Structurally confined ultrafine NiO nanoparticles on graphene as a highly efficient and durable electrode material for supercapacitors

Abstract

The most significant challenge in the development of ultrafine oxide based supercapacitors is the poor microstructure stability due to the rapid agglomeration of the fine nanoparticles (NPs). Here, we developed novel amorphous MnO_x structurally confined ultrafine NiO NPs (∼2.3 nm) supported on graphene, NiO@MnO_x via a simple and facile self-assembly process with the assistance of microwave sintering. NiO@MnO_x with a NiO : MnO_x weight ratio of 1 : 0.2 achieves a high capacitance of 966 F g⁻¹ based on total electrode materials and 3222 F g⁻¹ based on active materials at a discharge current density of 2 A g⁻¹. Remarkably, the materials retain 100% capacitance after 2000 cycles at a charge and discharge current of 10 A g⁻¹. In contrast, the durability of ultrafine NiO NPs without MnO_x confinement is very poor, with 94% of the capacitance lost under identical cyclic conditions despite the initial high capacitance of 3696 F g⁻¹. The substantially enhanced capacitance, durability and high rate capacity contribute to the formation of a nanoporous and amorphous MnO_x layer on ultrafine NiO NPs, which provides the extraordinary structural confinement and enhances the mass transfer process. The results provide a new strategy to develop highly efficient and durable ultrafine nanosized electrode materials for supercapacitors.

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1 Introduction

Metal oxides, such as RuO₂,¹ MnO₂,^2,3 NiO,⁴ Co₃O₄,⁵ SnO₂,⁶ TiO₂,⁷ V₂O₅,⁸ which exhibit an order of magnitude higher specific capacitance (C_s) than carbon electrodes and conductive polymers, have been considered to be the most promising materials for developing high-performance supercapacitors.⁹ However, the practical application of metal oxides is hindered by the high charge transfer resistance due to the fact that most of the oxides are semiconductors with a wide band gap and the cracking of transition metal oxide based electrodes results from the strain developed during the charge–discharge processes. One strategy is to combine electronically conductive carbon materials with nanosized metal oxides, where the carbon nanostructures not only serve as the physical support for metal oxides but also provide the channels for charge transport.^2,10–16

Intensive work has been devoted to develop carbon–metal oxides hybrids with the aim to achieve high capacity for energy storage. Different nano-structured and hierarchy structures, such as nanoflowers,^17,18 nanotubes,¹⁹ nanoplates¹⁶ and core–shell structures²⁰ etc., have been developed to increase the efficiency. While conventional metal oxide structures normally exhibit mediocre specific capacitance and show low active material utilization due to the fact that ion insertion/extraction depth is limited in the range of few nanometers on the surface of electrode materials. Two dimensional transition metal oxides (2D TMOs) show promising opportunities due to the improvement in the utilization of active materials. For example, Zhu et al. prepared 2D nickel hydroxide nanosheets with micron-sized planar area and thickness of <2 nm and reported specific capacitance of 3650 F g⁻¹ at 2 A g⁻¹.²¹ Unfortunately, these 2D TMOs are hard to fabricate due to their diverse coordination styles.²²

Ultrafine nano-sized (less than 3 nm) metal oxide NPs are promising electrode materials of supercapacitors because it is relatively easy to control the particle size. But there are few studies in the development of ultrafine nano-sized metal oxide NPs due to the fact that ultrafine NPs inevitably exhibits high surface free energy, high electrical resistance and sluggish charge and proton transfer,²³ which could consequently promote the growth and aggregation of NPs and the degradation of the electrocatalytic activity of the materials. For example, Ni(OH)₂ NPs with average size ∼4.8 nm can only retain ∼59% of the initial capacitance after 1000 cyclic voltammetry cycles at a scan rate of 20 mV s⁻¹.¹⁴ Jahromi et al. found that NiO NPs with an average particle size of 8 nm only remain ∼60.6% of its initial capacitance after 1000 cyclic voltammetry cycles at 2 A g⁻¹.²⁴ Zhao et al.²⁵ studied the capacitive performance of NiO supported graphene oxide with NiO NPs size of 5–7 nm and initial capacitance of 525 F g⁻¹ and showed a high retention of 95.4% of initial capacitance after 1000 cycles but at a relatively low current density of 0.2 A g⁻¹.

The growth and agglomeration of ultrafine NPs could be reduced by the nano- or structure confinement. It has been shown that Pt-based NPs can be stabilized via pore confinement or encapsulation.^26–28 Shang et al. reported a facile and scalable wet-chemical process to prepare graphene-nanosheet-supported Pt NPs covered by mesoporous silica (mSiO₂) layers. The chemically and thermally stable mSiO₂ layer not only prevented the aggregation and restacking of graphene nanosheets resulted from the π–π stacking interactions, but also provided spaces for confining metal NPs.²⁹ Galeano et al. synthesized Pt NPs with size of 3–4 nm encapsulated within the hollow graphitic spheres (HGS) with a specific surface area exceeding 1000 m² g⁻¹ and showed good long-term stability and high activity.²⁶ The increased durability of Pt NPs was attributed to the structural confinement of HGS to suppress the agglomeration.

