**Meet the editor**

Zoran M. Stevic (1958, Serbia) is a full-time professor at the University of Belgrade, Serbia, and he is the chief of Department at the Technical Faculty in Bor. He received his PhD diploma from the Faculty of Electrical Engineering in Belgrade, University of Belgrade. His research areas include energy efficiency, renewable energy sources, system modeling, computer measurement and process

control, sensors, power electronics, optoelectronics, supercapacitors, electrochemistry and IR thermography. He is a member of IEEE and IBPSA. He published over 300 papers (over 40 refereed full papers in scientific journals), 5 books, and 4 chapters. Also, he was a project coordinator and a member of working group at many scientific and technical projects.

## Contents

**Preface XI**


#### **X** Contents

#### **Section 2 Supercapacitor Applications 119**

Chapter 7 **Development and On‐Orbit Demonstration of Lithium‐Ion Capacitor‐Based Power System for Small Spacecraft 121** Masatoshi Uno and Akio Kukita

#### Chapter 8 **Interdigitated MEMS Supercapacitor for Powering Heart Pacemaker 145** Hafzaliza Erny Zainal Abidin, Azrul Azlan Hamzah, Jumril Yunas, Mohd Ambri Mohamed and Burhanuddin Yeop Majlis

#### Chapter 9 **Power Management in Supercapacitor-Based Wireless Sensor Nodes 165** Hengzhao Yang and Ying Zhang

## Preface

**Section 2 Supercapacitor Applications 119**

**VI** Contents

**Pacemaker 145**

**Sensor Nodes 165**

Masatoshi Uno and Akio Kukita

Hengzhao Yang and Ying Zhang

Chapter 7 **Development and On‐Orbit Demonstration of Lithium‐Ion**

Chapter 8 **Interdigitated MEMS Supercapacitor for Powering Heart**

Chapter 9 **Power Management in Supercapacitor-Based Wireless**

Mohd Ambri Mohamed and Burhanuddin Yeop Majlis

**Capacitor‐Based Power System for Small Spacecraft 121**

Hafzaliza Erny Zainal Abidin, Azrul Azlan Hamzah, Jumril Yunas,

Standard systems for electrical energy storage are capacitors and rechargeable batteries. Re‐ cently, more attention was paid to supercapacitors as a qualitatively new type of capacitors and fuel cells as sources of electricity. With the development of new materials and technolo‐ gies, a very large surface area and very small interelectrode distances have been reached. This equates to an extremely high capacitance (several orders of magnitude greater than conventional capacitors), so such systems are called supercapacitors.

A lot of teams and laboratories around the world are working on the development of super‐ capacitors, while their better and better performance enables wider usage. The goal of this book is to bring closer to the readers new supercapacitor technologies that are changing the present and the future of electricity storage.

The book is divided into two sections. The first section deals with the current status and trends of development of supercapacitors, while the second section deals with applications of supercapacitors.

> **Prof. Dr. Zoran Stevic** University of Belgrade Technical Faculty Bor Belgrade, Serbia

**Supercapacitor Design**

#### **High Volumetric Performance Supercapacitors with Controlled Nanomorphology High Volumetric Performance Supercapacitors with Controlled Nanomorphology**

Yue Zhou and Qiming Zhang Yue Zhou and Qiming Zhang

Additional information is available at the end of the chapter Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/65186

#### **Abstract**

Supercapacitor is one of the promising energy storage devices due to its relatively higher energy density compared with dielectric capacitor and higher power density and longer cycle life time (>millions) than conventional battery. In order to satisfy various requirements for energy technologies, supercapacitors with higher energy and power densities are required. In this chapter, we improved the electrochemical performance largely compared with commercial product through controlling the nanomorphology of cells. Meanwhile, although many past research programs have focused mainly on gravimetric energy densities, here we have also devoted efforts to study and develop nanomorphologic structures to realize high volumetric energy and power densities, since device volume is another critical and key performance parameter. Moreover, fundamental studies have been carried out on the mobile ion transport and storage in the nanostructures developed in this chapter.

**Keywords:** supercapacitor, carbon nanotubes, graphene, volumetric performance, nanomorphology control

## **1. Introduction**

In order to meet different application requirements and also for fundamental studies of ion transport and storage in nanoporous media, we selected carbon-based electrodes with unique and controlled nanomorphologies: highly aligned carbon nanotube (A-CNT) forests. As synthesized, A-CNTs have low volume fraction of CNT (~1%). Traditional method to achieve high volumetric performance of A-CNTs was to employ liquid surface tension densification, which did not have control on the nanomorphology of A-CNTs after densification. In this

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

work, making use of and improving upon the mechanical densification method, we achieved 40% volumetric density of A-CNTs. A series of supercapacitor electrodes have been developed and characterized. As can be seen in the section, the high volume fraction of A-CNTs and highly aligned ion channels in the nanoporous electrodes lead to the superior performance of the supercapacitors compared with any CNT-based supercapacitor electrodes studied earlier. The supercapacitors exhibit a volumetric power density, 25 kW/L (and gravimetric power density 50 kW/kg) for the capacitor cell with 0.8-mm thick A-CNTs, compared with the similar capacitors using A-CNTs densified by the liquid collapsing method, 13.4 kW/L (24 kW/kg) for cells with 0.5-mm thick A-CNTs. The study also shows the importance of the ionic conductivity of electrolytes in controlling the power and energy densities of the supercapacitors.

Since the energy and power densities of supercapacitors are directly proportional to the square of cell operation voltage V (~ V2 ), raising the cell operation voltage will have great potential to enhance the energy and power densities. Asymmetric supercapacitors, which allow for optimization of both cathode and anode simultaneously, provide an attractive approach to raise the cell operation voltage, besides other properties. In this chapter, we have investigated asymmetric supercapacitor configurations for carbon-based electrodes for one electrode and conducting polymer (CP)-coated A-CNTs for the other one, based on their electrochemical windows. Here, we investigate the asymmetric supercapacitors where both electrodes are tailored, respectively, to improve the device electrochemical performances such as specific capacitance and the electrochemical window. Hence, operating voltage is increased. The conformal vapor is utilized to deposit CP on the A-CNTs, enhancing the charge storage capability of the electrode, while the aligned nanowire morphology of the composite electrode exhibits straight fast ion transport pathways to enhance power. The a-graphene electrode, which is fabricated through a self-assembly process, shows the high active material density. Combining with a high specific surface area of 3000 m2 /g, the electrode yields very high specific volumetric capacitance, energy, and power densities. As a result, the asymmetric supercapacitors show an energy density 113 Wh/L (176 Wh/kg), which is the highest among all carbonbased supercapacitors, and a power density 149 kW/L (233 kW/kg).

#### **2. Symmetric supercapacitors with controlled unique nanomorphology**

Presently, most supercapacitors are fabricated from activated carbon (AC), which possesses a very large specific surface area (1000–2000 m2 /g). Recent advances have demonstrated many attractive features of utilizing A-CNTs for supercapacitors with nanoporous electrodes, especially the parallel ion channels formed by the A-CNTs that improve the ion transport, as schematically illustrated in **Figure 1(a)**, compared with randomly arranged nanoporous electrodes from AC, forming tortuous ion transport pathways [1, 2]. Consequently, supercapacitor cells with A-CNTs exhibit higher power and energy density than that from AC. Since as grown A-CNT forests have CNT volumetric density <5 vol%, A-CNTs should be densified to reach higher A-CNT volumetric density for practical supercapacitor applications. In the past decade, many works have been conducted to produce aligned A-CNTs with high CNT density to achieve high volumetric capacitance, energy density, and power density, which are critical for modern electric and electronic systems to realize compact device size and increased functions within given device volumes. For example, Futaba et al. [3] employed a liquid collapsing method to densify A-CNTs with a high density (~50 vol%). Here, the mechanical densification method has several advantages compared with the liquid collapsing method. This method provides a precise control on the density of the final A-CNTs, ranging from the original 1 vol% A-CNTs to >50 vol%. Besides, A-CNT samples with different sizes can be densified with precisely controlled nanomorphology (alignment), and hence this method provides a realistic pathway for scaling up high-density A-CNTs for large-scale manufacturing of supercapacitors from A-CNTs of ultrahigh volume density [4]. In addition, the availability of A-CNTs with different density and hence different ion channel sizes also creates a great opportunity to study how the ion channel size formed by the A-CNTs affect ion transport and storage, which is of great importance in developing supercapacitors with high energy, power density, and tailored performance. We have also studied the influence of the ionic conductivity of electrolyte on the electrochemical performance of the supercapacitors. An imidazolium-based ionic liquid (1-ethyl-3-methylimidazolium-tetrafluoroborate, EMI-BF4) was chosen for the study. Imidazolium ionic liquids (ILs) due to their high ionic conductivity and wide elec-

work, making use of and improving upon the mechanical densification method, we achieved 40% volumetric density of A-CNTs. A series of supercapacitor electrodes have been developed and characterized. As can be seen in the section, the high volume fraction of A-CNTs and highly aligned ion channels in the nanoporous electrodes lead to the superior performance of the supercapacitors compared with any CNT-based supercapacitor electrodes studied earlier. The supercapacitors exhibit a volumetric power density, 25 kW/L (and gravimetric power density 50 kW/kg) for the capacitor cell with 0.8-mm thick A-CNTs, compared with the similar capacitors using A-CNTs densified by the liquid collapsing method, 13.4 kW/L (24 kW/kg) for cells with 0.5-mm thick A-CNTs. The study also shows the importance of the ionic conductivity of electrolytes in controlling the power and energy densities of the

Since the energy and power densities of supercapacitors are directly proportional to the square

enhance the energy and power densities. Asymmetric supercapacitors, which allow for optimization of both cathode and anode simultaneously, provide an attractive approach to raise the cell operation voltage, besides other properties. In this chapter, we have investigated asymmetric supercapacitor configurations for carbon-based electrodes for one electrode and conducting polymer (CP)-coated A-CNTs for the other one, based on their electrochemical windows. Here, we investigate the asymmetric supercapacitors where both electrodes are tailored, respectively, to improve the device electrochemical performances such as specific capacitance and the electrochemical window. Hence, operating voltage is increased. The conformal vapor is utilized to deposit CP on the A-CNTs, enhancing the charge storage capability of the electrode, while the aligned nanowire morphology of the composite electrode exhibits straight fast ion transport pathways to enhance power. The a-graphene electrode, which is fabricated through a self-assembly process, shows the high active material density.

volumetric capacitance, energy, and power densities. As a result, the asymmetric supercapacitors show an energy density 113 Wh/L (176 Wh/kg), which is the highest among all carbon-

**2. Symmetric supercapacitors with controlled unique nanomorphology**

Presently, most supercapacitors are fabricated from activated carbon (AC), which possesses

many attractive features of utilizing A-CNTs for supercapacitors with nanoporous electrodes, especially the parallel ion channels formed by the A-CNTs that improve the ion transport, as schematically illustrated in **Figure 1(a)**, compared with randomly arranged nanoporous electrodes from AC, forming tortuous ion transport pathways [1, 2]. Consequently, supercapacitor cells with A-CNTs exhibit higher power and energy density than that from AC. Since as grown A-CNT forests have CNT volumetric density <5 vol%, A-CNTs should be densified to reach higher A-CNT volumetric density for practical supercapacitor applications. In the past decade, many works have been conducted to produce

), raising the cell operation voltage will have great potential to

/g, the electrode yields very high specific

/g). Recent advances have demonstrated

supercapacitors.

of cell operation voltage V (~ V2

4 Supercapacitor Design and Applications

Combining with a high specific surface area of 3000 m2

a very large specific surface area (1000–2000 m2

based supercapacitors, and a power density 149 kW/L (233 kW/kg).

**Figure 1.** (a) The illustration of the tortuous ion transport paths in the activated carbon electrodes as well as parallel ion pathways in the A-CNTs. (b) Optic images, mechanical densification process, and SEM images of 1% Vf and 40% Vf A-CNTs.

trochemical window have been investigated very extensively as the electrolytes for supercapacitors. A mixture of an IL such as EMI-BF4 and molecular liquid such as propylene carbonate (PC) can lead to marked enhancement (more than two times higher) in the ionic conductivity compared with the pure EMI-BF4. The experimental results show that increased ionic conductivity of the electrolyte can lead to a large increase in the power density (more than double the power density using EMI-BF4/PC compared with pure EMI-BF4) of the supercapacitors.

The A-CNTs in this work were fabricated through a modified chemical vapor deposition (CVD) method on silicon wafers and iron (Fe) on alumina was used as the catalyst. The as-grown carbon nanotube forests have a 1% volume fraction (Vf) with density of 109 –1010 CNTs cm−2. The average diameter of nanotubes is 8 nm with 3–5 multiwalls. The spacing between nanotubes is approximately 80 nm in the forest. For the high Vf A-CNT synthesis, the freestanding CNT forests were released from the silicon wafer using a razor blade. And then as shown in **Figure 1(b)**, the forests were subjected to a mechanical biaxial densification process in two orthogonal directions. In this method, the A-CNT forest was densified along one direction firstly to a fixed distance by utilizing a mechanical bar, and then another mechanical bar in the orthogonal direction was employed to compress the A-CNT forest to the final volume fraction. By varying the inter-CNT distance in densification process, A-CNT forests with different Vfs can be obtained.

**Figure 2.** Electrochemical performance of supercapacitors with A-CNT electrodes and EMI-BF4/PC electrolyte at 4 V: (a) gravimetric cyclic voltammograms and (b) volumetric cyclic voltammograms of A-CNT electrodes with 1% and 40% Vf at 100 mV s−1.

The A-CNT forests were used as the electrodes of the supercapacitors. 3 M EMI-BF4 (1-ethyl-3 methylimidazolium tetrafluoroborate) in propylene carbonate (PC) was used as the electrolyte to improve ionic conductivity. Compared with pure ionic liquid, the ionic liquids/molecular liquids (IL/ML) mixture will have higher conductivity. Polypropylene porous membrane (Celgard 3501, Celgrad LLC) with 25 μm thickness was used as the separator. The sandwich architecture (electrode/separator/electrode) was placed between two pieces of Au-coated steel plates, which served as the current collectors. Then, the whole cell was housed in a Teflon holder. As a comparison, supercapacitor electrodes, which are made from the activated carbon powders mixing with 10 wt% PTFE and 10 wt% carbon black, were also fabricated using the standard method.

trochemical window have been investigated very extensively as the electrolytes for supercapacitors. A mixture of an IL such as EMI-BF4 and molecular liquid such as propylene carbonate (PC) can lead to marked enhancement (more than two times higher) in the ionic conductivity compared with the pure EMI-BF4. The experimental results show that increased ionic conductivity of the electrolyte can lead to a large increase in the power density (more than double the power density using EMI-BF4/PC compared with pure EMI-BF4)

The A-CNTs in this work were fabricated through a modified chemical vapor deposition (CVD) method on silicon wafers and iron (Fe) on alumina was used as the catalyst. The as-grown

The average diameter of nanotubes is 8 nm with 3–5 multiwalls. The spacing between nanotubes is approximately 80 nm in the forest. For the high Vf A-CNT synthesis, the freestanding CNT forests were released from the silicon wafer using a razor blade. And then as shown in **Figure 1(b)**, the forests were subjected to a mechanical biaxial densification process in two orthogonal directions. In this method, the A-CNT forest was densified along one direction firstly to a fixed distance by utilizing a mechanical bar, and then another mechanical bar in the orthogonal direction was employed to compress the A-CNT forest to the final volume fraction. By varying the inter-CNT distance in densification process, A-CNT forests with different Vfs

**Figure 2.** Electrochemical performance of supercapacitors with A-CNT electrodes and EMI-BF4/PC electrolyte at 4 V: (a) gravimetric cyclic voltammograms and (b) volumetric cyclic voltammograms of A-CNT electrodes with 1% and

The A-CNT forests were used as the electrodes of the supercapacitors. 3 M EMI-BF4 (1-ethyl-3 methylimidazolium tetrafluoroborate) in propylene carbonate (PC) was used as the electrolyte to improve ionic conductivity. Compared with pure ionic liquid, the ionic liquids/molecular liquids (IL/ML) mixture will have higher conductivity. Polypropylene porous membrane (Celgard 3501, Celgrad LLC) with 25 μm thickness was used as the separator. The sandwich architecture (electrode/separator/electrode) was placed between two pieces of Au-coated steel plates, which served as the current collectors. Then, the whole cell was housed in a Teflon

–1010 CNTs cm−2.

carbon nanotube forests have a 1% volume fraction (Vf) with density of 109

of the supercapacitors.

6 Supercapacitor Design and Applications

can be obtained.

40% Vf at 100 mV s−1.

**Figure 2(a)** presents the cyclic voltammetry (CV) curves at 100 mV s−1 scan rate of supercapacitors with 1% Vf and 40% Vf of A-CNTs as electrodes in 3 M EMI-BF4/PC electrolyte. As shown in the figure, there is very little difference in the gravimetric capacitance drop as the A-CNT Vf increases from 1% to 40%, showing that the densification process can still maintains the aligned nanomorphology with parallel ion pathways. On the other hand, increasing the A-CNT Vf from 1% to 40% leads to a large increase considering the volumetric capacitance performance, as shown in **Figure 2(b)**.

