b) and positive (Figure S19 c and d) electrodes after 5000 cycles demonstrate that the nanosheets and nanoflakes structures of positive electrode as well as the core-shell structure of negative electrode are retained after prolonged cycling. However, a lower density of nanosheets is observed on the positive electrode which might be due to the loss of active materials via the partial water splitting at higher potentials. Figure 11 (a-i) demonstrates the assembled prototype cell powering LEDs. The cells consisted of 1 cm2 positive and negative electrodes, a 1.5 cm2 piece of KOH-soaked filter paper, and two nickel foil current collectors placed inside of a plastic pouch cell (Figure 11a) and were charged at 10 mA cm‒2 prior to the demonstrations. For sufficient voltage to power a single red LED (5 mm; Vf = 1.9 V), two ASC cells were placed in series (Figure 11 d-e). Figure 11c shows these two cells can light a red LED over a 2 h period. Additionally, we show that the same two cells in series can power 17 paralleled LEDs (5 mm; Vf = 1.9-2.0 V) over a 5 min period (Figure 11 g-i). These experiments illustrate the high energy density of our materials and the full device. Considering the superior volumetric energy densities of our device, these prototype devices demonstrate the promise of our modification of current collector approach toward enabling high performance ACSs devices. A way to provide a realistic picture of high power and high energy density exhibited with NiO/NiOOH||PPy@rECFC full-cell device is to plot our results in a Ragone plot (Figure 12).


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This plot demonstrates the relationship between the specific power and specific energy for our full-cell device with recently reported literature devices as well as several commercially available energy storage devices. Our full-cell device can achieve an excellent specific energy of 4.12 mWh cm−3 at the high specific power of 9.17 mW cm−3 and impressive retention of 2.60 mWh cm−3 when the power is increased to 425.58 mW cm−3. This is a better high energy/power performance than many recent reported values in the literature, such as Co3(PO4)2·8H2O//ACSupercapattery (1.17 mWh cm−3 at 18.74 mW cm−3),73 CF-Co(OH)2//CF-WO3@PPy-ASC (1.02 mWh cm−3 at 45 mW cm−3),19 Ni/GF@MnO2//Ni/GF@PPy-ASC (1.23 mWh cm−3 at 4.5 mW cm−3),74 NF@FeOOH//NF@Co–Ni-DH-ASC (0.723 mWh cm−3 at 15.3 mW cm−3),75 CuS NWs//AC-ASC (1.11 mWh cm−3 at 37 mW cm−3),76 Ni/GF/H-CoMoO4//Ni/GF/H-Fe2O3 (1.13 mWh cm−3 at 39 mW cm−3),77 Ni/GF/MnO2//Ni/GF/PPy-ASC (1.23 mWh cm−3 at 4.5 mW cm−3),78 MnO2//WON-ASC (1.27 mWh cm−3 at 650 mW cm−3),79 MnO2//Ti-Fe2O3@PEDOTASC (0.89 mWh cm−3 at 380 mW cm−3),80 and PPy/CPH//PPy/CPH-ASSPS (1.161 mWh cm−3 at 4.0 mW cm−3).81 Additionally, our device performs considerably better than a commercially available supercapacitor (5.5 V/100 mF, 0.55 mWh cm−3) and is well positioned to bridge the gap to high energy density (≈ 8 mWh cm−3 at 1.5 mW cm−3) commercial thin-film lithium-ion batteries (4 V/500 μAh), which have low power densities (0.4 mWh cm−3 at 5 mW cm−3).51, 82 In addition, we have reported the gravimetric energy and power density for the full-cell device which comprises of the mass of positive (28.1 mg cm–2) and negative (24.2 mg cm–2) electrodes (i.e. total electrode mass), separator (3.5 mg cm–2), and KOH electrolyte (80 mg: used in device fabrication). From this perspective, the specific energy of the full-cell device as constructed is 1.73 Wh kg−1 at a specific power of 284 W kg−1 and 2.75 Wh kg−1 at 6.1 W kg−1.



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4. Conclusion In contrast to the considerable efforts being made towards bulk synthesis of new energy storage materials, we focused on the modification of commercially available current collectors: nickel foam and carbon cloth, which themselves can participate as the active charge storage materials. This is of great importance from the industrial/commercial point of view since the current collectors can have a significant contribution to the total volume and mass of the full-cell devices, a fact that is commonly overlooked in most publications in the field. A soft templating/solvothermal treatment route was employed to generate NiO/NiOOH nanosheets from the Ni of the NF current collectors to produce the positive electrode. An essential step is the use of zirconium hydroxide as a time/temperature based dynamic-softtemplate. It was found that the shrinking and the redox mediation of the zirconium hydroxide gel generates the chemical makeup and unique morphology of the nickel oxide/hydroxide materials during the synthesis. For the negative electrode, a carbon fiber cloth was modified through an electrochemical oxidation/reduction route to generate a reduced exfoliated core-shell structure (rECFC) followed by electropolymerization of pyrrole into the shell structure to produce the active PPy@rECFC material. When examined in a three-electrode configuration, the NiO/NiOOH@NF electrodes show typical Faradaic redox properties consistent with battery-type materials, while PPy@rECFC electrodes store charge though both pseudocapacitive and EDLC mechanisms. Analysis of charge storage contribution at each electrode indicated that in both electrodes more than half of the total charge storage capacity comes from the inner bulk, indicating the high accessibility and usability of these materials for charge storage. When these materials are paired to produce an asymmetric supercapacitor, the full-cell device delivers promising volumetric energy densities in the range of 2.60 to 4.12 mWh cm−3 and power densities in the range of 9.17 to 425.58 mW cm−3. Such performance is comparable to 34 ACS Paragon Plus Environment