Here, we developed a new strategy to structurally confine ultrafine NiO NPs of 2.3 nm with an amorphous MnO_x thin layer homogenously supported on graphene sheets, NiO@MnO_x. The best results were obtained on NiO@MnO_x NPs with NiO to MnO_x weight ratio of 1 : 0.2, achieving a high specific capacitance of 3222 F g⁻¹ (based on the active materials), high material utilization, outstanding rate capability and columbic efficiency as well as excellent durability.

2 Experimental

2.1 Materials

The graphene used in present study was obtained from the Graphene Supermarket (1.6 nm flakes, USA). The as-received graphene was purified following the procedure reported:³⁰ graphene (50 mg) was dispersed in 50 mL HCl acid (30 wt%) and was ultrasonicated for 1 h. The dispersion was separated and the sludge was dispersed in a fresh 50 mL HCl acid (30 wt%), followed by being stirred overnight. Then the dispersion was separated using a centrifuge before being washed by DI water for 3 times. The purified graphene was dried and collected for further use. Nitric acid (65%, Fluka), ethanol (Sigma-Aldrich), nickel(II) acetylacetonate (Sigma-Aldrich), potassium permanganate (KMnO₄, Sigma-Aldrich), ethylene glycol (EG, Sigma-Aldrich), Nafion solution (5% in isoproponal and water, Sigma-Aldrich), poly(ethyleneimine) (PEI, molecular weight ∼ 1300, Sigma-Aldrich) were used without further purification.

2.2 Preparation of graphene supported NiO@MnO_x

The purified graphene was functionalized with PEI following the procedure reported:^31,32 50 mg graphene was sonicated in 400 mL Milli-Q water in the presence of 0.5 wt% PEI for 2 h, then the dispersion was stirred for overnight before it was filtrated using a nylon membrane and washed for several times to remove the excess PEI. The as-prepared solid was then dried in vacuum oven for 24 h at 71 °C.

PEI functionalized graphene (30 mg) was first ultrasonicated in 100 mL EG solution for 1 h, followed by the addition of 33.8 mg nickel(II) acetylacetonate. The dispersion was ultrasonicated for 15 min and then stirred for another hour before being placed in a microwave oven (1000 W) and heated for 4 min, followed by stirring overnight. The solution was then filtered using a nylon filter membrane and washed for several times. The designed loading of NiO on graphene was 25 wt%, and the as-prepared catalysts were denoted as NiO.

NiO (10 mg) was dispersed in 40 mL Milli-Q water under ultrasonic for 30 min and appropriate amount of KMnO₄ solution (1 mg mL⁻¹) was added. The dispersion was stirred at room temperature for 30 min followed by adding 40 mL ethanol. The mixture then were refluxed at 100 °C in an oil bath for 2 h and filtered using a nylon filter membrane (0.2 μm) and washed for several times with ethanol. The loading of MnO₂ formed was controlled as 5, 10 and 25 wt% on graphene. This resulted in the synthesized NiO@MnO_x with NiO to MnO_x weight ratio of 1 : 0.2, 1 : 0.4 and 1 : 1 supported on graphene, which was denoted as NiO@MnO_x (1 : 0.2), NiO@MnO_x (1 : 0.4) and NiO@MnO_x (1 : 1), respectively. For comparison, 25 wt% MnO_x on PEI-functionalized graphene was prepared in the same way and the product was denoted as MnO_x.

2.3 Characterization

Thermogravimetric (Q5000, USA) analysis (TGA) was performed to investigate the loading of NiO and MnO_x under air flow upon equilibration at 100 °C for 15 min, followed by a ramp of 10 °C min⁻¹ up to 800 °C. The X-ray photoelectron spectroscopy (XPS) measurements were carried out on a XPS apparatus (Kratos AXIS Ultra DLD) with pass energy of 40 eV, the XPS spectra are referenced to the C 1s feature of the substrates at a binding energy (BE) of 284.5 eV. Microstructure of NiO and NiO@MnO_x on graphene was studied using a transmission electron microscope (TEM, JEOL JEM-2000EX) operating at 80 kV and a high angle annular dark field STEM (HAADF-STEM) and EDS elemental mapping on a FEI Titan G2 80-200 with ChemiSTEM Technology at 80 kV. The average sizes of the NPs were estimated by measuring 200 randomly chosen NPs on the TEM or STEM images. The BET surface area is measured using a Gemini 2360 surface analyzer.

NiO, MnO_x or NiO@MnO_x on graphene (10 mg) was ultrasonically mixed in 2 mL of Nafion solution to form a homogeneous ink, followed by pasting 100 μL of catalyst ink onto the surface of a carbon electrode with diameter of 1 cm⁻². The total electrode material loading was 0.5 mg cm⁻². Pt foil (3.0 cm²) and saturated camel electrode (SCE) electrodes were used as the counter and reference electrodes, respectively. All potentials in the present study were given versus SCE reference electrode and all the experiment were performed in a 2 M KOH aqueous solution. Cyclic voltammetry (CV) and galvanostatic charge–discharge cyclic tests were conducted with different scan rates and current densities, using a Princeton potentiostat (Versastat 3, USA) in a standard electrochemical cell. Electrochemical impedance spectroscopy (EIS) measurements were conducted in the frequency range from 0.01 Hz to 100 kHz at open circuit potential with signal amplitude of 5.0 mV. The stability of electrode materials was tested at a charge–discharge of 10 A g⁻¹ for 2000 cycles, and the electrodes after cycling test were soaked in DI water and then collected for STEM and elemental mapping analysis. The specific capacitance C_s was calculated from the CV curves and galvanostatic charge discharge curves.