**Figure 3.** (a) Galvanostatic charge and discharge curves at 1 A g−1 for supercapacitors with electrodes of 1% and 40% Vf forests and EMI-BF4/PC as electrolyte. (b) Gravimetric and (c) volumetric specific capacitances for the supercapacitors with electrodes of 1% and 40% Vf forests with different discharge rates. (d) Gravimetric and (e) volumetric Ragone plots for the supercapacitors with electrodes of 1% and 40% Vf A-CNTs. (f) Cycle retention performance of 40% Vf A-CNT supercapacitor with a voltage of 4 V under the charge and discharge current density of 5 A g−1.

The galvanostatic cycles between 0 and 4 V for the supercapacitors with 1% Vf and 40% VfA-CNT electrodes under 1 A g−1 current density are presented in **Figure 3(a)**. The capacitance of the cell can be determined:

$$\mathbf{C} = \mathbf{I} / \left(dV / dt\right) \tag{1}$$

where *I* is the constant current density, *V* is potential, and *t* is discharge time. **Figure 3(b)** and **3(c)** shows the specific gravimetric and volumetric capacitances with different discharge currents, respectively. Although the cell with electrodes of 1% Vf show a little higher specific gravimetric capacitance which is larger than 270 F g−1, their volumetric capacitance is very low (around 3 F cm−3).

It should be mentioned that the 1% Vf A-CNT electrode has a very low active material density with 0.013 g/cm3 . For this kind of electrode, the majority of the electrode volume is filled with electrolytes, whose mass is not included when evaluating the gravimetric electrochemical performance. Instead, the volumetric values should be utilized to investigate when comparing electrodes with large difference about the active material density. It could be found that the specific volumetric capacitance of 40% Vf A-CNT electrodes is about 40 times higher than that of 1% Vf A-CNT electrodes, exhibiting that the nanomorphology of the A-CNTs is preserved by utilizing the mechanical densification method developed here.

The maximum power density and energy density of the supercapacitor cells can be calculated based on the equivalent series resistance (ESR) and the specific capacitance. The gravimetric and volumetric Ragone plots for the supercapacitor cells are shown in **Figure 3(d)** and **3(e)**, respectively. The energy of the supercapacitor cell under each discharge current can be calculated by integrating the discharge curves with time.

$$E = \begin{bmatrix} IV\begin{pmatrix} t \end{pmatrix} dt \end{pmatrix} \tag{2}$$

The ESR can be calculated based on the equation below:

$$ESR = \Delta V \,/\,\Delta I \,\tag{3}$$

where ∆*V* is the voltage drop as the current is switched from a positive value to a negative value, such as from 1 Ag−1 to −1 Ag−1 (Δ*I* = 2 Ag−1). The maximum power density, hence, can be deduced from:

$$P = V^2 / \left(4 \times ESR\right) \tag{4}$$

The active material density of the electrodes in many recently developed nanoporous electrodes, such as the A-CNT electrodes described here, can vary over a broad range. The traditional method to evaluate the performance of supercapacitor cell such as the gravimetric energy and power densities will not accurately reflect the device performance because the only mass of the conductive electrode material is included and the active material density is usually very low. For example, due to a very large "empty space" (80 nm between nanotubes) in the ion pathway, a very high power density of 100 kW kg−1 was obtained in the supercapacitor cell with 1% Vf A-CNT forests as the cell electrodes. As a comparison, 50 kW kg−1 was obtained in the supercapacitor cell based on 40% Vf A-CNT forests as electrodes due to the smaller pore size. However, the supercapacitor cell with 1% Vf A-CNT forests shows a very low gravimetric maximum power and energy density compared with that of the supercapacitor cell based on 40% Vf A-CNT forests if the total electrode mass, including all the elements such as the active materials and the electrolytes, is considered.

where *I* is the constant current density, *V* is potential, and *t* is discharge time. **Figure 3(b)** and **3(c)** shows the specific gravimetric and volumetric capacitances with different discharge currents, respectively. Although the cell with electrodes of 1% Vf show a little higher specific gravimetric capacitance which is larger than 270 F g−1, their volumetric capacitance is very low

It should be mentioned that the 1% Vf A-CNT electrode has a very low active material density

electrolytes, whose mass is not included when evaluating the gravimetric electrochemical performance. Instead, the volumetric values should be utilized to investigate when comparing electrodes with large difference about the active material density. It could be found that the specific volumetric capacitance of 40% Vf A-CNT electrodes is about 40 times higher than that of 1% Vf A-CNT electrodes, exhibiting that the nanomorphology of the A-CNTs is preserved

The maximum power density and energy density of the supercapacitor cells can be calculated based on the equivalent series resistance (ESR) and the specific capacitance. The gravimetric and volumetric Ragone plots for the supercapacitor cells are shown in **Figure 3(d)** and **3(e)**, respectively. The energy of the supercapacitor cell under each discharge current can be

where ∆*V* is the voltage drop as the current is switched from a positive value to a negative value, such as from 1 Ag−1 to −1 Ag−1 (Δ*I* = 2 Ag−1). The maximum power density, hence, can be

The active material density of the electrodes in many recently developed nanoporous electrodes, such as the A-CNT electrodes described here, can vary over a broad range. The traditional method to evaluate the performance of supercapacitor cell such as the gravimetric energy and power densities will not accurately reflect the device performance because the only mass of the conductive electrode material is included and the active material density is usually very low. For example, due to a very large "empty space" (80 nm between nanotubes) in the ion pathway, a very high power density of 100 kW kg−1 was obtained in the supercapacitor cell with 1% Vf A-CNT forests as the cell electrodes. As a comparison, 50 kW kg−1 was obtained in the supercapacitor cell based on 40% Vf A-CNT forests as electrodes due to the smaller pore

by utilizing the mechanical densification method developed here.

calculated by integrating the discharge curves with time.

The ESR can be calculated based on the equation below:

. For this kind of electrode, the majority of the electrode volume is filled with

*E IV t dt* = ò ( ) (2)

*ESR V I* =D D/ (3)

*P V ESR* ( ) <sup>2</sup> = ´ / 4 (4)

(around 3 F cm−3).

8 Supercapacitor Design and Applications

with 0.013 g/cm3

deduced from:

For the supercapacitor cells with 1% Vf of A-CNT forests as electrodes, the specific gravimetric capacitance will decrease to 4.3 Fg−1 when the total electrode mass, including both the 1% Vf active material (A-CNTs) and 99% Vf electrolyte, is used for the calculation. This is much smaller than 270 F g−1 calculated when only the mass of the A-CNTs is included. In contrast, for the 40% Vf A-CNTs, a gravimetric capacitance of 270 F g−1 for the active material alone is equivalent to a 139.8 F g−1 when all the electrode mass is included.

The supercapacitors based on 40% Vf A-CNTs also exhibit an excellent cycling life as shown in **Figure 3(f)**. The data were acquired over 5000 cycles by repeating the galvanostatic charge and discharge process between 0 and 4 V under an alternate current densities of 5 and −5 A g −1, which show an excellent electrochemical stability. Capacitance retention of 98% after 5000 cycles was obtained based on the supercapacitor cell with the ultra-high density A-CNTs as 40% Vf.

**Figure 4.** (a) Nyquist plots of supercapacitors based on A-CNTs with the two volume fractions in the range of 100 kHz to 10 mHz. (b) Gravimetric and (c) volumetric Ragone plots of supercapacitor cells with electrodes based on 40% Vf A-CNT forests and with electrodes using activated carbon. (d) Under several discharge current densities, the relationship of specific capacitance of cells based on 40%Vf A-CNT forests.

The electrochemical impedance spectroscopy (EIS) analysis was carried out to investigate the possible influence of the ultra-high density A-CNT forests on the electrochemical performance of supercapacitor electrodes. **Figure 4(a)** shows the EIS figure in the frequency range of 100 kHz–10 mHz. The Nyquist plots of supercapacitor cells with electrodes based on 1% and 40% Vf A-CNT forests show both semicircles in the high/middle frequency and sharp rises at low frequency range. The semicircle behavior is due to the charge transfer resistance of the electrodes, and the sharp increase at low frequency range is resulted from the ideal capacitive performance of the electrode. It could be found that the cell with electrodes based on 40% Vf A-CNT forests also show much smaller resistance (Z') when normalized with the area of the capacitors (Ω cm−2) compared with that of electrodes based on 1% Vf A-CNT forests.

As a comparison, activated carbon with the thickness of 800 μm was fabricated. The maximum power density, energy density, and electrochemical performances of cells based on activated carbon are shown in **Figure 4(b)** and **4(c)**. A much lower volumetric energy density and power density (20 and 1.1 kWL−1, under 4 V) were obtained based on the activated carbon electrodes, compared with the performance with 40% Vf of A-CNT forests (75 and >25 kW L−1 under 4 V). These results indicate the superior electrochemical performance of the ultra-high-density A-CNT electrodes fabricated from the mechanical densification method developed here. The nanomorphology of the aligned ion pathways leads to the fast charge/discharge rate and high power/energy densities.

**Figure 4(b)** and **4(c)** presents that there is very little increase in the power density while increasing the voltage from 3 to 4 V leads to a large increase in the energy density, from 15 to 75 Wh L−1 (gravimetric 30–150 Wh kg−1). The large increase in the energy density results from the increase in the specific capacitance of the electrodes with the increasing of voltage, as shown in **Figure 4(d)**. The specific capacitance is obtained as 260 F g−1 (135 F cm−3) at 4 V. These values are much higher than those reported earlier for the supercapacitors utilizing densified A-CNTs through the liquid collapsing method. On the other hand, it is noted that the maximum power density depends on the ESR (see Equation (4)), as well as the applied voltage. The results exhibit that there is a large increase in the ESR as the operation voltage has increased from 3 to 4 V, which is consistent with the results of an earlier study in our group [5]. Diffusion process and drifting process have dominated the transport of mobile ion in ionic devices such as supercapacitors. Diffusion is relatively slower and hence represents much higher ESR compared with drifting process.

## **3. Asymmetric supercapacitor with high electrochemical performance**

In this work, an asymmetric supercapacitor, exploiting nm-scale conformal coating of a conducting polymer, poly(ethylenedioxythiophene) (PEDOT) on aligned carbon nanotubes (A-CNTs) as one electrode and an ultra-high density aligned graphene sheets (a-graphene) as the other electrode, has been developed. The asymmetric configuration of the supercapacitor allows both electrodes to be separately tailored, increasing device capacitance and the electrochemical window, and thereby operating voltage. As a result of complementary three-dimensional nanotailoring of the asymmetric electrodes, the device exhibits a wide 4V electrochemical window and high electrochemical performance [6].

The electrochemical impedance spectroscopy (EIS) analysis was carried out to investigate the possible influence of the ultra-high density A-CNT forests on the electrochemical performance of supercapacitor electrodes. **Figure 4(a)** shows the EIS figure in the frequency range of 100 kHz–10 mHz. The Nyquist plots of supercapacitor cells with electrodes based on 1% and 40% Vf A-CNT forests show both semicircles in the high/middle frequency and sharp rises at low frequency range. The semicircle behavior is due to the charge transfer resistance of the electrodes, and the sharp increase at low frequency range is resulted from the ideal capacitive performance of the electrode. It could be found that the cell with electrodes based on 40% Vf A-CNT forests also show much smaller resistance (Z') when normalized with the area of the

capacitors (Ω cm−2) compared with that of electrodes based on 1% Vf A-CNT forests.

power/energy densities.

10 Supercapacitor Design and Applications

drifting process.

As a comparison, activated carbon with the thickness of 800 μm was fabricated. The maximum power density, energy density, and electrochemical performances of cells based on activated carbon are shown in **Figure 4(b)** and **4(c)**. A much lower volumetric energy density and power density (20 and 1.1 kWL−1, under 4 V) were obtained based on the activated carbon electrodes, compared with the performance with 40% Vf of A-CNT forests (75 and >25 kW L−1 under 4 V). These results indicate the superior electrochemical performance of the ultra-high-density A-CNT electrodes fabricated from the mechanical densification method developed here. The nanomorphology of the aligned ion pathways leads to the fast charge/discharge rate and high

**Figure 4(b)** and **4(c)** presents that there is very little increase in the power density while increasing the voltage from 3 to 4 V leads to a large increase in the energy density, from 15 to 75 Wh L−1 (gravimetric 30–150 Wh kg−1). The large increase in the energy density results from the increase in the specific capacitance of the electrodes with the increasing of voltage, as shown in **Figure 4(d)**. The specific capacitance is obtained as 260 F g−1 (135 F cm−3) at 4 V. These values are much higher than those reported earlier for the supercapacitors utilizing densified A-CNTs through the liquid collapsing method. On the other hand, it is noted that the maximum power density depends on the ESR (see Equation (4)), as well as the applied voltage. The results exhibit that there is a large increase in the ESR as the operation voltage has increased from 3 to 4 V, which is consistent with the results of an earlier study in our group [5]. Diffusion process and drifting process have dominated the transport of mobile ion in ionic devices such as supercapacitors. Diffusion is relatively slower and hence represents much higher ESR compared with

**3. Asymmetric supercapacitor with high electrochemical performance**

In this work, an asymmetric supercapacitor, exploiting nm-scale conformal coating of a conducting polymer, poly(ethylenedioxythiophene) (PEDOT) on aligned carbon nanotubes (A-CNTs) as one electrode and an ultra-high density aligned graphene sheets (a-graphene) as the other electrode, has been developed. The asymmetric configuration of the supercapacitor allows both electrodes to be separately tailored, increasing device capacitance and the electrochemical window, and thereby operating voltage. As a result of complementary For supercapacitors, it is well known that the energy density (*E*) is related to the gravimetric or volumetric cell capacitance (*C*) and operation voltage (*V*), i.e.

$$E = \frac{1}{2}CV^2\tag{5}$$

And the maximum power density P is determined by Equation (4). Equations (4) and (5) have shown that one of the most effective ways to increase both the power and energy densities is to raise the cell operation voltage. In general, the operation voltage of supercapacitor cell has the relationship with the electrochemical window, which is determined by the interface between the electrode and electrolyte. As a promising way, asymmetric supercapacitor can be assembled to make full use of the electrochemical windows of both electrodes to increase the maximum cell operation voltage in the supercapacitor cell. Morphology control of the electrodes via nanoscale tailoring is shown to be an effective way to increase supercapacitor performance (gravimetric and volumetric power and energy) via increasing ECW and capacitance and reducing ESR.

**Figure 5.** Nanostructured electrodes in asymmetric supercapacitors. Left, Low and high magnification TEM micrographs of the electrode, composed of conformal oCVD PEDOT on A-CNTs, and right, SEM images of activated graphene electrode.

Recent advances in the conformal coating of conducting polymer PEDOT by oxidative chemical vapor deposition (oCVD) onto nanowire arrays and development of graphene with relatively high gravimetric surface area create unique opportunities for developing high performance asymmetric supercapacitors. As schematically illustrated in **Figure 5**, the combination of the aligned ion transport pathways formed by the aligned nanowire arrays that provide fast ion transport in the electrode that reduces ESR of the electrode and the conformal coating of conducting polymer PEDOT on the A-CNTs that enhances the charge storage capability (a large C) contributes to both high energy and power densities of the cells. PEDOT was selected as the conducting polymer because of its environmental stability, high electrical conductivity, and a wide ECW. As shown in **Figure 6(a)**, conformally coated oCVD PEDOT on the A-CNTs yields a stable ECW from −1.0 to 1.8 V, when using an ionic liquid/ molecular liquid mixture, 2 M 1-butyl-3-methylimidazolium tetrafluoroborate (BMI BF4)/ propylene carbonate (PC), as the electrolyte. The high ECW of 1.8 V makes it as an excellent positive electrode material in the asymmetric supercapacitors.

**Figure 6.** Performance of two electrodes: (a) CV curves of PEDOT/A-CNTs composite at 5 mV s−1 in 2 M BMIBF4/PC. (b) Galvanostatic charge/discharge curves of PEDOT/A-CNT composite at current densities of 2 A/g. (c) Specific capacitance at different discharge densities of PEDOT/A-CNT electrode. (d) CV curves of a-graphene at 5 mV/s in 2 M BMIBF4/PC. (e) Galvanostatic charge/discharge curves of a-graphene at current densities of 2 A/g. (f) Specific capacitance at different discharge densities of a-graphene electrode.