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lithium thin films (0.3-10 mWh cm−3) and better than some commercial supercapacitors (< 1 mWh cm−3). These values are even more impressive since the entire cell volume (active materials, separator, and both current collectors), as opposed to the active material volume alone, has been considered. In addition, it was found that by increasing the charge ratio (Q−/Q+) of the full-cell device from 1 to 1.34, the overoxidation of PPy is suppressed and the cyclability performance of the full-cell asymmetric supercapacitor is significantly improved. The cycling performance of the full-cell device demonstrates excellent stability with capacitance retentions of 96 % and 87 % after 2000 and 5000 cycles, respectively. Additionally, the device achieved an excellent 71 % retention of the capacity at high current densities of 50 mA cm–2 compared to the capacity at 1 mA cm–2, which demonstrates the capability of the full-cell device for high-rate energy storage applications. Overall, our report demonstrates the utility of directly modifying the current collectors to form active material and provides a new approach for fabricating high performance asymmetric supercapacitor devices.

ASSOCIATED CONTENT Supporting Information Schematic diagram of the full-cell device; synthesis scheme of the nickel foam-supported NiO/NiOOH nanosheets; SEM images of bare NF and zirconia-coated NF; EDX line scan of the Zr-NF-250 sample; XPS spectra of NiO/NiOOH nanosheets and Zr-NF-250 sample; SEM images of Zr-NF-250 sample heat treated for 30 min and 60 min; SEM images of pristine nickel foam heat treated in KOH at 250 °C; SEM images of zirconia-coated-NF samples subjected to heat treatment in KOH at 200 °C and 300 °C; synthesis scheme for the PPy@rECFC core-shell electrode; FT-IR and Raman spectra of CFC, ECFC, and rECFC samples; FT-IR spectra of



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PPy@rECFC sample; XPS spectrum of PPy@rECFC sample; galvanostatic charge-discharge curve for the positive electrode, negative electrode, and full-cell device at a current density of 1 mA cm−2; the capacitance decay comparison of positive and negative electrodes after 1000 cycles; Nyquist plots with corresponding equivalent circuits for the positive electrode, negative electrode, and full-cell device; dependence of 1/C on v1/2 and dependence of C on v−1/2 for the positive and negative electrodes; SEM images of positive and negative electrodes after 5000 cycles.

AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] Notes: The authors declare no competing financial interest

ACKNOWLEDGMENTS The authors acknowledge the Department of Chemistry at Mississippi State University for the financial support of this work.


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Figure 1. XRD pattern and the standard JCPDS cards for (a) zirconium hydroxide-coated NF, and (b) the Zr-NF-250 sample. Peaks labeled by asterisk represent the nickel foam XRD peaks.


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Figure 2. Low- and high-magnification SEM images (a-i) of a Zr-NF-250 sample composed of rose-like (i) and flake-like (g and h) domains.



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Figure 3. (a, b, and c) HRTEM images of NiO/NiOOH nanosheets from Zr-NF-250 sample, and (d) (upper two images) rose-like NiOOH nanosheets and (lower two images) flake-like NiO nanosheets.


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Figure 4. Optical microscopy of exfoliated core-shell carbon fiber cloth (ECFC). (a-d) show translucent red-brown graphitic-hydrogel shells around the carbon fiber cores, (a, inset) magnified region for an individual exfoliated CF in (a), (b) shows the end of an individual CF having total loss of the carbon fiber core, (b, inset) illustration of geometric influence of CFs on the exfoliation process, (c) shows CF tips (without core) and shafts (with core), and (d) shows CFs from mechanically broken ECFC electrodes revealing their core-shell structure.



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Figure 5. SEM images of (a and b) pristine CFC, (c) PPy@CFC, (d) ECFC revealing the coreshell structure, (e and f) magnified cross-sections of ECFC fiber, (g and h) rECFC, and (i) PPy@rECFC.


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Figure 6. SEM and EDX images of PPy@rECFC, (a) side view of PPy@rECFC fiber, (b) magnified view of (a) with false coloration, (c, d, and e) cross section of PPy@rECFC, (f) carbon EDX of PPy@rECFC cross section, (g) sulfur EDX of PPy@rECFC cross section, (h) potassium EDX of PPy@rECFC cross section, (i) potassium EDX of PPy@rECFC cross section overlay with (e). 43 ACS Paragon Plus Environment


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Figure 9. (a) Cycling performance of the full-cell devices with Q−/Q+ = 1 and 1.34, (b) CV curves of the as-prepared positive electrode and detached positive electrodes operated for 5000 cycles in a full cell device with Q−/Q+ = 1 and 1.34, (c) CV curves of the newly prepared negative electrode and disassembled negative electrodes after 5000 cycles in a full cell device with Q−/Q+ = 1 and 1.34, and (d) CV curves of the newly prepared PPy@rECFC electrode with different upper potential limits. (CV scan rate = 10 mV s−1, charge-discharge current density = 30 mA cm−2).


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Figure 11. (a, b, and c) Photographic images of the assembled ASCs device, (c) demonstration of two full cells in series powering an LED for 2 h, (d) electrical schematic for the demo, (e) photograph of the demo circuit, (f-i) additional demo of two series cells powering 17 LEDs wired in parallel. For all demos, the cells were fully charged (10 mA cm‒2).


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1e-2 Our Work (NiO/NiOOH@NF//PPy@rECFC)

3

Energy Density (Wh/cm )

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Li Thin- Film Battery

High E Commercial 5.5V/100 mF SC

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CuS NWs//AC-ASC 77 Ni/GF/H-CoMoO4//Ni/GF/H-Fe2O3

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Power Density (W/cm ) Figure 12. Ragone plot of the proposed full-cell device compared with commercially available energy storage devices and most recently reported full cell devices.



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