3 Results and discussion

3.1 Synthesis of NiO@MnO_x via self-assembly

Fig. 1 shows the procedure of the synthesis of NiO@MnO_x nanostructures supported on PEI-functionalized graphene. The zeta potential of purified graphene was −15.3 ± 3.4 mV, and became positive +25.2 ± 4.3 mV after the functionalization with positively charged linear PEI. In solution, nickel acetylacetone precursor dissociates into C₅H₇O₂⁻ and Ni²⁺ ions. Thus, self-assembly between the positively charged PEI-functionalized graphene and the negatively charged deprotonated acetylacetone, C₅H₇O₂⁻, occurs through the electrostatic force, followed by the instantaneous self-assembly of positively charged Ni precursors, Ni²⁺. The self-assembled Ni²⁺ precursors were in situ reduced with assistance of EG and microwave heating treatment, forming NiO NPs. The loading of NiO was 25% as confirmed by TGA (Fig. S1†). The zeta potential of NiO supported on graphene was 26.4 ± 4.7 mV, which allow the self-assembly of negatively charged MnO₄⁻ ions onto the NiO NPs surface due to the electrostatic attraction. The MnO₄⁻ ions were reduced by ethanol due to the fast reaction between KMnO₄ and ethanol, forming MnO₂ or MnO_x:

3C₂H₅OH + 4KMnO₄ → 3HC₂H₃O₂ + 4MnO₂ + 4KOH + H₂O (1)

Fig. 1 Synthesis of NiO@MnO_x nanostructures supported on PEI-functionalized graphene through self-assembly route assisted by microwave sintering.

After the deposition of MnO_x, the zeta potential changed from 26.4 mV for NiO to 8.4, −2.4 and −17.5 mV for the deposition of 5%, 10% and 25% MnO_x, respectively, indicating the self-assembly of negatively charged MnO₄⁻ ions.

3.2 Microstructure of NiO@MnO_x

Fig. 2A displays the XRD patterns of NiO and NiO@MnO_x. The broad peak at 62.9° for NiO and NiO@MnO_x are originated from NiO NPs. The deposition of MnO_x on NiO NPs does not change the XRD patterns of NiO for NiO@MnO_x (1 : 0.2) and NiO@MnO_x (1 : 0.4), indicating the MnO_x covered on the surface is most likely amorphous in nature. No obvious peaks were observed for NiO@MnO_x (1 : 1) likely due to the coverage of NiO NPs by a thicker MnO_x layer. The average crystallite size of nickel oxide for NiO, NiO@MnO_x (1 : 0.2), and NiO@MnO_x (1 : 0.4) was calculated using the peak at 62.9° based on the Scherrer formula,³³ and is 1.35, 1.31 and 1.47 nm, respectively. This indicates that the crystalline sizes of NiO NPs do not change after the deposition of MnO_x.

Fig. 2 (A) The XRD patterns and (B) high resolution XPS patterns of Ni 2p and Mn 2p regions of NiO and NiO@MnO_x supported on graphene.

The nature of the NiO and NiO@MnO_x nanostructures was further investigated by XPS (Fig. 2B). The Ni 2p XPS spectrum possesses two shakeup satellites at 855.9 eV (Ni 2p_3/2) and 873.4 eV (Ni 2p_1/2) for NiO@MnO_x (1 : 0.2), about 0.1 eV lower than those for NiO with BE of 856.0 eV and 873.5 eV respectively. As the loading of NiO : MnO_x reaches 1 : 1, the Ni 2p_3/2 and Ni 2p_1/2 peaks are observed at 855.6 eV and 873.1 eV, which is about 0.4 eV lower than that of NiO. The Mn 2p XPS spectrum displays two characteristic peaks at 642.9 eV and 654.6 eV, corresponding to Mn 2p_3/2 and Mn 2p_1/2 spin–orbit peaks for NiO@MnO_x (1 : 0.2), about 0.7 eV higher than those of MnO_x. The BE of Mn 2p_3/2 and Mn 2p_1/2 is 642.6 eV and 654.3 eV for NiO@MnO_x (1 : 1), about 0.4 eV higher than those of MnO_x. The negatively BE shift of Ni 2p and the positive BE shift of Mn 2p of NiO@MnO_x indicate the electron donation from Mn to Ni.

Fig. 3 shows the TEM images of the NiO and NiO@MnO_x. The graphene sheet was homogenously covered by high density fine NPs (Fig. 3A) as compared with the pristine graphene (Fig. S2†). No obvious large particles observed, indicating that the self-assembly method is an effective method to obtain ultrafine nanostructures. The high resolution TEM image discloses that the average size of NiO NPs is 2.3 nm (Fig. 3A). This is larger than the crystallite size of 1.3 nm obtained from the XRD. The selected-area electron diffraction (SAED) pattern of the NiO reveals well-defined diffraction rings (Fig. 3A), suggesting the polycrystalline nature of NiO. This is consistent with the HRTEM image, which shows the NiO with a lattice space 0.21 nm (the inset, Fig. 3A).

Fig. 3 TEM micrographs and histograms of (A) NiO, (B) NiO@MnO_x (1 : 0.2), (C) NiO@MnO_x (1 : 0.4) and (D) NiO@MnO_x (1 : 1) on graphene.