Due to their favorable ECWs, carbon-based electrodes such as activated carbon have been used for the negative in the asymmetric supercapacitors. In this study, a new class of carbon material, the a-graphene, was selected due to its superior gravimetric surface area compared with activated carbon. The a-graphene, first reported by Zhu et al., presented a very large specific surface area (as large as 3100 m2 /g) with nanosized pores and demonstrated a gravimetric capacitance as high as 200 F/g when assembling supercapacitors [7]. However, the simple mechanical packing of the a-graphene flakes led to a relatively low density (~0.3 g/cm3 ) compared with the graphite density of 2.2 g/cm3 . A low volumetric efficiency of supercapacitors was obtained by this configuration (the specific volumetric capacitance was 60 F/cm3 ). When randomly packing these a-graphene flakes, which have lateral dimension of a few microns and a thickness of a few nanometers, it is inevitable to include micron-sized pores in the electrodes, reducing the density. Self-assembly processes are quite effective in increasing the density of graphene-based materials. Here, by employing a vacuum-assisted self-assembly process, which enabled a-graphene flakes aligned in parallel and stacked successively on top of each other, as shown by the SEM image of **Figure 5**, we fabricated the a-graphene electrodes with high density while preserving the nanoporous morphology of each a-graphene flake. The ECW of the a-graphene was also characterized, and as presented in **Figure 6(d)**, the a-graphene has a stable ECW from −2.2 V to 1 V with an electrolyte of 2M BMIBF4/PC. The combination of high specific gravimetric surface area and high density of the a-graphenes as the negative electrode increases the ECW and results in high volumetric power and energy densities, besides the long cycle lifetime and high capacitive retention.

electrical conductivity, and a wide ECW. As shown in **Figure 6(a)**, conformally coated oCVD PEDOT on the A-CNTs yields a stable ECW from −1.0 to 1.8 V, when using an ionic liquid/ molecular liquid mixture, 2 M 1-butyl-3-methylimidazolium tetrafluoroborate (BMI BF4)/ propylene carbonate (PC), as the electrolyte. The high ECW of 1.8 V makes it as an excellent

**Figure 6.** Performance of two electrodes: (a) CV curves of PEDOT/A-CNTs composite at 5 mV s−1 in 2 M BMIBF4/PC. (b) Galvanostatic charge/discharge curves of PEDOT/A-CNT composite at current densities of 2 A/g. (c) Specific capacitance at different discharge densities of PEDOT/A-CNT electrode. (d) CV curves of a-graphene at 5 mV/s in 2 M BMIBF4/PC. (e) Galvanostatic charge/discharge curves of a-graphene at current densities of 2 A/g. (f) Specific capaci-

Due to their favorable ECWs, carbon-based electrodes such as activated carbon have been used for the negative in the asymmetric supercapacitors. In this study, a new class of carbon material, the a-graphene, was selected due to its superior gravimetric surface area compared with activated carbon. The a-graphene, first reported by Zhu et al., presented a very large specific

tance at different discharge densities of a-graphene electrode.

positive electrode material in the asymmetric supercapacitors.

12 Supercapacitor Design and Applications

The high density (through packing) of aligned carbon nanotube (A-CNT) forests are distinctively advantageous as the conductive networks to support the CP coating layer in supercapacitors, when we compare with randomly packed morphologies. Besides the direct (and thereby fast) ion transport in aligned channel to reduce ESR illustrated in **Figure 5**, the PEDOT/ A-CNT composite also provides better mechanical stability and hence higher retention property after many charge and discharge cycles, compared with the electrodes of the PEDOT/ randomly packed CNT networks. In the extant literature, most electrodes of CP/CNTs were fabricated by electrochemical methods, which will lead to nonuniform CP layers coating on the CNTs. As shown in **Figure 6**, the thin (~5 to 10 nm) oxidative chemical vapor deposition (oCVD), PEDOT layers form a conformal coating on very high aspect ratio A-CNTs (0.2 mm long).

The electrochemical performance of PEDOT/A-CNT composite electrode were investigated by cyclic voltammetry (CV) and galvanostatic charge-discharge tests using a screen-printed electrode system (Dropsens) with the PEDOT/A-CNT composite as the working electrode, while Ag and Pt were employed as the reference and counter electrodes, respectively. **Figure 6(a)** shows a CV curve of the PEDOT/A-CNT forest composite electrode in 2 M BMIBF4/PC electrolyte under a scan rate of 5 mV s−1, which shows an ECW from −1 V to +1.8 V. The galvanostatic cycles for the PEDOT/A-CNT electrode at the alternate currents of 2 and −2A g−1 are shown in **Figure 6(b)**. The symmetric and linear charge and discharge characteristics with time reveal a rapid I-V response and reversible electrochemical reaction, leading to an superior capacitive behavior. The specific capacitance of the electrode can be determined from Equation (1). A high specific gravimetric capacitance of 230 F/g was obtained at 2 A/g. **Figure 6(c)** presents the specific capacitance at different discharge current density, from 0.5 to 10 A/g. Capacitance retention of 74.2% was obtained at 10 A/g (from 260.8 F/g at 0.5 A/g to 193.5 F/g at 10 A/g), indicating that the PEDOT/A-CNT electrode provide reliable capacitive performance for high power applications. This relatively high retention mechanistically arises from the conformal coating of oCVD PEDOT on A-CNTs. The parallel ion transport pathways formed by the PEDOT/A-CNTs and the high electronic conductivity of the A-CNTs improve the ion transport and result in low ESR and therefore high power density. The cycling stability of the PEDOT/A-CNT electrodes was characterized and compared with that of the electrodes of PEDOT deposited on randomly packed CNT networks. Symmetric supercapacitors made of the PEDOT/A-CNTs had a retention of 89% after 1000 cycles of 2 V voltage cycle compared with a retention of 73% after 1000 cycles from PEDOT on random CNT morphologies. In the randomly packed CNT networks, there are CP layers in the gaps between nanotubes. The mechanical failure of CP layers in these gaps will cause disruption of the electric conduction paths between nanotubes and reduce the conductivity of CNT networks after long charge/ discharge cycles. As a result, the capacitance and other electrochemical performances will be influenced. In contrast, the electric conduction path of the continuously aligned CNT forests would not be disrupted by the mechanical failure of the CP coating layers due to this nanomorphology. Hence, the A-PEDOT/A-CNT electrodes exhibit more robust mechanical stability and high retention of the capacitance, compared with the electrodes of the CP deposited on randomly packed CNT networks.

The electrochemical performance of the a-graphene electrode was characterized as above, including using 2 M BMIBF4/PC as the electrolyte. **Figure 6(d)** presents a CV curve of the agraphene electrode at a scan rate of 5 mV/s, showing an ECW of −2.2 V to +1 V. The slope of the discharge curve in **Figure 6(e)** yields a specific capacitance of 165 F/g at 2 A/g. The specific capacitances of the a-graphene electrode with different discharge currents are presented in **Figure 6(f)**. The a-graphene electrode exhibits high specific capacitance, ranging from 186.4 to 148.2 F/g as the discharge current increases from 0.5 A/g to 10 A/g. Moreover, the high density of the a-graphene electrodes results here in a specific volumetric capacitance with 175 F/cm3 from the discharge curve of constant current of 1 A/g, which is the highest among all the carbonbased electrodes.

Both electrodes are independently tailored in asymmetric supercapacitors to operate under more optimal conditions. Here, the PEDOT/A-CNTs electrode and a-graphene electrode were assembled, separated by a porous paper of 40 μm thick. 2 M BMIBF4/PC was used as the electrolyte due to its high ionic conductivity. By properly tuning the mass ratio of the two electrodes, the asymmetric capacitor can be operated at the full 4 V cell operation voltage, reaching the maximum voltages from both electrodes (=1.8 V (A-CNT/PEDOT composite, V+) + 2.2 V (a-graphene flakes, V−)). Based on the consideration that charge stored at the two electrodes should be equal in magnitude with opposite sign ( q+ <sup>=</sup> q− ), the mass ratio between the two electrodes can be determined as following equation if we consider the stored charge q at the electrode is *q* = *C*∆*Vm*, where *C* is the specific gravimetric capacitance, ∆*V* is the maximum potential range allowed by the ECW, and *m* is the mass of the electrode.

m CV m CV + -- - ++ <sup>D</sup> <sup>=</sup> <sup>D</sup> (6)

High Volumetric Performance Supercapacitors with Controlled Nanomorphology http://dx.doi.org/10.5772/65186 15

**Figure 7.** Cell performance: (a) CV curves of asymmetric cell under different scan rates from 5 to 100 mV/s between 0 and 4 V using 2 M BMIBF4/PC as the electrolyte. (b) Galvanostatic charge/discharge curves of asymmetric cell under a current density of 2 A/g. (c) Cell capacitances of asymmetric cell at different discharge current densities. (d) Cycle capacitance retention of asymmetric supercapacitor under a voltage of 4 V at a current density of 5 A/g in 2 M BMIBF4/PC electrolyte.

From the specific capacitances of the two electrodes, 230 and 165 F/g, respectively, under a constant discharge current of 2 A/g, and ΔV+ =1.8 V and ΔV− = −2.2 V, Equation (6) can lead to the mass ratio (m+/m−) of 0.88, for a full 4 volts cell operation voltage, providing the design

characteristics of the asymmetric supercapacitor assembled here.

performance for high power applications. This relatively high retention mechanistically arises from the conformal coating of oCVD PEDOT on A-CNTs. The parallel ion transport pathways formed by the PEDOT/A-CNTs and the high electronic conductivity of the A-CNTs improve the ion transport and result in low ESR and therefore high power density. The cycling stability of the PEDOT/A-CNT electrodes was characterized and compared with that of the electrodes of PEDOT deposited on randomly packed CNT networks. Symmetric supercapacitors made of the PEDOT/A-CNTs had a retention of 89% after 1000 cycles of 2 V voltage cycle compared with a retention of 73% after 1000 cycles from PEDOT on random CNT morphologies. In the randomly packed CNT networks, there are CP layers in the gaps between nanotubes. The mechanical failure of CP layers in these gaps will cause disruption of the electric conduction paths between nanotubes and reduce the conductivity of CNT networks after long charge/ discharge cycles. As a result, the capacitance and other electrochemical performances will be influenced. In contrast, the electric conduction path of the continuously aligned CNT forests would not be disrupted by the mechanical failure of the CP coating layers due to this nanomorphology. Hence, the A-PEDOT/A-CNT electrodes exhibit more robust mechanical stability and high retention of the capacitance, compared with the electrodes of the CP deposited on

The electrochemical performance of the a-graphene electrode was characterized as above, including using 2 M BMIBF4/PC as the electrolyte. **Figure 6(d)** presents a CV curve of the agraphene electrode at a scan rate of 5 mV/s, showing an ECW of −2.2 V to +1 V. The slope of the discharge curve in **Figure 6(e)** yields a specific capacitance of 165 F/g at 2 A/g. The specific capacitances of the a-graphene electrode with different discharge currents are presented in **Figure 6(f)**. The a-graphene electrode exhibits high specific capacitance, ranging from 186.4 to 148.2 F/g as the discharge current increases from 0.5 A/g to 10 A/g. Moreover, the high density of the a-graphene electrodes results here in a specific volumetric capacitance with 175 F/cm3 from the discharge curve of constant current of 1 A/g, which is the highest among all the carbon-

Both electrodes are independently tailored in asymmetric supercapacitors to operate under more optimal conditions. Here, the PEDOT/A-CNTs electrode and a-graphene electrode were assembled, separated by a porous paper of 40 μm thick. 2 M BMIBF4/PC was used as the electrolyte due to its high ionic conductivity. By properly tuning the mass ratio of the two electrodes, the asymmetric capacitor can be operated at the full 4 V cell operation voltage, reaching the maximum voltages from both electrodes (=1.8 V (A-CNT/PEDOT composite, V+) + 2.2 V (a-graphene flakes, V−)). Based on the consideration that charge stored at the two electrodes should be equal in magnitude with opposite sign ( q+ <sup>=</sup> q− ), the mass ratio between the two electrodes can be determined as following equation if we consider the stored charge q at the electrode is *q* = *C*∆*Vm*, where *C* is the specific gravimetric capacitance, ∆*V* is the

maximum potential range allowed by the ECW, and *m* is the mass of the electrode.

m CV m CV + -- - ++

<sup>D</sup> <sup>=</sup> <sup>D</sup> (6)

randomly packed CNT networks.

14 Supercapacitor Design and Applications

based electrodes.

**Figure 7(a)** presents the CV curves of the fabricated asymmetric supercapacitors at scan rates from 5 to 100 mV s−1 using the 2 M BMIBF4/PC as electrolyte. The capacitors display near rectangular CV curves, especially for the lower scan rates. In order to evaluate the capacitive performance of the cell further, galvanostatic charge/discharge curves at different current densities were characterized. The galvanostatic cycles at alternate charge/discharge current densities of 2 and –2 A/g are presented in **Figure 7(b)**, from which the cell capacitance was determined (Equation (1)). **Figure 7(c)** presents the cell gravimetric and volumetric capacitances at different discharge currents. It should be noted that the calculated cell capacitance was based on the total mass of the active materials (both positive and negative electrodes) because it is not meaningful to deduce the specific capacitance of a single electrode for the asymmetric supercapacitor. The cell capacitance is 81.6 F/g at 0.2 A/g and becomes 55.4 F/g as the discharge current density increases to 10 A/g, indicating relatively good capacitance retention. The cell capacitance obtained here is higher than that of a-graphene-based symmetric supercapacitors and other conducting polymer-based asymmetric supercapacitors. Cycling retention performance of the asymmetric supercapacitors was investigated by continuously performing the galvanostatic charge/discharge process between 0 and 4 V at the alternative current densities of 5 and −5 A/g for more than 1000 cycles. The asymmetric supercapacitor maintains electrochemical retention of 94% after 1000 cycles as presented in **Figure 7(d)**. The small capacitance loss is likely attributable to capacitance decay of the PEDOT/A-CNT electrode.

**Figure 8.** Asymmetric cell absolute and relative performance: (a) Nyquist plot of the asymmetric cell. (b) Internal resistance drop at different current densities. (c) Ragone plot of PEDOT/A-CNTs//a-graphene asymmetric supercapacitor in gravimetric unit. (d) Ragone plot of PEDOT/A-CNTs//a-graphene asymmetric supercapacitor in volumetric unit.

The electrochemical performance of the asymmetric supercapacitor cell was further characterized by electrochemical impedance spectroscopy (EIS). Nyquist plot, as shown in **Figure 8(a)**, is achieved in the frequency range of 100 kHz–10 mHz of 5 mV applied voltage. It could be found that a semicircle in the high-frequency region and a sharp rise of the imaginary part of the electric impedance shown in the figure reflecting the dominance of the cell capacitance in the-low frequency region. As shown in the Nyquist plot, the semicircle at the middle/high frequency is due to the charge transfer resistance in the porous electrodes. The high-frequency intersection on the real axis of the Nyquist plot shown in the figure represents the internal resistances. Internal resistance of 0.1 Ω cm2 of the cell is obtained in the figure when normalized with the area of the current collector of the capacitors indicating a high electrical conductivity and low ESR of the cells.

The maximum power density of the asymmetric supercapacitor cell is determined from Equation (4). **Figure 8(c)** presents the Ragone plot (gravimetric power density versus energy density) of the asymmetric supercapacitor derived from the galvanostatic discharge curves measured at different charge/discharge current densities following standard practice. In addition to gravimetric performance, the maximum power density is derived from Equation (3), where *V* is the operation voltage, which is 4 V here.

The volumetric energy and power densities are more important in practical applications. The cells exhibit both high volumetric and gravimetric power and energy density at 149 kW/L (233 kW/kg) and 113.2 Wh/L (176.6 Wh/kg), respectively. These values are significantly higher than those of other reported carbon-based symmetric supercapacitors, conducting polymer-based supercapacitors and other devices reported previously [8–10]. The pseudocapacitor nature of PEDOT has lower charge/discharge speed compared with that of the pure EDLC supercapacitors and hence the power density is lower than that of the A-aMEGO supercapacitors, which have been presented in the proceeding section.

In summary, an asymmetric supercapacitor, employing the conformal coating of PEDOT/A-CNT composite as one electrode and high density a-graphene flakes as the other electrode, has been developed in this paper. PEDOT/A-CNT composite combines fast ion transport pathways, enhances charge storage capability, and reduces ESR while a-graphene electrode fabricated from a self-assembly process, which possesses exceedingly high specific gravimetric and volumetric capacitance. The two electrodes are individually tailored to control the nanomorphology and work synergistically together in the asymmetric cell configuration. The ECW has been expended up to 4 V. Tailoring of the two electrodes materials at a scale approaching that of the ions can allow asymmetric supercapacitor performance to be further expanded to meet the requirement of a broad range of energy storage applications.

## **Acknowledgements**

galvanostatic charge/discharge process between 0 and 4 V at the alternative current densities of 5 and −5 A/g for more than 1000 cycles. The asymmetric supercapacitor maintains electrochemical retention of 94% after 1000 cycles as presented in **Figure 7(d)**. The small capacitance

**Figure 8.** Asymmetric cell absolute and relative performance: (a) Nyquist plot of the asymmetric cell. (b) Internal resistance drop at different current densities. (c) Ragone plot of PEDOT/A-CNTs//a-graphene asymmetric supercapacitor in gravimetric unit. (d) Ragone plot of PEDOT/A-CNTs//a-graphene asymmetric supercapacitor in volumetric unit.

The electrochemical performance of the asymmetric supercapacitor cell was further characterized by electrochemical impedance spectroscopy (EIS). Nyquist plot, as shown in **Figure 8(a)**, is achieved in the frequency range of 100 kHz–10 mHz of 5 mV applied voltage. It could be found that a semicircle in the high-frequency region and a sharp rise of the imaginary part of the electric impedance shown in the figure reflecting the dominance of the cell capacitance in the-low frequency region. As shown in the Nyquist plot, the semicircle at the middle/high frequency is due to the charge transfer resistance in the porous electrodes. The high-frequency intersection on the real axis of the Nyquist plot shown in the figure represents the internal resistances. Internal resistance of 0.1 Ω cm2 of the cell is obtained in the figure when normalized with the area of the current collector of the capacitors indicating a high electrical conduc-

The maximum power density of the asymmetric supercapacitor cell is determined from Equation (4). **Figure 8(c)** presents the Ragone plot (gravimetric power density versus energy density) of the asymmetric supercapacitor derived from the galvanostatic discharge curves measured at different charge/discharge current densities following standard practice. In addition to gravimetric performance, the maximum power density is derived from Equation

tivity and low ESR of the cells.

16 Supercapacitor Design and Applications

(3), where *V* is the operation voltage, which is 4 V here.

loss is likely attributable to capacitance decay of the PEDOT/A-CNT electrode.

This work was supported by AFOSR under Grant no. FA9550-11-1-0192. Reproduced from Ref. [4] with permission from ELSEVIER. Reproduced from Ref. [6] with permission from the Royal Society of Chemistry.

## **Author details**

Yue Zhou\* and Qiming Zhang

\*Address all correspondence to: yuezhou26@gmail.com

Department of Electrical Engineering, Pennsylvania State University, University Park, Pennsylvania, USA

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#### **Electrochemical Materials Design for Micro-Supercapacitors Electrochemical Materials Design for Micro-Supercapacitors**

Can Liu and Zhengjun Zhang Can Liu and Zhengjun Zhang

Additional information is available at the end of the chapter Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/64986

#### **Abstract**

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Micro–supercapacitors (m–SC) arise from the demand of developing micro–power system for MEMS devices, attracting much research interest in recent years. As m–SC has to achieve high areal energy and power densities, the volumetric capacitance and the rate capability of the electrode materials have become the most important concern. This review compares the intrinsic electrochemical properties of the state–of–art electrode materials for m–SC, reporting the recent advances in the three types of electrode materials. For carbon electrode materials, two developing trends are identified: one is to enhance volumetric capacitance through a proper film fabrication process, while the other one is to further promote its fast response rate by making open– structured devices. For pseudocapacitive oxides, in order to achieve better rate capability and cyclability, the relationship between the electrochemical property and the structure is worth further exploration. As an example, the composition, microstructure, and morphology of the molybdenum oxide film were optimized to realize superior electrochemical performance as an electrode material for m–SC. Architecture design is another important factor for m–SC. In–plane interdigital architectures have proven its success to fabricate fast response devices. Further study on the interplay effect between such architecture and pseudocapacitive materials is in need.

**Keywords:** micro-supercapacitors, electrode materials, EDLC, pseudocapacitance

## **1. Introduction**

Since the microelectromechanical systems (MEMS) develop rapidly toward standalone microsensors, actuators, and various functional devices, the design of power supply has received more and more research interests [1]. The conventional bulky batteries severely limit

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

the advantages of these smart systems, a micropower system, i.e., generating power directly from microstructures, thus has to be developed [2]. After the earliest explorations on the mirointernal-combustion engine which requires complicated micromachining processes and high manufacturing cost [2–4], researchers turn to microbattery system which is potentially of low cost and high capacity and more desirable for MEMS devices [1]. A complete micropower system should consist of energy conversion and storage units integrated on chip. The energy conversion devices include microscaled fuel cells and solar cells, while the energy storage devices mainly refer to rechargeable microbatteries, which have been remarkably advanced under many research efforts [5]. Microbatteries, or thin-film batteries, have become commercially available with a rapid expanding market. Nevertheless, similar to the features of macroscaled batteries, the shortages of microbatteries are limited lifetime and low power density, which bring economic and environmental challenges to systems that they power.

Micro-supercapacitors (m-SC) appeared later as another important energy storage unit. Also known as an electrochemical capacitor (EC), a supercapacitor works through the accumulation of the electrostatic charge within an electrochemical double layer at the electrode/electrolyte interface. It could present high specific capacitance, which mainly depends on the high surface area of the electrode materials, or some pseudo-Faradaic charge transfer process. Compared with the batteries, supercapacitors possess inferior energy density but superior power density, i.e., they can be fully charged/discharged in seconds or minutes. Another prominent advantage of supercapacitors is the long cycle life, which is rather comparable with that of the functional devices. For the large-scale application system, supercapacitors are usually used combined with batteries so that both the high energy density of the batteries and the high power density of superapacitors could be utilized to ensure sufficient power supply. Similarly, m-SC is also complementary but indispensable when high power density is required to support the MEMS devices. More importantly, it could even replace the microbatteries when the cycle life of the device is preferred over the energy density for the whole system. As a matter of fact, with the great development of energy harvesters and nanogenerators, i.e., microscaled energy conversion devices that harvest energy from the ambient environment such as solar power, wind, water flow, vibrational energy, and thermal energy from waste heat, m-SC has been much more competitive as an alternative to batteries to play the role of energy storage in the self-powered micro/nanosystems [6]. **Figure 1** illustrates such a sustainable self-powered system, which consists of five different modules, namely energy harvester, energy storage, sensor, data processor/controller, and data transmitter/receiver [7].

M-SC is originally targeted at high power delivery and robust cyclability, and thus, carbonbased electrode materials are first employed to design on-chip electrochemical double-layer capacitors (EDLC), but the capacitance is relatively low [8–10]. The research interest grows quickly since 2010, when several works were reported on the improved design of the carbon electrode materials, especially the carbide-derived carbon film that possesses high volumetric capacitance and compatible with microfabrication [11–13]. Thanks to the fruitful development of the conventional supercapacitors and nanomaterials, m-SCs also received rapid advances when more research groups turn their interest to the on-chip devices, employing various nanomaterials and designing different fabrication protocols even without being limited to conventional MEMS fabrication routes [14–19].

the advantages of these smart systems, a micropower system, i.e., generating power directly from microstructures, thus has to be developed [2]. After the earliest explorations on the mirointernal-combustion engine which requires complicated micromachining processes and high manufacturing cost [2–4], researchers turn to microbattery system which is potentially of low cost and high capacity and more desirable for MEMS devices [1]. A complete micropower system should consist of energy conversion and storage units integrated on chip. The energy conversion devices include microscaled fuel cells and solar cells, while the energy storage devices mainly refer to rechargeable microbatteries, which have been remarkably advanced under many research efforts [5]. Microbatteries, or thin-film batteries, have become commercially available with a rapid expanding market. Nevertheless, similar to the features of macroscaled batteries, the shortages of microbatteries are limited lifetime and low power density,

20 Supercapacitor Design and Applications

which bring economic and environmental challenges to systems that they power.

processor/controller, and data transmitter/receiver [7].

Micro-supercapacitors (m-SC) appeared later as another important energy storage unit. Also known as an electrochemical capacitor (EC), a supercapacitor works through the accumulation of the electrostatic charge within an electrochemical double layer at the electrode/electrolyte interface. It could present high specific capacitance, which mainly depends on the high surface area of the electrode materials, or some pseudo-Faradaic charge transfer process. Compared with the batteries, supercapacitors possess inferior energy density but superior power density, i.e., they can be fully charged/discharged in seconds or minutes. Another prominent advantage of supercapacitors is the long cycle life, which is rather comparable with that of the functional devices. For the large-scale application system, supercapacitors are usually used combined with batteries so that both the high energy density of the batteries and the high power density of superapacitors could be utilized to ensure sufficient power supply. Similarly, m-SC is also complementary but indispensable when high power density is required to support the MEMS devices. More importantly, it could even replace the microbatteries when the cycle life of the device is preferred over the energy density for the whole system. As a matter of fact, with the great development of energy harvesters and nanogenerators, i.e., microscaled energy conversion devices that harvest energy from the ambient environment such as solar power, wind, water flow, vibrational energy, and thermal energy from waste heat, m-SC has been much more competitive as an alternative to batteries to play the role of energy storage in the self-powered micro/nanosystems [6]. **Figure 1** illustrates such a sustainable self-powered system, which consists of five different modules, namely energy harvester, energy storage, sensor, data

M-SC is originally targeted at high power delivery and robust cyclability, and thus, carbonbased electrode materials are first employed to design on-chip electrochemical double-layer capacitors (EDLC), but the capacitance is relatively low [8–10]. The research interest grows quickly since 2010, when several works were reported on the improved design of the carbon electrode materials, especially the carbide-derived carbon film that possesses high volumetric capacitance and compatible with microfabrication [11–13]. Thanks to the fruitful development of the conventional supercapacitors and nanomaterials, m-SCs also received rapid advances when more research groups turn their interest to the on-chip devices, employing various

**Figure 1.** (a) Schematic diagram of the integrated self-powered system showing five modules: energy harvester, energy storage, sensors, data processor & controller, and data transmitter & receiver. (b) Prototype of an integrated self-powered system using a nanogenerator as the energy harvester and a capacitor as the storage unit [7].

After several years of active research, the term of micro-supercapacitor is formally defined and the performance metrics are well recognized [6]. According to the definition given by Beidaghi and Gogotsi, micro-supercapacitors, or electrochemical micro-capacitors, refer to miniaturized supercapacitors that are designed and fabricated to serve as power sources or energy storage units in microelectronic devices. Due to the purpose of being specifically assembled to microelectronic devices, there comes the confinement in the fabrication methods that should be compatible with the current techniques in the semiconductor industry. Hence, the general appearance of m-SCs is a device taking up a footprint area in the millimeter or centimeter scale and a thickness of less than 10 μm. The configurations include sandwiched assemblies consisting of thin-film electrodes, planar arrays of microelectrodes like interdigital electrodes, and three-dimensional (3D) architectures of nanoscale building blocks [6, 12]. The former two configurations are commonly adopted in the present laboratory prototypes, while the latter one, 3D architecture, is a composed idea of the next generation whose realization still requires innovations.

## **2. Performance metrics for micro-supercapacitors**

As perceived to power the microelectronic devices in a self-powered system, sufficient and energy have to be delivered by the m-SCs. The specific requirement of the power depends on the functional devices varying from environmental sensors to personal electronics, which may be in the range of 1–100 μW. Meanwhile, the duration of the power delivery is also required, meaning certain energy has to be supplied. Thus, power and energy together determine the suitability of an m-SC to power a microelectronic system. It should be mentioned that there is another direction emerging for the application of m-SCs other than micropower units, namely replacing electrolytic capacitors in electronic circuits such as alternating current line filtering. In order to utilize the advantage of miniaturized size of m-SCs to replace the conventional bulky electrolytic capacitors, the m-SCs have to response fast with a relatively large capacitance, i.e., ideal capacitive behavior with a small resistor-capacitor (RC) time constant (e.g., <8.3 ms for ac line filtering). In such case, the capacitance is a critical parameter for the evaluation, and it is also important whether the capacitance is well kept under faster charge/ discharge conditions, which may be termed as rate capability, a terminology from the battery field. In addition, good reversibility (usually assessed by the Coulombic efficiency) and long cycle life should persist for the m-SCs.

As a matter of fact, assuming that an m-SC with a constant capacitance of C is charged from 0 V to an ultimate voltage of U in a duration of *t*, the stored energy *E* is calculated by *E* = 1/2 CU2 . And the power *P* is calculated by *P* = *E*/*t*, which is further written as *P* = 1/2 IU, if the charge current is I. Therefore, the energy depends on the capacitance and the voltage. Meanwhile, the power depends on the voltage and the working current that is chosen for the operation of the device. However, the current is not chosen arbitrarily, as the capacitance usually decreases with the increase of current. In other words, the energy shows a declining relation with the power, which is usually described as a Ragone plot. In a word, the performance of an m-SC device could be represented by a Ragone plot, or alternatively, it could be described more intrinsically by the voltage and the capacitance versus the working current. Since the footprint area and the occupied volume are limited for m-SCs when integrated into the system, normalized parameters, i.e., the areal or volumetric energy and power densities and capacitance, are the most important in evaluating the m-SCs. The device performance is determined by the intrinsic properties of the electrode materials, the electrolyte, and the device architecture, wherein the electrode materials play the most critical role and attract abundant studies. In order to compare the properties of different electrode materials effectively, the volumetric energy and power densities as well as the volumetric capacitance should be assessed. This is different from the performance metrics for conventional macro-supercapacitors, where gravimetric capacitance, energy, and power densities are emphasized. Due to the microscaled size, the weight of the electrode materials becomes almost negligible, while the volumetric parameters are the most important concern in developing practical devices. On the other hand, mass density is not a limiting factor, which remarkably expands the choices of the electrode materials for novel m-SCs.

#### **3. Electrode materials for micro-supercapacitors**

Selection of suitable electrode materials is the kernel in the design of m-SC for certain application. According to the electrochemical energy storage mechanism, the electrode materials currently studied could be divided into three types: carbon-based materials, pseudocapacitive oxides, and conducting polymers. The main purpose of this review is to give a clear comparison among the intrinsic electrochemical properties of the state-of-art electrode materials for m-SC, expecting to open up helpful strategies on developing advanced m-SCs in future.

#### **3.1. Carbon-based materials**

In order to utilize the advantage of miniaturized size of m-SCs to replace the conventional bulky electrolytic capacitors, the m-SCs have to response fast with a relatively large capacitance, i.e., ideal capacitive behavior with a small resistor-capacitor (RC) time constant (e.g., <8.3 ms for ac line filtering). In such case, the capacitance is a critical parameter for the evaluation, and it is also important whether the capacitance is well kept under faster charge/ discharge conditions, which may be termed as rate capability, a terminology from the battery field. In addition, good reversibility (usually assessed by the Coulombic efficiency) and long

As a matter of fact, assuming that an m-SC with a constant capacitance of C is charged from 0 V to an ultimate voltage of U in a duration of *t*, the stored energy *E* is calculated by *E* = 1/2

Selection of suitable electrode materials is the kernel in the design of m-SC for certain application. According to the electrochemical energy storage mechanism, the electrode materials currently studied could be divided into three types: carbon-based materials, pseudocapacitive oxides, and conducting polymers. The main purpose of this review is to give a clear comparison among the intrinsic electrochemical properties of the state-of-art electrode materials for m-SC,

expecting to open up helpful strategies on developing advanced m-SCs in future.

. And the power *P* is calculated by *P* = *E*/*t*, which is further written as *P* = 1/2 IU, if the charge current is I. Therefore, the energy depends on the capacitance and the voltage. Meanwhile, the power depends on the voltage and the working current that is chosen for the operation of the device. However, the current is not chosen arbitrarily, as the capacitance usually decreases with the increase of current. In other words, the energy shows a declining relation with the power, which is usually described as a Ragone plot. In a word, the performance of an m-SC device could be represented by a Ragone plot, or alternatively, it could be described more intrinsically by the voltage and the capacitance versus the working current. Since the footprint area and the occupied volume are limited for m-SCs when integrated into the system, normalized parameters, i.e., the areal or volumetric energy and power densities and capacitance, are the most important in evaluating the m-SCs. The device performance is determined by the intrinsic properties of the electrode materials, the electrolyte, and the device architecture, wherein the electrode materials play the most critical role and attract abundant studies. In order to compare the properties of different electrode materials effectively, the volumetric energy and power densities as well as the volumetric capacitance should be assessed. This is different from the performance metrics for conventional macro-supercapacitors, where gravimetric capacitance, energy, and power densities are emphasized. Due to the microscaled size, the weight of the electrode materials becomes almost negligible, while the volumetric parameters are the most important concern in developing practical devices. On the other hand, mass density is not a limiting factor, which remarkably expands the choices of the

cycle life should persist for the m-SCs.

22 Supercapacitor Design and Applications

electrode materials for novel m-SCs.

**3. Electrode materials for micro-supercapacitors**

CU2

Carbon is a typical electrode material for EDLC, which possesses good reversibility and cyclability but limited specific capacitance. For conventional supercapacitors, carbon materials are usually fabricated as powders with a loose structure to facilitate the ions diffusion, so as to acquire a high gravimetric capacitance and fast response rate, but the volumetric capacitance is sacrificed. In addition, it is difficult to prepare uniform and qualified thin films from powders, which hinders the direct application of carbonaceous powders in the microelectrodes. Aiming at carbon-based m-SCs, Chimiola et al. tackle the problem by embracing the technological hurdles [12]. They found the carbide-derived carbon (CDC) films produced by selective etching from metal carbides exhibiting an unprecedentedly high volumetric capacity, holding the promise of developing advanced m-SCs. As a first step, they demonstrated that a thin CDC film on a bulk TiC ceramic plate with strong interface adherence was fabricated by direct chlorination at elevated temperatures. Ti is extracted from TiC as TiClO4, forming a porous carbon film, while the conductive TiC plays as both the substrate and the current collector. Both the porosity of the CDC film and the etching speed are closely related to the chlorination temperature [20]. After 15 s etching at 500°C, a CDC film of 2 μm thick, with a pore size of about 0.7 nm, was synthesized. Its volumetric capacitance reaches 180 and 160 F/cm3 in an acetonitrile solution dissolving 1 M tetraethylammonium tetrafluoroborate (TEABF4) and an aqueous solution of 1 M H2SO4, respectively, which are contributed by almost pure EDLC. They proposed the fabrication processes for m-SCs based on CDC films, which is schematically shown in **Figure 2**. The deposition of the precursor carbides and gold collectors could be conducted by well-known chemical and physical vapor depositions (CDC and PVD). The chlorination and the plasma etching of the photolithography are wellestablished techniques. Thus, the good compatibility with the semiconductor industry highlighted the promise of CDC-based m-SC devices.