In the case of NiO@MnO_x (1 : 0.2), average particle size of NPs is 6.1 nm and uniformly distributed on graphene (Fig. 3B). As shown by the XRD results (Fig. 2A), the size of NiO NPs does not change after the deposition of MnO_x layer, hence the large particle size of NiO@MnO_x (1 : 0.2) indicates that the NiO NPs are covered by a thin layer of amorphous MnO_x. Based on the average size of 2.3 nm NiO NPs, thickness of the amorphous MnO_x layer deposited can be estimated to be ∼1.9 nm. The inset HRTEM in Fig. 3B further indicates that the MnO_x appears to be porous, exhibiting large amount of nanopores of ∼0.5 nm in diameter. With the increase of the MnO_x loading, the size of the NiO@MnO_x increases (Fig. 3C and D), indicating the formation of NiO@MnO_x nanostructures with thick MnO_x layer. In the case of NiO@MnO_x (1 : 0.4), the size of NPs is 9–13 nm, and the size of NiO@MnO_x (1 : 1) nanostructure is larger than 20 nm and is interconnected (Fig. 3D). The continuous and interconnected NiO@MnO_x nanostructure would make it hard to calculate the thickness of the MnO_x films. Nevertheless, based on the fact that the size of NiO NPs does not change during the deposition of MnO_x, the thickness of MnO_x could be estimated to be 3–4 nm for NiO@MnO_x (1 : 0.4) and 8–10 nm for NiO@MnO_x (1 : 1), respectively. The SAED diffraction patterns of NiO@MnO_x are identical with fcc NiO and no diffraction rings attributed to MnO_x were found (Fig. 3B–D). Both XRD and HRTEM results indicate that deposited MnO_x layer is amorphous and nanoporous. A schematic structure model of NiO@MnO_x NPs is shown in Fig. 1.

Fig. 4 shows the elemental mapping using EDS coupled with STEM of NiO@MnO_x with different NiO/MnO_x ratios. The STEM images of the NiO@MnO_x (1 : 0.2) shows the ultrafine NiO NPs covered by a thin MnO_x shell, forming porous nanostructures (Fig. 4A–C). The EDS mapping displays all three elements, Mn, Ni and O, which further indicates that the NiO NPs are covered by a thin layer of amorphous MnO_x. With the increase of the MnO_x loading, the EDS peak intensity significantly increases and at the same time, the intensity of the color associated with Mn also increases, indicating an increase of MnO_x thickness on the NiO NP surface (Fig. 4D). In the case of NiO@MnO_x (1 : 0.4), the number of interconnected nanostructures increase (Fig. 4B). As the NiO/MnO_x ratio reaches to 1 : 1, the structure shows NiO NPs embedded in a thick and continuous amorphous MnO_x layer (Fig. 4C).

Fig. 4 HAADF-STEM-EDS mapping images (A) NiO@MnO_x (1 : 0.2), scale bar = 7 nm, (B) NiO@MnO_x (1 : 0.4), scale bar = 10 nm. (C) NiO@MnO_x (1 : 1), scale bar = 20 nm. (D) EDS spectra originated from mapping area of (a) NiO, (b) NiO@MnO_x (1 : 0.2), (c) NiO@MnO_x (1 : 0.4), (d) NiO@MnO_x (1 : 1), (E) nitrogen sorption isotherm profiles of NiO and NiO@MnO_x, and (F) schematic diagram showing the structure evolution of NiO@MnO_x with the increase of MnO_x loading.

The nitrogen sorption isotherm profiles of NiO and NiO@MnO_x show a shape similar with type-IV with a H3 hysteresis loop associated with slit-shaped pores (Fig. 4E).^34,35 The knee-point at low P/P₀ of isotherm profiles is due to the formation of micropores likely resulted from the inherent tunnels or pores of amorphous MnO_x. The BET surface area increase from 149.8 m² g⁻¹ for NiO to 280.1, 340.8 and 430.9 for NiO@MnO_x (1 : 0.2), NiO@MnO_x (1 : 0.4) and NiO@MnO_x (1 : 1), respectively. The increase of the BET surface area could be due to the connection and formation of hierarchy microstructures (the increase of hysteresis loop) and the micropores inherent from the porous structure of manganese oxide. Fig. 4F shows the formation of NiO@MnO_x nanostructures with different NiO/MnO_x ratios on graphene.

3.3 Supercapacitance performance of NiO@MnO_x

Fig. 5A presents typical CV curves of the NiO, NiO@MnO_x at a scan rate of 50 mV s⁻¹. All the CV curves consist of a pair of strong redox peaks resulted from the transformation between Ni(II) and Ni(III) involving the reversible process of insertion and extraction of protons/ions. The results indicate that the capacitance of NiO and NiO@MnO_x is mainly governed by faradaic redox reactions of the active material, NiO. The anodic and the cathodic peaks (E_a and E_c) of NiO are symmetric with the E_a and E_c observed at 0.357 V and 0.183 V vs. SCE. On the other hand, both E_a and E_c values are negatively shifted to 0.295 and 0.159 V for the reaction on NiO@MnO_x (1 : 0.2), 0.277 and 0.144 V on NiO@MnO_x (1 : 0.4) and 0.24 and 0.165 V for NiO@MnO_x (1 : 1), respectively, as compared to that on NiO. Consequently, the ΔE (i.e., E_a − E_b) decreases from 174 mV for NiO to 136, 133 and 75 mV for NiO@MnO_x (1 : 0.2), NiO@MnO_x (1 : 0.4) and NiO@MnO_x (1 : 1) respectively, indicating the improved reversibility of the transformation reaction on NiO@MnO_x. The reversibility increases with the increase of the thickness of the amorphous MnO_x layer likely due to the synergistic effect between the NiO and MnO_x resulted from the donation of electron form Mn to Ni (Fig. 2B). The downshift in the Mn 2p_3/2 leads to a decrease of the affinity and adsorption energy of MnO_x for oxygen species,^36–38 which may benefit the charge transfer process on the embedded NiO NPs.