Heon et al. [21] continued the exploration of the electrochemical property of the CDC films, aiming at the CDC-based m-SCs. The synthesis of uniform and adherent porous CDC films on various substrates by reactive DC magnetron sputtering and chlorination was realized, and the high volumetric capacitance of ∼180 F/cm3 in 1.5 M TEABF4/acetronitrile electrolyte was achieved. Later, the on-chip m-SC from CDC films was fabricated and tested by Huang et al. [22]. The preparation process is generally similar to that proposed in **Figure 2a–d** except that Photolithography step is applied on the TiC film to produce interdigitated electrodes before the chlorination and the current collectors deposited on the CDC electrodes are Ti/Au layers instead of the single Au layer. The active material, i.e., TiC-CDC film, was 1.6 μm thick. The device was dipped into 1 M NEt4BF4 in propylene carbonate (PC) electrolyte in glove box under Ar atmosphere for electrochemical tests, exhibiting the EDLC behavior over a potential window of 2 V with an areal capacitance of 1.5 mF/cm2 , a maximum energy density of 3.0 mJ/ cm2 , and a maximum power density of 84 mW/cm2 . At this point, the EDLC property of TiC-CDC film and the feasibility of manufacturing on-chip m-SCs that are to be integrated into MEMS and electronics have been fully demonstrated, which should be recognized as a representative research on the development of carbon-based m-SC device. Nevertheless, the energy and power performance of the TiC-CDC m-SC are within the range of values reported for other carbon-based micro-supercapacitors, in spite of the high volumetric capacitance of the TiC-CDC active material [10, 22–24]. This is fundamentally attributed to the intrinsic limitation of the capacitance inhibited from the EDLC behavior, which only involves the variation of the ion concentrations in the electric double layer, and the theoretical capacitance is 16–50 μF/cm2 .

**Figure 2.** (A-D) Schematic of the fabrication of an on-chip m-SC based on the CDC film process, using standard photolithography. (E) Schematic of CDC synthesis and a sandwiched device for electrochemical test. The SEM micrograph shows the interface of CDC/TiC. Reproduced with permission from ref. 12. Copyright (2010) American Association for the Advancement of Science.

To acquire high surface area within the limited electrode volume is still the most effective way to obtain high capacitance for the carbonaceous m-SCs. In this case, activated carbon is still a good choice. However, many well-developed fabrication routes of activated carbon powders are not easily applicable for synthesis of thin activated carbon films (ACFs) which is necessary for m-SCs. The challenges include the formation of cracks in the carbon film due to the shrink of polymers during heating and carbonization, the weak interface between the polymer and the substrate resulted from the large stress produced under the harsh synthesis and cooling conditions, and the possible damage to the brittle film in the photolithographical process. Wei et al reported an effective way to minimize some of the interface stresses in order to fabricate ACFs, namely catalyst-assisted low-temperature carbonization of an organic compound solution [25]. The sucrose and H2SO4 (as a catalyst) aqueous solution was spin-coated onto a silicon wafer, dried at room temperature, carbonized, annealed at 700°C in vacuum to remove decomposition products of the carbohydrate and catalyst residues, and activated by annealing at 900°C in CO2 to induce open porosity within carbon. ACFs of 1–2 μm thick were thus produced free of microcracks, while the interface adhesion to the SiO2/Si wafer could be reinforced by further annealing at 1100°C in Ar, which could survive lithographical patterning as evidenced experimentally. The electrochemical property of such ACF film electrodes was tested in a symmetric sandwich-type configuration with 1 M H2SO4 as the electrolyte solution. Exhibiting typical EDLC property by the rectangular cyclic voltammetry (CV) curves, an extremely high volumetric capacitance of 390 F/cm3 was obtained under the slow scan rate of 1 mV/s, which is the highest value reported for carbon film electrode materials at present (see **Figure 3**). Moreover, **Figure 3** shows clearly that the performance was strongly affected by the activation time. As the film becomes more porous after longer activation time, the capacitance increases and the rate capability improves as well, which should be due to easier accessibility of ions into the films. It means that the volumetric metrics of the carbonaceous electrode materials could be greatly enhanced when the materials structure is carefully optimized with proper fabrication techniques. But it is also worth noting that the fabrication process is still quite harsh as several times of high temperature annealing are required, which probably causes difficulties for the practical manufacturing of the devices integrated on the chips, although the realization of uniform and robust ACF films represents a great progress itself.

for other carbon-based micro-supercapacitors, in spite of the high volumetric capacitance of the TiC-CDC active material [10, 22–24]. This is fundamentally attributed to the intrinsic limitation of the capacitance inhibited from the EDLC behavior, which only involves the variation of the ion concentrations in the electric double layer, and the theoretical capacitance

**Figure 2.** (A-D) Schematic of the fabrication of an on-chip m-SC based on the CDC film process, using standard photolithography. (E) Schematic of CDC synthesis and a sandwiched device for electrochemical test. The SEM micrograph shows the interface of CDC/TiC. Reproduced with permission from ref. 12. Copyright (2010) American Association for

To acquire high surface area within the limited electrode volume is still the most effective way to obtain high capacitance for the carbonaceous m-SCs. In this case, activated carbon is still a good choice. However, many well-developed fabrication routes of activated carbon powders are not easily applicable for synthesis of thin activated carbon films (ACFs) which is necessary for m-SCs. The challenges include the formation of cracks in the carbon film due to the shrink of polymers during heating and carbonization, the weak interface between the polymer and the substrate resulted from the large stress produced under the harsh synthesis and cooling conditions, and the possible damage to the brittle film in the photolithographical process. Wei et al reported an effective way to minimize some of the interface stresses in order to fabricate ACFs, namely catalyst-assisted low-temperature carbonization of an organic compound solution [25]. The sucrose and H2SO4 (as a catalyst) aqueous solution was spin-coated onto a silicon wafer, dried at room temperature, carbonized, annealed at 700°C in vacuum to remove

is 16–50 μF/cm2

the Advancement of Science.

.

24 Supercapacitor Design and Applications

**Figure 3.** Rate performance of symmetric ACF electrode cells in 1 MH2SO4 electrolyte: volumetric capacitance of ACFs as a function of CV scan rate. Reproduced with permission from ref. 25. Copyright (2013) American Chemical Society.

As a matter of fact, there have also been great advances in the research of carbon-based electrode materials for macro-scaled supercapacitors that pursue higher volumetric property in recent years. For example, liquid-mediated dense graphenes [26] and nitrogen-doped mesoporous carbon [27] were reported to have unprecedentedly high volumetric capacitance that is comparable with pseudocapacitive materials. Pseudocapacitance has been induced on these modified carbon materials actually [28]. Yang et al. tackle the problem by considering the paradox between gravimetric capacitance *C*wt and packing density of the carbon *ρ* [26]. The specific volumetric capacitance *C*vol is simply calculated by *C*vol = *C*wt× *ρ*, while *C*wt is always compromised with the increase of *ρ*. They addressed this challenge with liquid electrolytemediated chemically converted graphene (EM-CCG) films. They started with the chemically reduced graphene oxide, namely chemically converted graphene (CCG) sheets, which are well dispersed in water and could self-assemble to form an oriented hydrogel film through a directional flow-induced bottom-up assembly process. With the CCG sheets remaining largely separated, a high *C*wt of over 200 F/g was obtained, while the packing density was only ∼0.069 g/cm3 , resulting in a mediocre *C*vol of ∼18 F/cm3 . In order to compress the CCG hydrogel films, they were exchanged with a miscible mixture of volatile and nonvolatile liquids and then subjected to removal of the volatile liquid by vacuum evaporation. As a consequence, the film thickness was reduced, while the sheets remained solvated by the nonvolatile liquid (e.g., sulfuric acid). Through electrochemical tests in 1 M H2SO4 electrolyte, they found that the *C*vol of the EM-CCG films was nearly proportional to *ρ*, and the highly compact one with *ρ* of 1.33 g/cm3 yielded a *C*vol of 255.5 F/cm3 at 0.1 A/g, which is much higher than previous porous carbon materials for conventional SCs. Although the adaptable intersheet spacing among the graphene sheets is particularly emphasized in this work to optimize the *C*vol, other factors such as surface wettability and the pore interconnectivity are also important to realize superior capacitive property for the carbon materials [28]. In other words, both surface chemical property and structural configuration play significant roles in determining the volumetric capacitance of the carbonaceous electrode materials. Lin et al. made a breakthrough in the chemical way, finding that a nitrogen-doped ordered mesoporous few-layer carbon has an extremely high specific capacitance of 855 F/g when tested in 0.5 M H2SO4 electrolyte at 1 A/g (comparing with the 200 F/g for the CCG sheets) [27]. The extra capacitance comes from the pseudocapacitance contributed by the doped N in the pyrrolic and pyridine sites incorporating protons. To be assembled into supercapacitors, they studied variation in the volumetric capacitance with the mass loading of this N-doped mesoporous carbon materials, which showed that the highest value of 560 F/cm3 was reached at the loading of 6 mg/cm2 . Although great advances have been accomplished in the carbonaceous materials for the macro-scaled SCs, they are in the form of either self-standing membranes or powders, difficult to be integrated into the on-chip m-SCs.

Besides the attempts to enhance the volumetric capacitance for energy storage through a proper film fabrication process, there is another important direction for carbon-based m-SC, i.e., to utilize its fast response rate to replace electrolytic capacitors by making open-structured devices [29, 30]. Pech et al. produced m-SCs through electrophoretic deposition of a severalmicrometre-thick layer of nanostructured carbon onions (OLC) with diameters of 6–7 nm onto the interdigital Au current collectors patterned on silicon wafers. The OLC particles were prepared by annealing nanodiamond powder at 1800°C. A stable capacitive behavior was obtained for the microdevice over a 3 V potential window in a 1 M solution of tetraethylammonium tetrafluoroborate in PC, with a linear dependence of the discharge current on the scan rate and low resistive contributions up to 100 V/s, which is about three orders of magnitude higher than conventional SCs. Such a microdevice preserves an areal capacitance of 0.9 mF/cm2 at 100 V/s, which is comparable to values usually reported at much lower scan rates (1–100 mV/s) for microscaled EDLC devices (0.4–2 mF/cm2 ) [9, 11, 31]. The most appealing feature of such OLC m-SC is the extremely small characteristic relaxation time constant *τ*<sup>0</sup> (the minimum time needed to discharge all the energy from the device with an efficiency of greater than 50%), which is only 26 ms, much lower than that of the AC-based microdevice (*τ*0 = 700 ms) or OLC-based macroscopic devices (*τ*0 > 1 s) [32]. **Figure 4** shows the Ragone plot of several typical energy storage devices designed for power microelectronics applications, including a 500-μAh thin-film lithium battery, a 25-mF supercapacitor, and an electrolytic capacitor of the same absolute capacitance, as well as the m-SCs composed of AC, OCL and graphene-based materials. It could be seen that the power density of the OLC-based m-SC has reached that of the electrolytic capacitors, but the energy density is more than one order of magnitude higher than that of latter.

specific volumetric capacitance *C*vol is simply calculated by *C*vol = *C*wt× *ρ*, while *C*wt is always compromised with the increase of *ρ*. They addressed this challenge with liquid electrolytemediated chemically converted graphene (EM-CCG) films. They started with the chemically reduced graphene oxide, namely chemically converted graphene (CCG) sheets, which are well dispersed in water and could self-assemble to form an oriented hydrogel film through a directional flow-induced bottom-up assembly process. With the CCG sheets remaining largely separated, a high *C*wt of over 200 F/g was obtained, while the packing density was only ∼0.069

they were exchanged with a miscible mixture of volatile and nonvolatile liquids and then subjected to removal of the volatile liquid by vacuum evaporation. As a consequence, the film thickness was reduced, while the sheets remained solvated by the nonvolatile liquid (e.g., sulfuric acid). Through electrochemical tests in 1 M H2SO4 electrolyte, they found that the *C*vol of the EM-CCG films was nearly proportional to *ρ*, and the highly compact one with *ρ* of 1.33

materials for conventional SCs. Although the adaptable intersheet spacing among the graphene sheets is particularly emphasized in this work to optimize the *C*vol, other factors such as surface wettability and the pore interconnectivity are also important to realize superior capacitive property for the carbon materials [28]. In other words, both surface chemical property and structural configuration play significant roles in determining the volumetric capacitance of the carbonaceous electrode materials. Lin et al. made a breakthrough in the chemical way, finding that a nitrogen-doped ordered mesoporous few-layer carbon has an extremely high specific capacitance of 855 F/g when tested in 0.5 M H2SO4 electrolyte at 1 A/g (comparing with the 200 F/g for the CCG sheets) [27]. The extra capacitance comes from the pseudocapacitance contributed by the doped N in the pyrrolic and pyridine sites incorporating protons. To be assembled into supercapacitors, they studied variation in the volumetric capacitance with the mass loading of this N-doped mesoporous carbon materials, which

showed that the highest value of 560 F/cm3 was reached at the loading of 6 mg/cm2

great advances have been accomplished in the carbonaceous materials for the macro-scaled SCs, they are in the form of either self-standing membranes or powders, difficult to be

Besides the attempts to enhance the volumetric capacitance for energy storage through a proper film fabrication process, there is another important direction for carbon-based m-SC, i.e., to utilize its fast response rate to replace electrolytic capacitors by making open-structured devices [29, 30]. Pech et al. produced m-SCs through electrophoretic deposition of a severalmicrometre-thick layer of nanostructured carbon onions (OLC) with diameters of 6–7 nm onto the interdigital Au current collectors patterned on silicon wafers. The OLC particles were prepared by annealing nanodiamond powder at 1800°C. A stable capacitive behavior was obtained for the microdevice over a 3 V potential window in a 1 M solution of tetraethylammonium tetrafluoroborate in PC, with a linear dependence of the discharge current on the scan rate and low resistive contributions up to 100 V/s, which is about three orders of magnitude higher than conventional SCs. Such a microdevice preserves an areal capacitance of 0.9 mF/cm2 at 100 V/s, which is comparable to values usually reported at much lower scan rates

. In order to compress the CCG hydrogel films,

. Although

) [9, 11, 31]. The most appealing

at 0.1 A/g, which is much higher than previous porous carbon

g/cm3

g/cm3

, resulting in a mediocre *C*vol of ∼18 F/cm3

yielded a *C*vol of 255.5 F/cm3

26 Supercapacitor Design and Applications

integrated into the on-chip m-SCs.

(1–100 mV/s) for microscaled EDLC devices (0.4–2 mF/cm2

**Figure 4.** A Ragone plot showing the relationship between the volumetric energy density and power density of typical electrolytic capacitors, supercapacitors, batteries, and the m-SCs with AC and OLC electrode [29], as well as the m-SCs with various graphene films [17, 33] and graphene-CNT (rGO-CNT) composite electrode materials [34]. The dashed ellipsoid generally describes the best high power performance currently achieved by these state-of-art m-SCs assembled with conductive carbon electrode materials.

Thereafter, several research papers reported high-power m-SCs fabricated from graphene, whose performance reaches a similar level with that of the OLC-based device (see **Figure 4**). For example, interdigitated graphene m-SCs were produced through laser burning along designed patterns on a graphene oxide (GO) film supported on a PET sheet which was inserted into a LightScribe DVD drive [17]. Due to the photo-thermal effect under laser radiation, the exposed GO was converted into graphene, constructing the positive and negative electrodes, while the unexposed GO remained insulating and served as a separator. A hydrogel-polymer electrolyte, poly(vinyl alcohol) (PVA)-H2SO4, was then drop-cast on the patterned area to create a planar m-SC. Such a device using reduced GO (rGO) as electrode materials exhibits an areal capacitance of 2.32 mF/cm2 and a volumetric capacitance of 3.05 F/cm3 , with a characteristic relaxation time of only 19 ms and a high power density of nearly 200 W/cm3 . There is another planar device using rGO and carbon nanotube (CNT) composites as the electrode material, which is prepared by combining electrostatic spray deposition (ESD) and photolithography lift-off methods [33]. The m-SC delivers an areal capacitance of 6.1 mF/cm2 at 10 mV/s, and a value of 2.8 mF/cm2 is still preserved at 50 V/s, corresponding to 3.1 F/cm3 . Its characteristic time constant is only 4.8 ms. An even faster m-SC device is made from graphene films of only 6–100 nm thick, whose maximum capacitance is 0.807 mF/cm2 and 17.9 F/cm3 (specific values of 0.323 mF/cm2 and 71.6 F/cm3 for the electrode material), with the maximum power density reaching 495 W/cm3 , the maximum energy density 2.5 mWh/cm3 , and the characteristic time constant as short as 0.28 ms [34]. It is concluded from these researches that the electronic conductivity of the electrodes has to be enhanced in order to acquire fast response performance. The most straightforward route is to reduce micropores within the electrodes and enlarge the open area in direct contact with the electrolyte, as the outmost surface is the most easily accessible with the ions for charge/discharge processes. However, such a design is usually at the cost of volumetric capacitance, and the film thickness should be thin as well, which thus limits the areal capacitance of the device.

#### **3.2. Pseudocapacitive oxides**

For the pseudocapacitive materials, Faradaic charge transfer occurs on the electrode/electrolyte interface during charge/discharge processes, giving much higher areal capacitance than EDLC does. Many transition metal oxides exhibit pseudocapacitive behavior in certain aqueous electrolyte. RuO2 is the first discovered pseudocapacitive oxide and still the most ideal candidate [35, 36]. Within the potential range of 0–1.4 V vs. SHE, RuO2 continuously changes its valence from Ru2+ to Ru4+, following the reaction of RuO2 <sup>+</sup> H+ <sup>+</sup> e− RuO2 − (OH), where 0 ≤ *x* ≤ 2 [37]. The process undergoes through both electron transfer and proton incorporation in RuO2 particles. Because of the good electronic conductivity and proton conductivity for hydrated RuO2, fast and reversible charge/discharge pseudocapacitance according to adsorption isotherm model [38] is observed, with a specific capacitance value over 600 F/g. The practical capacitance is closely related to the crystallinity of the material. For crystalline RuO2, protons only adsorb on the surface instead of entering the grains, which provides a capacitance per real surface area of 339–490 μF/cm2 and an overall specific capacitance of about 380 F/g [37]. For amorphous RuO2·*x*H2O, protons could easily transport inside the domains, thus presenting a much higher specific capacitance.

There have been researches on RuO2 thin-film electrodes ever since 1990s. Jow and Zheng coated an amorphous RuO2 film onto Ti substrate through sol-gel method. In spite of inferior uniformity and many cracks, the film still shows a capacitance of 40 mF/cm2 as tested at 50 mV/s in 0.5 M H2SO4, which decreases by only 10% at 500 mV/s [37]. Zheng et al. further employed pulsed laser deposition (PLD) to prepare the RuO2 films [37]. The amorphous film deposited at room temperature possessed the highest capacitance (6.3 mF/cm2 ). As the deposition temperature increased to above 200°C, the capacitance reduced to only 0.3 mF/ cm2 . The diffusion length of protons in amorphous RuO2 film was estimated to be about 5.8 nm, while it reached >11.1 nm in RuO2·*x*H2O. Assuming the diffusion coefficient to be >10−8 cm2 /s, the proton in and out diffusion from RuO2 film is in the order of 10 μs for a diffusion length less than 11.1 nm. This explains the good rate capability of the RuO2 electrode (i.e., it could be charged/discharged at a rate of over 500 mV/s without loss of the capacitance). Thus, for the RuO2 electrode, the charge/discharge rate is mainly limited by the electric resistance and the proton transport in the electrolyte, rather than the proton diffusion inside the RuO2 material.

which is prepared by combining electrostatic spray deposition (ESD) and photolithography lift-off methods [33]. The m-SC delivers an areal capacitance of 6.1 mF/cm2 at 10 mV/s, and a

is still preserved at 50 V/s, corresponding to 3.1 F/cm3

and 71.6 F/cm3 for the electrode material), with the maximum power density

time constant is only 4.8 ms. An even faster m-SC device is made from graphene films of only

constant as short as 0.28 ms [34]. It is concluded from these researches that the electronic conductivity of the electrodes has to be enhanced in order to acquire fast response performance. The most straightforward route is to reduce micropores within the electrodes and enlarge the open area in direct contact with the electrolyte, as the outmost surface is the most easily accessible with the ions for charge/discharge processes. However, such a design is usually at the cost of volumetric capacitance, and the film thickness should be thin as well, which thus

For the pseudocapacitive materials, Faradaic charge transfer occurs on the electrode/electrolyte interface during charge/discharge processes, giving much higher areal capacitance than EDLC does. Many transition metal oxides exhibit pseudocapacitive behavior in certain aqueous electrolyte. RuO2 is the first discovered pseudocapacitive oxide and still the most ideal candidate [35, 36]. Within the potential range of 0–1.4 V vs. SHE, RuO2 continuously changes its valence from Ru2+ to Ru4+, following the reaction of RuO2 <sup>+</sup> H+ <sup>+</sup> e− RuO2 − (OH), where 0 ≤ *x* ≤ 2 [37]. The process undergoes through both electron transfer and proton incorporation in RuO2 particles. Because of the good electronic conductivity and proton conductivity for hydrated RuO2, fast and reversible charge/discharge pseudocapacitance according to adsorption isotherm model [38] is observed, with a specific capacitance value over 600 F/g. The practical capacitance is closely related to the crystallinity of the material. For crystalline RuO2, protons only adsorb on the surface instead of entering the grains, which

tance of about 380 F/g [37]. For amorphous RuO2·*x*H2O, protons could easily transport inside

There have been researches on RuO2 thin-film electrodes ever since 1990s. Jow and Zheng coated an amorphous RuO2 film onto Ti substrate through sol-gel method. In spite of inferior uniformity and many cracks, the film still shows a capacitance of 40 mF/cm2 as tested at 50 mV/s in 0.5 M H2SO4, which decreases by only 10% at 500 mV/s [37]. Zheng et al. further employed pulsed laser deposition (PLD) to prepare the RuO2 films [37]. The amorphous film

deposition temperature increased to above 200°C, the capacitance reduced to only 0.3 mF/

. The diffusion length of protons in amorphous RuO2 film was estimated to be about 5.8 nm, while it reached >11.1 nm in RuO2·*x*H2O. Assuming the diffusion coefficient to be >10−8

/s, the proton in and out diffusion from RuO2 film is in the order of 10 μs for a diffusion length less than 11.1 nm. This explains the good rate capability of the RuO2 electrode (i.e., it

deposited at room temperature possessed the highest capacitance (6.3 mF/cm2

, the maximum energy density 2.5 mWh/cm3

6–100 nm thick, whose maximum capacitance is 0.807 mF/cm2

provides a capacitance per real surface area of 339–490 μF/cm2

the domains, thus presenting a much higher specific capacitance.

. Its characteristic

and 17.9 F/cm3 (specific values

, and the characteristic time

and an overall specific capaci-

). As the

value of 2.8 mF/cm2

28 Supercapacitor Design and Applications

reaching 495 W/cm3

limits the areal capacitance of the device.

**3.2. Pseudocapacitive oxides**

of 0.323 mF/cm2

cm2

cm2

**Figure 5.** (a) Cross-section SEM image of a CNW film (bottom) electrodeposited by hRuO2 on its top. (b) Schematic of the vertically aligned CNW decorated with hRuO2 particles. (c) Schematic diagram of on-chip m-SC with 2D architecture. (d) A Ragone plot showing the energy and power density of the CNW/hRuO2-based m-SC, compared with other advanced m-SCs and microbatteries. Reproduced with permission from ref. 39. Copyright (2014) Elsevier Ltd.

The development of RuO2-based m-SCs emerged in recent years. Liu et al. [36] fabricated the planar device through depositing RuO2 nanorods onto the patterned stack layer of Ru/Au/Ti/ SiO2 on silicon wafer, which was subjected to electrodeposition for another layer of hydrous RuO2 on top. It worked well in 0.5 M H2SO4 electrolyte, providing a capacitance of 21.4 mF/cm2 at 50 mV/s and 14.9 mF/cm2 at 500 mV/s. Makino et al. [18] reported the fabrication of an m-SC with ordered mesoporous RuO*x* as the electrode material, which was produced by controlled electro-deposition using a lyotropic liquid crystal template method and subsequent electro-oxidation on an interdigital electrode array. The device exhibited good capacitive property with maximum capacitance of 12.6 mF/cm2 and maximum energy of 1.49 μWh/cm2 at the slowest discharge rate of 0.38 mA/cm2 and maximum power delivery of 750 μW/cm2 at 2.88 mA/cm2 . More recently, a new m-SC device based on hydrous RuO2/carbon nanowalls hierarchical structured composite electrode was proposed by Dinescu et al., which showed an exceptionally high capacitance [39]. Carbon nanowalls (CNW), or vertically oriented graphene sheets, is a good EDLC material with a large surface area, good electronic conductivity, and excellent chemical stability, while RuO2 is an ideal pseudocapacitive materials with a high specific capacitance. A silicon wafer coated with an insulating Si3N4 layer was first deposited with a 40 nm Cr/200 nm Pt layer by evaporation as the current collector and subjected to the CNW layers growth by PECVD at 700°C, and then, electrodeposition of hydrous RuO2 (hRuO2) onto the CNW was carried out afterwards, after which the samples were annealed in air at 150 °C. The pristine CNW layer is 12 μm thick, with a capacitance of 5.7 mF/cm2 in 0.5 M H2SO4 electrolyte, close to the values of other carbonaceous electrodes. When about half of the CNW layer was decorated with hRuO2 (see **Figure 5a, b**), the hybrid electrode exhibited an extremely high capacitance of 1094 mF/cm2 at 2 mV/s, which is three orders of magnitude higher than that of the state-of-the-art graphene-based m-SCs [34], and also far larger than most other advanced m-SC electrodes [22, 25, 33, 40]. An all-solid-state m-SC in a stack configuration was realized with a solid-polymer electrolyte sandwiched between two CNW/ hRuO2 electrodes (see **Figure 5c**), delivering an energy density of 49 μWh/cm2 , i.e., 20 mWh/ cm3 . Such a value is even comparable to the state-of-the-art lithium ion microbatteries [41–43], but its power density and cycle life (more than 90% is retained after 2000 cycles) are much higher than that of the latter, which is shown in **Figure 5d**.

In spite of the ideal pseudocapacitve property, RuO2 is too expensive for large-scale application. Cheaper oxides have been widely researched, such as MnO2 and NiO, wherein MnO2 attracts the most attention [44, 45]. MnO2 works in neutral aqueous solutions, with the potential window within 0.8 V and above 0 V vs. Ag/AgCl, and thus is suitable to serve as a positive electrode in asymmetric devices. Its working mechanism is the surface adsorption/desorption of the electrolytic cations and protons in the solution, which is described as the reaction MnO2 <sup>+</sup> C+ <sup>+</sup> H+ <sup>+</sup> <sup>+</sup> e− MnOOCH . [46, 47]. It is obvious from the CV curve on **Figure 6** that the charge/discharge behavior of MnO2 is similar to that of EDLC. Besides, the abundant resource and the safe working condition of neutral solutions also boost the wide research on MnO2, although the specific capacitance of MnO2 powders or micrometer-thick films is only 150 F/g.

**Figure 6.** The CV curve of MnO2 electrode tested in 2 M Li2SO4 aqueous electrolyte solution..

The intrinsic electrochemical properties of MnO2 films produced by different preparation conditions and of different morphologies and structures have been explored [48–51]. For example, the MnO2 film deposited by cathodic electrodeposition is poorly crystalline and porous, whose specific capacitance strongly decreases with the scan rate and the film thickness [51]. The specific capacitance of a film bearing a deposited mass of 45 μg/cm2 is tested to be 353 F/g (15.9 mF/cm2 ) at 2 mV/s and 135 F/g (6.1 mF/cm2 ) at 100 mV/s. For films annealed under 200 °C, whose porosity is reduced and crystallinity increased, the maximum specific capacitance is decreased, while rate capability is improved more or less. PLD has also been utilized to prepare manganese oxide films, with amorphous MnO*x*, crystalline Mn2O3, and Mn3O4 to be produced under different deposition temperatures and the partial pressure of oxygen [50]. The crystalline Mn2O3 film possesses the highest specific capacitance, 210 F/g at 1 mV/s for a film of 120 nm, while the Mn3O4 film has the lowest value.

(hRuO2) onto the CNW was carried out afterwards, after which the samples were annealed in air at 150 °C. The pristine CNW layer is 12 μm thick, with a capacitance of 5.7 mF/cm2

M H2SO4 electrolyte, close to the values of other carbonaceous electrodes. When about half of the CNW layer was decorated with hRuO2 (see **Figure 5a, b**), the hybrid electrode exhibited an extremely high capacitance of 1094 mF/cm2 at 2 mV/s, which is three orders of magnitude higher than that of the state-of-the-art graphene-based m-SCs [34], and also far larger than most other advanced m-SC electrodes [22, 25, 33, 40]. An all-solid-state m-SC in a stack configuration was realized with a solid-polymer electrolyte sandwiched between two CNW/

. Such a value is even comparable to the state-of-the-art lithium ion microbatteries [41–43], but its power density and cycle life (more than 90% is retained after 2000 cycles) are much

In spite of the ideal pseudocapacitve property, RuO2 is too expensive for large-scale application. Cheaper oxides have been widely researched, such as MnO2 and NiO, wherein MnO2 attracts the most attention [44, 45]. MnO2 works in neutral aqueous solutions, with the potential window within 0.8 V and above 0 V vs. Ag/AgCl, and thus is suitable to serve as a positive electrode in asymmetric devices. Its working mechanism is the surface adsorption/desorption of the electrolytic cations and protons in the solution, which is described as the reaction

**Figure 6** that the charge/discharge behavior of MnO2 is similar to that of EDLC. Besides, the abundant resource and the safe working condition of neutral solutions also boost the wide research on MnO2, although the specific capacitance of MnO2 powders or micrometer-thick

**Figure 6.** The CV curve of MnO2 electrode tested in 2 M Li2SO4 aqueous electrolyte solution..

The intrinsic electrochemical properties of MnO2 films produced by different preparation conditions and of different morphologies and structures have been explored [48–51]. For example, the MnO2 film deposited by cathodic electrodeposition is poorly crystalline and porous, whose specific capacitance strongly decreases with the scan rate and the film thick-

hRuO2 electrodes (see **Figure 5c**), delivering an energy density of 49 μWh/cm2

higher than that of the latter, which is shown in **Figure 5d**.

MnO2 <sup>+</sup> C+ <sup>+</sup> H+ <sup>+</sup> <sup>+</sup> e− MnOOCH

films is only 150 F/g.

cm3

30 Supercapacitor Design and Applications

in 0.5

, i.e., 20 mWh/

. [46, 47]. It is obvious from the CV curve on

Because of the insulating property of MnO2 and the increased difficulty for ions to access into a thicker film, the rate capability of MnO2 film is always unsatisfying. Research efforts have been devoted to enhance the electronic conductivity so as to improve the rate capability. Si et al. [52] deposited the MnO*x*/Au multilayer film, which showed a capacitance of 32.8 F/cm3 , higher than that of the pure MnO*<sup>x</sup>* film electrode (19.9 F/cm3 ). Two kinds of interdigital m-SCs were prepared using these two films. The device of MnO*x*/Au multilayer possessed a smaller equivalent series resistance (ESR) and a shorter characteristic relaxation time (5 ms). Doping is another way to adjust the electrochemical property. By doping Mo into the electrodeposited manganese oxide, a film of MnMo6+0.18O1.18(OH)0.59(H2O)0.25 was obtained, whose specific capacitance (190.9 F/g and 18.5 mF/cm2 at 5 mV/s), cycling reversibility, and rate capability were all improved as compared with the undoped film [53]. It was found that the electronic resistivity of a ∼0.5-μm-thick film was reduced from 5.0 × 104 to 3.0 × 101 Ω cm after being doped with Mo, which suggested that the enhancement of the electrochemical property is mainly attributed to the increase in electronic conductivity. Similar effect was also discovered in doping Co into the electrodeposited manganese oxide film and in doping V into the PLDdeposited amorphous manganese oxide film [54, 55].

Other pseudocapacitve oxides include Co3O4, NiO, NiCo2O4, and so on [56–58]. They could reversibly form hydroxides in alkaline solutions, the process of which could be represented as M3O4 + OH<sup>−</sup> + H2O 2MOOH + e−, wherein M refers to elements such as Ni and Co. Their specific capacitances are reported to be very high (382–1400 F/g); however, phase transformations are always involved during the reaction processes, and the charge/discharge behaviors are more like that of batteries, for example, with potential plateaus, short potential windows (0.4–0.5 V), limited rate capabilities [56].

Overall, pseudocapacitive oxides could exhibit high specific capacitance, but poor rate capability and cycling stability hinder their application, which is mainly attributed to their insulating property. For the film electrodes, when thickness increases, the areal capacitance rarely scales up as expected. To disperse the active materials onto a 3D current collector is a common way to acquire certain areal capacitance with acceptable rate capability, which relies on the innovation of nanofabrication techniques. Nevertheless, the intrinsic relationship between the comprehensive electrochemical property and the microstructure of the oxides is still worth studying, which may help to optimize the intrinsic electrochemical performance of the film electrodes, contributing to the development of m-SCs compatible with conventional microelectronics manufacturing techniques.

**Figure 7.** (a) Schematic of the multi-phased microstructure of the MoO*<sup>x</sup>* film deposited at 150°C by magnetron sputtering. (b) CV curves of MoO2+*x* and MnO2 films in 2 M Li2SO4 electrolyte, at 50 mV/s. (c) Illustrative diagram for the working process and (d) a Ragone plot showing energy density vs. power density of MoO2+*x*(−)//2 M Li2SO4//MnO2(+) microdevice. The performances of representative m-SCs of EDLC type (symmetric device based on CDC electrode [22]) and other pseudocapacitive type (asymmetric device based on VN and NiO electrodes [61]) are also plotted for reference [59, 60] Reproduced with permission from ref. 60. Copyright (2014) Elsevier Ltd.