Fig. 5 (A) cyclic voltammetry (CV) curves of NiO and NiO@MnO_x at scan rate of 50 mV s⁻¹ in 2 M KOH solution; (B) plots of C_s (based on total materials) calculated based on CV curves as a function of scan rates, (C) charge and discharge curves of NiO and NiO@MnO_x at a current density of 2 A g⁻¹, (D) plots of C_s (based on total materials) calculated based on charge and discharge curves as a function of current densities, (E) Nyquist plots of NiO and NiO@MnO_x, measured at 0 V, and (F) CV measured at 50 mV s⁻¹ and charge and discharge curves measured at 2 A g⁻¹ of MnO_x supported on graphene.

With the increase of NiO/MnO_x ratio, the intensity of anodic and the cathodic peaks decreased because the covering of the NiO with a thicker layer of MnO_x would increase the ionic insertion length and leads to a lower active material utilization. The anodic and cathodic peaks were symmetric, suggesting excellent reversibility for both NiO and NiO@MnO_x electrodes. The minimal changes in the shape of the CV curves as the scan rate increases from 10 to 100 mV s⁻¹ reveal the excellent electron conductivity of NiO and NiO@MnO_x supported on graphene (Fig. S3†). The anodic and cathodic peaks, however, are positively shifted with increasing scan rate due to the internal resistance of the electrode. The anodic peak of NiO shifted by 70 mV from 0.28 V at the scan rate of 10 mV s⁻¹ to 0.35 V at 100 mV s⁻¹, which is higher than a positively shift of 60, 55 and 30 mV for NiO@MnO_x (1 : 0.2), NiO@MnO_x (1 : 0.4) and NiO@MnO_x (1 : 1) measured under the same conditions. This indicates the coverage of MnO_x on the NiO nanostructures surface improve the reversibility and rate capacity.

The specific capacitance, C_s of NiO calculated from the CV curves at scan rate of 10 mV s⁻¹ is 865 F g⁻¹ based on the total mass of the electrode material (C_total, NiO + graphene) and 3460 F g⁻¹ based on the active material NiO (C_active). The C_total and C_active is 1037 F g⁻¹ and 3456 F g⁻¹, respectively, for NiO@MnO_x (1 : 0.2). However, with the increased loading of MnO_x, the C_total and C_active decreased to 807 and 2305 F g⁻¹ for NiO@MnO_x (1 : 0.4), and 297 and 594 F g⁻¹ for NiO@MnO_x (1 : 1), respectively. The similar C_active between NiO and NiO@MnO_x (1 : 0.2) indicate that the deposition of a thin amorphous MnO_x shell (∼1.9 nm) would not impede the fast electron transport and ion-diffusion to reach the active NiO NPs. With the increase in the scan rate, the C_total decreases gradually (Fig. 5B), which can be attributed to the inaccessible of the electrolytic ions (i.e., OH⁻) to the active NiO NPs for the charge storage at high scan rates. In the case of NiO NPs, C_total decreases to 252.7 F g⁻¹ at scan rate of 200 mV s⁻¹, which is only 29% of 865 F g⁻¹ measured at scan rate of 10 mV s⁻¹. The C_total of NiO@MnO_x (1 : 0.2), NiO@MnO_x (1 : 0.4) and NiO@MnO_x (1 : 1) are 742.8, 588.7, and 230.8 F g⁻¹, respectively, at scan rate of 200 mV s⁻¹, which is 71.6%, 72.7% and 77.5% of that at scan rate of 10 mV s⁻¹, much higher than 29% observed on NiO NPs. The much higher reduction in supercapacitance at high scanning rates for the reaction on NiO indicates that the rate capability of the NiO@MnO_x electrodes is significantly higher than that of NiO electrode.

The electrochemical properties of NiO and NiO@MnO_x were further characterized by galvanostatic charging and discharging method (Fig. 5C). The C_s of the electrodes was calculated from the discharge curve. The C_total and C_active of the electrode material is 924 and 3696 F g⁻¹ for NiO, 966 and 3222 F g⁻¹ for NiO@MnO_x (1 : 0.2), 711 and 2031 F g⁻¹ for NiO@MnO_x (1 : 0.4), 263 and 526 F g⁻¹ for NiO@MnO_x (1 : 1), respectively, measured at a current density of 2 A g⁻¹. For the reaction on NiO without deposition of amorphous MnO_x, the C_s decreases rapidly with the increase of current density (Fig. 5D and S4†). C_total obtained at 40 A g⁻¹ is 136 F g⁻¹, which is 14.8% of 924 F g⁻¹ measured at a current density of 2 A g⁻¹. This indicates that pristine NiO NPs without the amorphous MnO_x layer have a very low rate capability, which is consistent with that reported in the literature. For example, Liu et al. reported that the C_s of Ni(OH)₂ NPs with average size ∼4.8 nm supported on reduced graphene oxide sheets decreased from 1578 F g⁻¹ at 1 A g⁻¹ to 905 at 10 A g⁻¹, a reduction of 43% in capacitance.¹⁴