Our group has systematically studied the fabrication and electrochemical property of molybdenum oxide thin film, which has various chemical valences with quite different properties [59]. The composition, microstructure, and morphology were controlled to enhance the electrochemical performance of the molybdenum oxide film, and its potential to be applied as a superior electrode material in m-SC is evaluated. We fabricated electronically conductive MoO2+*<sup>x</sup>* films via RF magnetron sputtering from a MoO3 target. Multi-valence composition and mixed-phased microstructure, i.e., coexistence of MoO2 nanocrystals and amorphous MoO*<sup>x</sup>* (2 < *x* ≤ 3), were acquired in these films (see **Figure 7a**), which exhibit excellent pseudocapacitance in Li2SO4 electrolyte [59]. The MoO*x* (*x* ≈ 2.3) film deposited at 150°C presented an areal capacitance of 31 mF/cm2 at 5 mV/s, corresponding to a volumetric value of 392 F/cm3 , superior to most of the advanced m-SC electrode materials previously discussed. Forty-seven percent of the capacitance was retained when the scan rate increases from 20 to 500 mV/s, meaning good rate capability. The cycling stability is excellent as well, with 100% preserved after 5000 cycles. The multi-phased microstructure of such MoO*<sup>x</sup>* films is quite interesting, as it intrinsically endows the material with superior electrochemical property. The pseudocapacitance originates from the cation (H+ and Li+ ) insertion/extrusion in the amorphous MoO*x*, in which H+ is more active. The MoO2 grains could also catalyze the decomposition of water combined on the surface, producing H atoms that could be reversibly stored in amorphous MoO*<sup>x</sup>* and thus promote the pseudocapacitive process. Furthermore, the crystalline MoO2 also improves the electronic conductivity and maintains a stable structure of the film [60]. Another interesting feature about MoO*<sup>x</sup>* film is its relatively negative potential window, i.e., between −1.1 and 0 V vs. SCE, which makes it a proper anode material relative to other pseudocapacitive oxides. Thus, an asymmetric microdevice of MoO2+*x*(−)//2M Li2SO4//MnO2(+) is successfully fabricated. It could be seen from **Figure 7b** that the working potential windows of the two electrodes well complement to each other. The working mechanism is described in the schematic of **Figure 7c**. A high energy density of 2.8 μWh/cm2 at a real power density of 0.35 mW/cm2 was obtained from this device, combined with good stability (no capacitance loss for 10,000 cycles). The Ragone plot of **Figure 7d** shows that the performance of this asymmetric device is much better than other typical state-of-art m-SCs.


a Carbide-derived carbon (CDC), reduced graphene oxide (rGO), carbon nanotubes (CNT). b Please note that *C*V here corresponds to the electrodes, and some of the values are read from the reported plots, or estimated from the cell data.

**Table 1.** Electrochemical properties of typical m-SC microelectrodes in literature [62].

**Figure 7.** (a) Schematic of the multi-phased microstructure of the MoO*<sup>x</sup>* film deposited at 150°C by magnetron sputtering. (b) CV curves of MoO2+*x* and MnO2 films in 2 M Li2SO4 electrolyte, at 50 mV/s. (c) Illustrative diagram for the working process and (d) a Ragone plot showing energy density vs. power density of MoO2+*x*(−)//2 M Li2SO4//MnO2(+) microdevice. The performances of representative m-SCs of EDLC type (symmetric device based on CDC electrode [22]) and other pseudocapacitive type (asymmetric device based on VN and NiO electrodes [61]) are also plotted for refer-

Our group has systematically studied the fabrication and electrochemical property of molybdenum oxide thin film, which has various chemical valences with quite different properties [59]. The composition, microstructure, and morphology were controlled to enhance the electrochemical performance of the molybdenum oxide film, and its potential to be applied as a superior electrode material in m-SC is evaluated. We fabricated electronically conductive MoO2+*<sup>x</sup>* films via RF magnetron sputtering from a MoO3 target. Multi-valence composition and mixed-phased microstructure, i.e., coexistence of MoO2 nanocrystals and amorphous MoO*<sup>x</sup>* (2 < *x* ≤ 3), were acquired in these films (see **Figure 7a**), which exhibit excellent pseudocapacitance in Li2SO4 electrolyte [59]. The MoO*x* (*x* ≈ 2.3) film deposited at 150°C presented an areal

to most of the advanced m-SC electrode materials previously discussed. Forty-seven percent of the capacitance was retained when the scan rate increases from 20 to 500 mV/s, meaning good rate capability. The cycling stability is excellent as well, with 100% preserved after 5000 cycles. The multi-phased microstructure of such MoO*<sup>x</sup>* films is quite interesting, as it intrinsically endows the material with superior electrochemical property. The pseudocapacitance

H+ is more active. The MoO2 grains could also catalyze the decomposition of water combined on the surface, producing H atoms that could be reversibly stored in amorphous MoO*<sup>x</sup>* and thus promote the pseudocapacitive process. Furthermore, the crystalline MoO2 also improves the electronic conductivity and maintains a stable structure of the film [60]. Another interesting feature about MoO*<sup>x</sup>* film is its relatively negative potential window, i.e., between −1.1 and 0 V vs. SCE, which makes it a proper anode material relative to other pseudocapacitive oxides. Thus, an asymmetric microdevice of MoO2+*x*(−)//2M Li2SO4//MnO2(+) is successfully fabricated.

at 5 mV/s, corresponding to a volumetric value of 392 F/cm3

) insertion/extrusion in the amorphous MoO*x*, in which

, superior

ence [59, 60] Reproduced with permission from ref. 60. Copyright (2014) Elsevier Ltd.

and Li+

capacitance of 31 mF/cm2

32 Supercapacitor Design and Applications

originates from the cation (H+

The pseudocapacitive MoO*<sup>x</sup>* film was further prepared by electrodeposition, with its electrochemical property adjusted by annealing under different conditions [62]. Optimal experimental parameters were determined to fabricate the film containing MoO2 nanocrystallites and amorphous MoO*x*. A film of 63 nm thick exhibited a high volumetric capacitance of 700 F/cm3 , good rate capability with a relaxation time constant of 11 ms, and excellent cycling stability of 99% capacitance retention after 4000 cycles. The performance is superior to other typical microelectrodes for m-SCs, the comparison of which is listed in **Table 1**. Furthermore, a 3D microelectrode was developed by electrodepositing MoO*<sup>x</sup>* on a Ti nanorod array prepared by oblique angle deposition [63]. An areal capacitance of 27 mF/cm2 , corresponding to a high volumetric capacitance of 643 F/cm3 , was obtained as well as satisfying cycling stability, which is rather attractive compared to other 3D microelectrodes. And post-annealing in reductive atmosphere improved its rate capability and response speed. Thus, the further improvement in electrochemical property of MoO*x* electrode by architecture design of employing current collectors with large specific area promotes its practical application in m-SCs.

#### **3.3. Conducting polymers**

Conducting polymers are another group of pseudocapacitive materials, which work through the fast redox reaction of ion doping [68, 69]. Currently studied conducting polymers mainly include P-type doping materials such as polypyrrole (PPy) and polyaniline (PANI), and Ntype doping materials such as polythiophene (PTH). They are electronically conductive, leading to low ESR, and their capacitance is generally 2–3 times as high as that of activated carbon materials. However, the cycling stability is always unsatisfying due to the large volume expansion and shrink during the charge/discharge processes. The application of conducting polymers in m-SCs is relatively pioneering, i.e., developed in 3D structures [16, 70, 71]. Beidaghi and Wang [71] fabricated such a 3D-structured interdigital m-SC through carbonmicroelectrochemical system (C-MEMS) technology. Briefly, the carbonization of patterned photoresist pillars produced microarrays of carbon pillars on the interdigital carbon layer supported on silicon wafer. The carbon arrays served as the C-MEMS current collectors, on which PPy was coated by electrodeposition. The resulted PPy/C-MEMS electrode presented a high areal capacitance of 162 mF/cm2 (volumetric capacitance estimated to be 11.6 F/cm3 ) and a power density of 1.62 mW/cm2 at 20 mV/s scan rate. Correspondingly, the entire symmetric m-SC device exhibited an average capacitance of 78 mF/cm2 and a power density of 0.63 mW/ cm2 . In spite of the high areal capacitance, the electrodes were not robust for consecutive cycling, as only 56% of the capacitance is retained after 1000 cycles.

#### **4. Summary**

Architecture design is especially important for m-SC, which significantly affects the comprehensive device performance. Structure of a separator layer sandwiched by thin-film electrodes is the traditional architecture, but the performance is closely related to the thickness of the electrode layer, which hinders it from scaling up to acquire higher areal energy density [12]. In-plane interdigital architecture is often employed in recent years, especially for fast response devices [13, 17]. Besides, integrated 3D architectures have also been proposed for m-SCs [19]. The microelectrodes with the structure of micro- or nanoarrays as cited above [16, 63, 70, 71] could be considered as ordered 3D structure, which is already realized by microfabrication techniques. However, a "true" 3D device is composed to consist of interpenetrating electrodes that are separated by a very thin layer of electrolyte, which is proposed by Long and coworkers [5]. Although several possible strategies for fabrication of electrochemical energy storage devices with 3D architectures have been proposed [72], there has been no work fully realized the concept of 3D m-SC [6].

Above all, structure control is still the most important factor in exploring the property limits of different kinds of electrode materials. Although excellent efforts have been made on developing carbon film electrode with better volumetric and areal capacitances, the value is still lower than that of pseudocapacitive materials, which is limited by the theoretical limits of EDLC. For the pseudocapacitive materials such as molybdenum oxide, it is possible to optimize the intrinsic electrochemical property by structure design, i.e., to obtain both good rate capability and cyclability as well as keep relatively high capacitance. However, the application of the oxides is seriously limited by the electrolyte condition, which is always a disadvantage compared with carbon materials. Pseudocapacitive carbon materials represent another promising trend to achieve balanced performance. In-plane m-SCs with interdigital architectures have proven its success to fabricate fast response devices with carbon electrodes. Thus, further study on the interplay effect between such architecture and pseudocapacitive materials is in need. Exploring 3D architecture for m-SC is a difficult but attracting challenge.

## **Acknowledgements**

The pseudocapacitive MoO*<sup>x</sup>* film was further prepared by electrodeposition, with its electrochemical property adjusted by annealing under different conditions [62]. Optimal experimental parameters were determined to fabricate the film containing MoO2 nanocrystallites and amorphous MoO*x*. A film of 63 nm thick exhibited a high volumetric capacitance of 700 F/cm3

good rate capability with a relaxation time constant of 11 ms, and excellent cycling stability of 99% capacitance retention after 4000 cycles. The performance is superior to other typical microelectrodes for m-SCs, the comparison of which is listed in **Table 1**. Furthermore, a 3D microelectrode was developed by electrodepositing MoO*<sup>x</sup>* on a Ti nanorod array prepared by

is rather attractive compared to other 3D microelectrodes. And post-annealing in reductive atmosphere improved its rate capability and response speed. Thus, the further improvement in electrochemical property of MoO*x* electrode by architecture design of employing current

Conducting polymers are another group of pseudocapacitive materials, which work through the fast redox reaction of ion doping [68, 69]. Currently studied conducting polymers mainly include P-type doping materials such as polypyrrole (PPy) and polyaniline (PANI), and Ntype doping materials such as polythiophene (PTH). They are electronically conductive, leading to low ESR, and their capacitance is generally 2–3 times as high as that of activated carbon materials. However, the cycling stability is always unsatisfying due to the large volume expansion and shrink during the charge/discharge processes. The application of conducting polymers in m-SCs is relatively pioneering, i.e., developed in 3D structures [16, 70, 71]. Beidaghi and Wang [71] fabricated such a 3D-structured interdigital m-SC through carbonmicroelectrochemical system (C-MEMS) technology. Briefly, the carbonization of patterned photoresist pillars produced microarrays of carbon pillars on the interdigital carbon layer supported on silicon wafer. The carbon arrays served as the C-MEMS current collectors, on which PPy was coated by electrodeposition. The resulted PPy/C-MEMS electrode presented a high areal capacitance of 162 mF/cm2 (volumetric capacitance estimated to be 11.6 F/cm3

m-SC device exhibited an average capacitance of 78 mF/cm2 and a power density of 0.63 mW/

Architecture design is especially important for m-SC, which significantly affects the comprehensive device performance. Structure of a separator layer sandwiched by thin-film electrodes is the traditional architecture, but the performance is closely related to the thickness of the electrode layer, which hinders it from scaling up to acquire higher areal energy density [12]. In-plane interdigital architecture is often employed in recent years, especially for fast response

cycling, as only 56% of the capacitance is retained after 1000 cycles.

. In spite of the high areal capacitance, the electrodes were not robust for consecutive

oblique angle deposition [63]. An areal capacitance of 27 mF/cm2

collectors with large specific area promotes its practical application in m-SCs.

volumetric capacitance of 643 F/cm3

34 Supercapacitor Design and Applications

**3.3. Conducting polymers**

a power density of 1.62 mW/cm2

cm2

**4. Summary**

,

) and

, corresponding to a high

, was obtained as well as satisfying cycling stability, which

at 20 mV/s scan rate. Correspondingly, the entire symmetric

The authors are grateful to the financial support by the National Basic Research Program of China (973 program, Grant No. 2013CB934301), the National Natural Science Foundation of China (Grant No. 51531006 and No. 51572148), the Research Project of Chinese Ministry of Education (Grant No. 113007A), and the Tsinghua University Initiative Scientific Research Program.

## **Author details**

Can Liu1 and Zhengjun Zhang2\*

\*Address all correspondence to: zjzhang@tsinghua.edu.cn

1 Institute for Catalysis, Hokkaido University, Sapporo, Japan

2 School of Materials Science and Engineering, Key Laboratory of Advanced Materials, Tsinghua University, Beijing, People's Republic of China

## **References**


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#### **Engineering Nanostructured MnO2 for High Performance Supercapacitors Engineering Nanostructured MnO2 for High Performance Supercapacitors**

#### Jian-Gan Wang Jian-Gan Wang

Additional information is available at the end of the chapter Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/65008

#### **Abstract**

Manganese oxides (MnO2) have particularly received increasing attention owing to their high theoretical specific capacitance of 1370 F/g, low-cost, natural abundance, and environmental benignity. However, MnO2 suffers from low electrical conductivity (10−5 to 10−6 S/cm), low ionic diffusion constant (~10−13 cm2 /V s), and low structural stability, which results in low electrochemical utilization and poor cycling life. It is therefore important to explore new strategies to improve the electrochemical performance of MnO2. The effective methods to maximize the performance involve (i) reducing MnO2 structures to a nanoscale range and (ii) compositing MnO2 with highly conductive materials. In this chapter, we will first introduce the rapid development of MnO2 nanostructures for supercapacitors. Then the fundamental charge storage mechanism of MnO2 will be specifically clarified. The preparation methods of MnO2 nanostructures and their composites will be subsequently summarized. Then, we will pay great attention to the most recent development of MnO2-based nanostructures for supercapacitors, which is the main body of this chapter. The practical application of MnO2 nanostructures for symmetric and asymmetric supercapacitors will be discussed. Finally, we will present a brief perspective regarding the rational design and synthesis of MnO2-based nanostructures.

**Keywords:** MnO2, nanostructure, composite, supercapacitor, high performance

### **1. Introduction**

According to the charge storage mechanism, supercapacitors can be categorized as electrochemical double layer capacitors (EDLCs) and pseudocapacitors [1]. The typical electroactive electrode materials for pseudocapacitors include transition-metal oxides and conducting

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

polymers. To this end, massive efforts have been dedicated to developing low-cost and ecofriendly alternatives, such as MnO2, Co3O4, and NiO. Among the alternative candidates, MnO2 has received intensive attention in the past decades. Lee and Goodenough reported the first use of MnO2 for supercapacitors in 1999 [2]. This finding paved a new avenue for exploring a huge variety of electroactive MnO2-based electrode materials for high performance supercapacitors. This is because MnO2 shows advantages of (i) a high theoretical specific capacitance of 1370 F/g; (ii) a wide operating potential window of about 1.0 V; and (iii) the ability to enable mild aqueous electrolytes with a much less chemical corrosion to current collectors or packages [3, 4]. In addition, MnO2 has additional merits of natural abundance, low-cost, and environmentally benignity [4]. These unique characteristics will enable MnO2-based electrodes to act as a high-performance, safe, and a low-cost replacement.

However, the specific capacitance of MnO2 is almost one magnitude lower than the theoretical value [4–6]. This is because MnO2 shows poor electronic conductivity (10−5 to 10−6 S/cm), low ionic diffusion constant (~10−13 cm2 /V s), and structure flexibility [7, 8]. Moreover, owing to a wide diversity of textural feature, crystal forms (e.g., *α*-, *β*-, *δ*-, *γ*-MnO2), and defect chemistry, MnO2 exhibits a variety of distinct electrochemical performance. The performance of MnO2 will be substantially improved with a better fundamental understanding of the charge storage mechanism at the electrode/electrolyte interfaces and a rational design of MnO2-based electrodes.

In recent years, extensive endeavors have been dedicated to synthesizing MnO2 nanostructures or nanocomposites with different morphologies and crystal forms. This chapter will introduce the state-of-the-art MnO2 electrodes for high performance supercapacitors. The content includes a brief discussion of the charge storage mechanisms of MnO2, a summary of the synthetic methodologies, a main focus on the MnO2-based electrodes and their application in supercapacitors.