On the other hand, in the case of NiO@MnO_x electrode materials, C_total measured at a current of 40 A g⁻¹ is 713, 585 and 214 F g⁻¹ for NiO@MnO_x (1 : 0.2), NiO@MnO_x (1 : 0.4) and NiO@MnO_x (1 : 1), respectively, which is 73.8%, 82.3% and 81.3% of that at current density of 2 A g⁻¹ (Fig. 5D). This shows that the NiO@MnO_x structure can retain 70–80% of its initial value when the discharge current density increases by 20 times, significantly better than 14.8% measured on NiO. The results show that the NiO@MnO_x structure with a thin layer of amorphous MnO_x can efficiently utilize the active materials even at high current density (high power demand). More importantly, the columbic efficiency (C_s discharge/C_s charge × 100%) of the NiO@MnO_x is significantly higher than that of NiO. For example, the columbic efficiency is 100% for NiO@MnO_x (1 : 0.2) at 40 A g⁻¹, which is higher than 87.5% of NiO measured under identical conditions.

The impedance responses of NiO are characterized by a semicircle at high frequencies and a low frequency linear tail attributed to the ion-diffusion process between Warburg diffusion and ideal capacitive ion diffusion (Fig. 5E).^11,16,39 The charge transfer resistance (R_ct) associated with the high frequency arc of NiO is significantly higher than that of NiO@MnO_x, revealing the decrease of charge transfer resistance after covering the NiO with amorphous MnO_x. Most interesting, the NiO@MnO_x (1 : 0.2) exhibits the low frequency tail with a slope of ∼60°, as compared with a slope of ∼45° for the reaction on NiO@MnO_x (1 : 0.4), NiO@MnO_x (1 : 1) and NiO. This indicates the lower Warburg impedance and better ion diffusion of NiO@MnO_x (1 : 0.2),^11,16 consistent with its high performance.

The capacitance behavior of MnO_x supported on graphene was also studied and the results are shown in Fig. 5F. The CV curve shows a rectangular shape, typical characteristics of double layer capacitance. Based on the discharge curve, the C_total and C_active is 12.4 and 49.6 F g⁻¹, lower than 110 F g⁻¹ at 2 A g⁻¹ reported on MnO₂ nanosheets dispersed on single walled carbon nanotubes.⁴⁰ The C_s of MnO_x is substantially lower than that of NiO@MnO_x, indicating that the high C_s of NiO@MnO_x is mainly due to the NiO NPs and the contribution of deposited amorphous MnO_x layer to the overall capacitance performance of NiO@MnO_x is very low and can be neglected.

3.4 Stability of NiO@MnO_x

The long-term stability of NiO and NiO@MnO_x were investigated by galvanostatic charging and discharging cyclic tests at a current density of 10 A g⁻¹ in 2 M KOH solution and the results are shown in Fig. 6. NiO displays very poor durability and the capacitance (C_total) decrease rapidly from 773 F g⁻¹ to ∼234 F g⁻¹ after 500 cycles and reached 46.4 F g⁻¹ after 2000 cycles, a reduction of 94% of the initial capacitance. The shape of the charge and discharge curves after 2000 cycles is also significantly distorted, indicating the very poor stability of the NiO NPs supported on graphene. The poor stability of NiO NPs in this study is likely due to the aggregation of ultrafine NiO NPs. In contrast, the shape of the charge and discharge curves remains unchanged with the cyclic tests for NiO@MnO_x core–shell structures. For example, the initial capacitance of NiO@MnO_x (1 : 0.2) is 705 F g⁻¹ and increases slightly to 825 F g⁻¹ after 200 cycles most likely due to the activation of NiO NPs. After tests for 2000 cycles, there is no decrease in the capacitance, indicating the excellent stability of NiO@MnO_x (1 : 0.2) electrode materials. Similar results were also obtained for NiO@MnO_x (1 : 0.4) and NiO@MnO_x (1 : 1) electrode materials (Fig. 6B).

Fig. 6 (A) The 1^st and 2000^th charge and discharge curves of NiO, NiO@MnO_x at current density of 10 A g⁻¹ in 2 M KOH and (B) plots of C_s (based on total materials) of NiO and NiO@MnO_x electrode materials as a function of the cycle numbers.

The microstructure of NiO, NiO@MnO_x (1 : 0.2) and NiO@MnO_x (1 : 0.4) after cyclic test was analyzed by HAADF-STEM, and the results are shown in Fig. 7. In the case of NiO, a significant change of morphology was observed. NiO NPs aggregated into large particles with size around 12–50 nm (average ∼27.4 nm) (Fig. 7A), increased by an order of magnitude as compared to the initial particle size of 2.3 nm. In the case of NiO@MnO_x, the change in the microstructure is much smaller. After the 2000 charge/discharge cycles, NiO NPs in the NiO@MnO_x matrix can be clearly identified as indicated by the significantly increase intensity of the color associated with Ni (green in this case), different to that before the test (Fig. 4A). This indicates that originally embedded NiO NPs may extrude through the amorphous MnO_x. In the case of NiO@MnO_x (1 : 0.2), NiO NPs also grew to 2.0–6.5 nm (average ∼ 3.6 nm, Fig. 7B). However, NiO NPs are still clearly separated and the agglomeration of embedded NiO NPs is substantially smaller than that observed in the case of NiO. This evidently demonstrates the effectiveness of the confinement effect of the amorphous MnO_x on the growth and agglomeration of NiO. Similar microstructure evolution was also observed for NiO@MnO_x (1 : 0.4) (Fig. 7C). In this case, the embedded NiO NPs size is in the range of 2.5–9 nm (average ∼ 4.4 nm).

Fig. 7 HAADF-STEM, EDS mapping and histograms of particle size distribution of (A) NiO, (B) NiO@MnO_x (1 : 0.2), (C) NiO@MnO_x (1 : 0.4) after 2000 cycles. In histograms, black curve is the particle sizes of NiO before cycling and red curve is the NiO sizes after 2000 cycles. Particle size after cycling was calculated based on the STEM images, and the mapping results indicate bright dots are NiO.

3.5 Structural confinement and mass transfer promotion effect of nanoporous and amorphous MnO_x

The best results in this study were obtained on amorphous MnO_x confined NiO with weight ratio of 1 : 0.2 supported on graphene, NiO@MnO_x (1 : 0.2). NiO@MnO_x (1 : 0.2) electrode materials not only exhibit high C_s (C_active = 3222 F g⁻¹), but also possess high active material utilization, outstanding rate capability and columbic efficiency as well as excellent structurally stability. For example, the C_active of the NiO@MnO_x (1 : 0.2) measured at 2 A g⁻¹ is 3222 F g⁻¹, which is comparable or significantly higher than those of advanced nanostructures reported (based on active materials),^{11,14,16,21,41–44} including NiO/ITO nanowire shell–core structure (1025 F g⁻¹ at 2 A g⁻¹),⁴¹ porous nickel hydroxide–manganese dioxide–reduced graphene oxide (1985 F g⁻¹ at 2 A g⁻¹),¹¹ porous nickel hydroxide–manganese dioxide hybrid nanosheets (2628 F g⁻¹ at 2 A g⁻¹),⁴² amorphous nickel hydroxide (2188 F g⁻¹ at 2 A g⁻¹),⁴³ Ni(OH)₂ supported on reduced graphene oxide (1717 F g⁻¹ at 2 A g⁻¹),¹⁴ ultrathin β-Ni(OH)₂ nanoplates on nickel-coated carbon nanotubes (1807 F g⁻¹ at 2 A g⁻¹)¹⁶ and NiO with hierarchical porous ball like surface morphology with uniform ripple-shaped pores (370 F g⁻¹ at 2 A g⁻¹).⁴⁵ The size of the porous or nanostructured NiO electrode materials is critical as it is widely known that ions insertion/extraction depth is limited in the range of few nanometers on the surface of electrode materials. Most significant, NiO@MnO_x (1 : 0.2) electrode materials show a much better rate of capacitance as compared to that reported in the literature. For example, Zhu et al. synthesized ultrathin nickel hydroxide nanosheets and reported a very high C_active value of 3650 F g⁻¹ measured at 2 A g⁻¹.²¹ However, with the increase of discharge current to 16 A g⁻¹, the C_active decreased rapidly to 2680 F g⁻¹, lower than 2781 F g⁻¹ at 20 A g⁻¹ for NiO@MnO_x (1 : 0.2) nanostructures. Fig. 8 compares the capacitance performance at selected discharge currents of the advanced NiO-based electrode materials reported in the literature with the NiO@MnO_x (1 : 0.2) electrode materials in this study.

Fig. 8 Comparison of specific capacitance based on active material of NiO@MnO_x (1 : 0.2) nanostructure in this study with those reported in the literature. Empty square (): hexagonal Ni(OH)₂ nanoplates on graphene,⁴⁴ empty circle (): porous nickel hydroxide–manganese oxide on graphene,¹¹ empty triangle (): amorphous nickel hydroxide,⁴³ empty reversed triangle(): ultrathin porous nickel hydroxide–manganese dioxide hybrid,⁴² empty diamond (): ultrathin β-Ni(OH)₂ nanoplates vertically grown on nickel-coated carbon nanotubes,¹⁶ empty star (): ultrathin nickel hydroxide and oxide nanosheets,²¹ solid circle () NiO NPs supported on graphene (this study), and solid square (): NiO@MnO_x (1 : 0.2) nanostructure supported on graphene (this study).

The high performance of NiO@MnO_x (1 : 0.2) is clearly due to the ultrafine NiO NPs and this is supported by the initially high capacitance of 3696 F g⁻¹ of ultrafine NiO NPs with size ∼2.3 nm supported on graphene. The ultrafine NiO NPs would allow short and efficient ion pathways for the charge and discharge process, leading to a high utilization of active material and favorable reaction kinetics. However, without the confinement of amorphous MnO_x, NiO NPs grow and agglomerate quickly during the charge/discharge cycles, forming large particles with size of 12–50 nm (Fig. 7A). This explains the rapid loss of the capacitance of NiO NPs during the durability test (Fig. 6). The deposition of a thin layer of amorphous MnO_x, such growth and agglomeration of embedded NiO NPs is significantly alleviated by the confinement of amorphous MnO_x, as indicated by the 100% retention of the capacitance after the durability test and substantially grain growth and aggregation of active NiO NPs of NiO@MnO_x (Fig. 6 and 7). Fig. 9 shows the structural model of agglomeration NiO NPs and structural confinement effect on NiO@MnO_x nanostructure.

Fig. 9 (A) Aggregation of ultrafine NiO NPs, and (B) structure confinement of the ultrafine NiO NPs core with amorphous MnO_x shell. The inset picture is the structure model of NiO@MnO_x.

The electricity storage and discharge of NiO based electrode materials is through the following reactions in alkaline solution:

Ni(OH)₂ + OH⁻ ↔ NiOOH + H₂O + e (2)

Thus, in the case of ultrafine NiO NPs, the rate capacity would strongly depend on the OH⁻ concentration at the surface. With the increase of scanning rate or charge/discharge currents, capacitance of NiO decreased significantly (Fig. 5B and D), due to the significantly depletion of OH⁻ concentration in the surface region. In the case of NiO@MnO_x electrode materials, there is significantly enhanced rate capacity and reversibility, indicated by the negatively shifted ΔE for the redox reaction on NiO@MnO_x and the much smaller reduction in C_s with the increase of discharge current density, as compared with NiO NPs.

Manganese oxides have been widely studied for energy storage and molecular sieve due to their inherent nanopores and interconnected channels, based on the MnO₆ octahedral building blocks.⁴⁶ The much higher rate capacity and reversibility of NiO@MnO_x electrode materials clearly indicate that the presence of a nanoporous and amorphous MnO_x layer substantially reduces the depletion of the reactant concentration, OH⁻ at the surface of embedded NiO NPs, as compared to the free standing NiO NPs, an indication of enhanced mass transfer through the nanoporous structure of the deposited amorphous MnO_x. Tagliazucchi and Szleifer proposed that ions transportation in nanopores would be enhanced driven by electrical field.⁴⁷ Liu et al. revealed that the shorter the nanopore and/or the smaller its radius, the faster the osmotic flow.⁴⁸ This implies that mass transfer rate within the nanochannels would be enhanced,^49,50 and the nanoporous and amorphous MnO_x layer could act as a “ion pump” to enhanced the mass transfer rate.⁵¹ This may explain the much less reduction in the capacitance of NiO@MnO_x electrode materials as a function of the scan rate and charge/discharge currents, as compared to that of NiO. The negatively shift of Ni 2p peaks and the positively shift of Mn 2p peaks for NiO@MnO_x (see Fig. 2B) also indicate that the MnO_x could donation electron to the NiO and facilitate the charge and discharge of embedded NiO NPs.⁵² The results of this study demonstrate that the presence of microporous MnO_x not only substantially inhibits the agglomeration of NiO NPs, but also enhances the mass transfer of the reactants for the charge/discharge, leading to highly active, durable, high utilization and high rate capacity of ultrafine NiO active materials for supercapacitors.

4 Conclusions

In this paper, we developed a new strategy to explore the excellent properties of nanoporous and amorphous MnO_x in the development of ultrafine NiO NPs active materials for supercapacitors. NiO with size 2.3 nm homogenously supported on graphene and structurally confined by a thin layer of nanoporous and amorphous MnO_x film is very active material with high C_s, high material utilization, outstanding rate capability, columbic efficiency as well as long-term durability. The capacitive performance of NiO@MnO_x depends on the weight ratio of NiO/MnO_x and the best results were obtained on NiO@MnO_x with NiO : MnO_x weight ratio of 1 : 0.2, achieving a high capacitance of 966 F g⁻¹ based on total electrode materials and 3222 F g⁻¹ based on active materials at discharge current density of 2 A g⁻¹. Most important, the materials retain 100% capacitance after 2000 cycles at charge and discharge current density of 10 A g⁻¹. As shown in Fig. 8, the performance and stability of NiO@MnO_x (1 : 0.2) are significantly better than the advanced NiO-based nanostructured electrode materials reported in the literature. The nanoporous and amorphous MnO_x shell not only structurally confines the ultrafine NiO NPs and prevents the active materials from aggregation during the cyclic charge and discharge but also functions as an “ionic pump”, promoting the mass transfer for the reaction and substantially enhancing the rate capacity. The structurally confined ultrafine active NPs provide an effective strategy to develop highly active and durable ultrafine nanosized materials with high rate capacity for supercapacitors.

Acknowledgements

This research was supported by the Australian Research Council under Discovery Project Funding Scheme (project number: DP150102044 & DP150102025) and the Major International (Regional) Joint Research Project of NNSCF, China (51210002). The authors acknowledge the facilities, and the scientific and technical assistance of the Australian Microscopy and Microanalysis Research Facility at the Centre for Microscopy, Characterisation & Analysis, the University of Western Australia, a facility funded by the University, State and Commonwealth Governments. The authors would also like to acknowledge the technical assistance of Dr Kane O'Donnell with XPS analysis and the WA X-Ray Surface Analysis Facility, funded by an Australian Research Council LIEF grant (LE120100026).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra04880c

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