## **2. Charge storage mechanism of MnO2**

The charge storage mechanism of MnO2 was first imitated from RuO2 and Zn/MnO2 cells in acidic electrolytes [9], which involve a reversible insertion/extraction of protons into/from MnO2 with a concomitant Mn3+–Mn4+ transition. However, the specific capacitance of MnO2 is found to be related to the species and concentrations of the alkali metal cations, such as Li+ , Na+ , and K+ , regardless of the pH value [10, 11]. Consequently, a more rational mechanism based on the chemisorption of alkali metal cations on the surface of MnO2 was proposed and can be written as [3, 11]:

$$\left(\text{MnO}\_2\right)\_{\text{surface}} + \text{M}^+ + \text{e}^- \rightarrow \left(\text{MnOOM}\right)\_{\text{surface}}\tag{1}$$

where M+ represents alkali metal cations, including Li+ , Na+ , and K+ . Later, Belanger and coworkers developed a new complementary storage mechanism, that is, the alkali metal cations can insert/extract into/from MnO2, as shown in Eq. (2) [11]:

$$\text{MnO}\_2 + \text{M}^+ + \text{e}^- \rightarrow \text{MnOOM} \tag{2}$$

It should be noted that MnO2 can crystallize in various forms. Can the above charge storage mechanisms be suitable for all MnO2 structures? Brousse et al. investigated a systematic comparison on the supercapacitive properties of MnO2 with five types of phase structures [11]. It is revealed that the capacitance is closely associated with the tunnel size of crystalline MnO2, i.e., the larger the tunnel size, the larger the specific capacitance. More specifically, the one-dimensional (1D) *α*-MnO2 (4.6 Å) and two-dimensional (2D) *δ*-MnO2 (7 Å) allow a fast insertion of hydrated K+ cations (3 Å) due to their large tunnel sizes, while *β*-MnO2 and *γ*-MnO2 with one-dimensional tunnel sizes smaller than cations limit the diffusion process, and the spinel *λ*-MnO2 with a more opened three-dimensional (3D) structure can permit a partial cationic diffusion. This suggests that the tunnel space should be large enough to allow a high-rate insertion/extraction of electrolyte ions for charge storage governed by Eq. (2), which is known as "tunnel storage mechanism (TSM)" [12]. Hence, the specific capacitance of MnO2 depends strongly on its crystallographic structure, which follows the decreasing order: *α*(*m*) > *α ≈ δ > γ > λ > β*. Furthermore, owing to the large tunnel, *α*-MnO2 also exhibits ideal pseudocapacitive behavior in the electrolytes containing Ca2+, Mg2+, or Ba2+, which can store more charge through a multivalent cation storage mechanism [5]. In addition to the metal cations, anions can also be the working species that compensate the Mn-valent state variation upon charge/discharge in aprotic ionic liquids (IL) [13]. In summary, the charge/discharge process of MnO2 typically involves (i) electrolyte ion transport; (ii) ion adsorption/desorption at the surface sites of electrodes, which are dependent on the ion size, the mobility of ion, and the dehydration/hydration rate; and (iii) ion insertion/extraction into/from bulk MnO2. Therefore, the charge storage process is not only associated with crystallographic structures, but also relates to other factors, such as specific surface area, electronic conductivity, and ionic conductivity, etc. Therefore, the charge storage mechanism of MnO2 basically involves a capacitive surface chemisorption process (Eq. (1)) and a bulk/subsurface Faradaic reaction (Eq. (2)).

polymers. To this end, massive efforts have been dedicated to developing low-cost and ecofriendly alternatives, such as MnO2, Co3O4, and NiO. Among the alternative candidates, MnO2 has received intensive attention in the past decades. Lee and Goodenough reported the first use of MnO2 for supercapacitors in 1999 [2]. This finding paved a new avenue for exploring a huge variety of electroactive MnO2-based electrode materials for high performance supercapacitors. This is because MnO2 shows advantages of (i) a high theoretical specific capacitance of 1370 F/g; (ii) a wide operating potential window of about 1.0 V; and (iii) the ability to enable mild aqueous electrolytes with a much less chemical corrosion to current collectors or packages [3, 4]. In addition, MnO2 has additional merits of natural abundance, low-cost, and environmentally benignity [4]. These unique characteristics will enable MnO2-based electrodes to act

However, the specific capacitance of MnO2 is almost one magnitude lower than the theoretical value [4–6]. This is because MnO2 shows poor electronic conductivity (10−5 to 10−6 S/cm), low

wide diversity of textural feature, crystal forms (e.g., *α*-, *β*-, *δ*-, *γ*-MnO2), and defect chemistry, MnO2 exhibits a variety of distinct electrochemical performance. The performance of MnO2 will be substantially improved with a better fundamental understanding of the charge storage mechanism at the electrode/electrolyte interfaces and a rational design of MnO2-based

In recent years, extensive endeavors have been dedicated to synthesizing MnO2 nanostructures or nanocomposites with different morphologies and crystal forms. This chapter will introduce the state-of-the-art MnO2 electrodes for high performance supercapacitors. The content includes a brief discussion of the charge storage mechanisms of MnO2, a summary of the synthetic methodologies, a main focus on the MnO2-based electrodes and their application in

The charge storage mechanism of MnO2 was first imitated from RuO2 and Zn/MnO2 cells in acidic electrolytes [9], which involve a reversible insertion/extraction of protons into/from MnO2 with a concomitant Mn3+–Mn4+ transition. However, the specific capacitance of MnO2 is found to be related to the species and concentrations of the alkali metal cations, such as Li+

based on the chemisorption of alkali metal cations on the surface of MnO2 was proposed and

workers developed a new complementary storage mechanism, that is, the alkali metal cations

( ) <sup>2</sup> ( ) MnO M e MnOOM surface surface

represents alkali metal cations, including Li+

can insert/extract into/from MnO2, as shown in Eq. (2) [11]:

, regardless of the pH value [10, 11]. Consequently, a more rational mechanism

+ - + +® (1)

, and K+

. Later, Belanger and co-

, Na+

/V s), and structure flexibility [7, 8]. Moreover, owing to a

,

as a high-performance, safe, and a low-cost replacement.

**2. Charge storage mechanism of MnO2**

ionic diffusion constant (~10−13 cm2

44 Supercapacitor Design and Applications

electrodes.

supercapacitors.

, and K+

where M+

can be written as [3, 11]:

Na+

Recently, mixed-valent chemistry of metal oxides charge compensation mechanisms at oxygen centers and surfaces were proposed as a new charge storage mechanism. The mixedvalent MnOx nanostructures exhibited a high-specific capacitance of 2530 F/g in a twoelectrode configuration, about double of the theoretical value (1370 F/g) [14]. It is revealed that a large portion of charge compensation might be originated from the hole state redistribution toward the O sites, rather than merely reduction of Mn ions. In addition, the ionic defects (vacancies and misplaced ions), electronic defects (electron and holes), and structural defects (cavities, stacking faults, etc.) in the mixed-valent MnOx can also boost the kinetics of the surface redox reactions and the transport of charged species. These findings may offer critical insights into the fundamental understanding of the relationship between mixed-valent structures and electrochemical properties and to rational design of a newgeneration supercapacitors.

## **3. Synthesis methods**

#### **3.1. Chemical precipitation**

MnO2 nanostructures can be easily prepared *via* various chemical co-precipitation methods. The most commonly one is based on a simple redox reaction between MnO4 − and Mn2+ [4]:

$$2\text{MnO}\_4^- + \text{ 3Mn}^{2+} + 2\text{H}\_2\text{O} \rightarrow \text{ 5MnO}\_2 + 4\text{H}^+ \tag{3}$$

Amorphous hydrated MnO2 powders are the resulting products at room temperature. Increasing the reaction temperature not only leads to an increase in the degree of MnO2 crystallinity, but also generates different morphologies, such as nanorods, nanoflowers, and nanotubes [15]. Because of the strong oxidation of MnO4 − , a large number of reducing species, including acid, polymer monomers, carbon, etc., can be used to obtaining MnO2 nanostructures [16]. In addition, MnO2 can also be gained by changing other oxidants, such as Na2S2O8, KClO3, etc. These flexible derived reactions facilitate the formation of MnO2-carbon and MnO2-conducting polymer nanocomposites [17].

#### **3.2. Electrodeposition**

#### *3.2.1. Anodic electrodeposition*

Anodic electrodeposition of MnO2 nanostructures involves an oriented diffusion of Mn2+ species to the anode surface under an applied electric field and a subsequent electro-oxidation of the species into MnO2. The anodic deposition reaction can be expressed as [18]:

$$\text{Mn}^{2+} + 2\text{H}\_2\text{O} \rightarrow \text{MnO}\_2 + 4\text{H}^+ + 2\text{e}^- \tag{4}$$

Three different deposition techniques have been developed: i.e., potentiostatic deposition, potentiodynamic deposition, and galvanostatic deposition. The morphology can be well controlled by modifying the deposition parameters (such as applied potentials, applied currents, scan rates, electrolyte ingredients/concentrations, and pH value, etc.) and by tuning the physiochemical nature of substrates (such as a porous structure, a hydrophilic surface, etc.).

In addition to pure MnO2, conducting polymer-MnO2 nanocomposites can also be fabricated using solution precursors containing both Mn2+ species and monomers [19]. One-step or twostep electrochemical co-deposition has been explored to effectively prepare a nanocomposite. The co-electrodeposition process primarily involves the anodic deposition of MnO2 and the simultaneous electropolymerization of conducting polymers. The morphology of the nanostructured hybrids can be controlled by modifying the deposition parameters.

#### *3.2.2. Cathodic electrodeposition*

**3. Synthesis methods**

46 Supercapacitor Design and Applications

**3.1. Chemical precipitation**

**3.2. Electrodeposition**

etc.).

*3.2.1. Anodic electrodeposition*

MnO2 nanostructures can be easily prepared *via* various chemical co-precipitation methods.

Amorphous hydrated MnO2 powders are the resulting products at room temperature. Increasing the reaction temperature not only leads to an increase in the degree of MnO2 crystallinity, but also generates different morphologies, such as nanorods, nanoflowers, and

including acid, polymer monomers, carbon, etc., can be used to obtaining MnO2 nanostructures [16]. In addition, MnO2 can also be gained by changing other oxidants, such as Na2S2O8, KClO3, etc. These flexible derived reactions facilitate the formation of MnO2-carbon and

Anodic electrodeposition of MnO2 nanostructures involves an oriented diffusion of Mn2+ species to the anode surface under an applied electric field and a subsequent electro-oxidation

Three different deposition techniques have been developed: i.e., potentiostatic deposition, potentiodynamic deposition, and galvanostatic deposition. The morphology can be well controlled by modifying the deposition parameters (such as applied potentials, applied currents, scan rates, electrolyte ingredients/concentrations, and pH value, etc.) and by tuning the physiochemical nature of substrates (such as a porous structure, a hydrophilic surface,

In addition to pure MnO2, conducting polymer-MnO2 nanocomposites can also be fabricated using solution precursors containing both Mn2+ species and monomers [19]. One-step or twostep electrochemical co-deposition has been explored to effectively prepare a nanocomposite. The co-electrodeposition process primarily involves the anodic deposition of MnO2 and the simultaneous electropolymerization of conducting polymers. The morphology of the nano-

structured hybrids can be controlled by modifying the deposition parameters.

of the species into MnO2. The anodic deposition reaction can be expressed as [18]:

<sup>2</sup> Mn 2H O MnO 4H 2e 2 2

4 22 2MnO 3Mn 2H O 5MnO 4H - + <sup>+</sup> + +® + (3)

−

<sup>+</sup> + - + ® ++ (4)

−

, a large number of reducing species,

and Mn2+ [4]:

The most commonly one is based on a simple redox reaction between MnO4

2

nanotubes [15]. Because of the strong oxidation of MnO4

MnO2-conducting polymer nanocomposites [17].

Cathodic electrodeposition of MnO2 nanostructures can take place at the cathode surface *via* two pathways based on Mn species in solution under a negative potential. The first way is based on the electrochemical processes in aqueous solution containing Mn2+ species, which include water electrolysis and oxygen reduction to generate OH− , followed by the precipitation of metastable Mn(OH)2 and a further oxidation into MnO2. The other pathway is on basis of the reduction of Mn7+ species from MnO4 − on the cathode surface, following reaction (5) [20]. The MnO2 nanostructures can be shape-controlled by tuning deposition variables of current density, potential, solution concentration, and pH value.

$$2\text{MnO}\_4^- + 2\text{H}\_2\text{O} + 3\text{e}^- \rightarrow \text{MnO}\_2 + 4\text{OH}^-\tag{5}$$

#### **3.3. Electrostatic interaction assembly**

Electrostatic interaction assembly is an effective technique for fabricating nanocomposites *via* a self-assembly of oppositely charged materials. The strong electrostatic attractive interactions between positively charged species and negatively charged species can ensure the robustness of resulting materials. The binding intensity of the electric force is determined by the Coulomb's law [21]:

$$\mathbf{F} = \mathbf{q}\_\* \mathbf{q}\_- / \mathbf{r}^2 \tag{6}$$

where *F* is the electric force, *q+* and *q−* are the charges of the negative and positive species, and *r* is the distance between the two species. The main challenge is to controllably modify the target material surfaces with negatively/positively charged species, such as negatively charged poly(styrene sulfonate) (PSS), and positively charged aminopropyltrimethoxysilane (APTS) moieties and poly(diallyldimethylammonium chloride) (PDDA) [22].

## **4. Electrochemical performance of pure MnO2 nanostructures**

#### **4.1. Amorphous MnO2 powder electrodes**

The MnO2 powders prepared at low temperatures are generally amorphous or poorly crystalline in nature [23]. In addition, the as-prepared samples often contain a certain amount of hydrated content, which exists in the form of residual structural water and/or hydrated trivalent MnOOH. The hydrated MnO2 powders have large surface area because of ultrafine nanoparticle morphology. Although the amorphous structure can be maintained up to 300°C, the hydrated amount and the surface area decrease significantly with increasing annealing temperature and dwelling time [8, 24].

The electrochemical properties of the amorphous MnO2 powder electrode are closely related to their surface area and hydrated amount. According to the Eq. (1), an increase in the surface area is rather favorable for enhancing specific capacitance. The hydrated species can also modify the MnO2 surface affinity, which facilitates the electrolyte access and ion transport for enhanced performance. It was revealed that the amorphous MnO2 had the largest surface area and also possessed the best electrochemical performance. A dramatic decrease in the surface area and water content were observed after annealing treatments, which, in turn, degraded the specific capacitance.

#### **4.2. Crystalline MnO2 powder electrodes**

As mentioned earlier, MnO2 can exist in various crystal forms. The arrangement of [MnO6] basic building unit enables the construction of one-dimensional, two-dimensional, and threedimensional tunnel structures. The tunnel space, depending on the tunnel sizes, can accommodate cations, such as K+ , Na+ , Li+ , Ca+ , Mg+ , etc., which favors the energy storage according to Eq. (2). Recently, Ghodbane et al. has carried out a systematic study dealing with the microstructural effects on electrochemical properties in crystalline MnO2 [7]. As shown in **Figure 1**, the three-dimensional spinel *λ*-MnO2 exhibits the highest specific capacitance followed by the two-dimensional *δ*-MnO2 and, finally, the one-dimensional group shows the lowest values. To some extent, the specific capacitance increases with an increase in the tunnel size. One exception is the Ni-todorokite, which has a large tunnel size but a low-specific capacitance. The reason is the presence of very short and strong hydrogen bonds, which stabilize the hydrated cations inside the cavity and impede the ion diffusion through the tunnels. Additionally, it was found the specific capacitance did not depend on the specific surface area, and the ionic conductivity exerted great influence on the electrochemical performance. The comparison suggests the charge storage in MnO2 is a fundamentally Faradaic process.

**Figure 1.** Comparison of the specific capacitance, ionic conductivity, and surface area of various MnO2 nanostructures [7].

Hydrothermal or solvothermal methods are appropriate techniques to prepare MnO2 nanostructures, including nanorods, nanotubes, and nanowires, [4, 15]. The as-prepared products showed distinct morphologies, but the electrochemical results are highly scattered. There is no consistent relationship among the synthesis conditions, microstructures, and specific capacitance. In general, MnO2 with high crystallinity shows a low-specific surface area, and the cations in the tunnel structure are difficult to be extracted for TSM, thereby resulting in limited specific capacitance.

## **4.3. Thin-film MnO2 electrodes**

The electrochemical properties of the amorphous MnO2 powder electrode are closely related to their surface area and hydrated amount. According to the Eq. (1), an increase in the surface area is rather favorable for enhancing specific capacitance. The hydrated species can also modify the MnO2 surface affinity, which facilitates the electrolyte access and ion transport for enhanced performance. It was revealed that the amorphous MnO2 had the largest surface area and also possessed the best electrochemical performance. A dramatic decrease in the surface area and water content were observed after annealing treatments, which, in turn, degraded

As mentioned earlier, MnO2 can exist in various crystal forms. The arrangement of [MnO6] basic building unit enables the construction of one-dimensional, two-dimensional, and threedimensional tunnel structures. The tunnel space, depending on the tunnel sizes, can accom-

to Eq. (2). Recently, Ghodbane et al. has carried out a systematic study dealing with the microstructural effects on electrochemical properties in crystalline MnO2 [7]. As shown in **Figure 1**, the three-dimensional spinel *λ*-MnO2 exhibits the highest specific capacitance followed by the two-dimensional *δ*-MnO2 and, finally, the one-dimensional group shows the lowest values. To some extent, the specific capacitance increases with an increase in the tunnel size. One exception is the Ni-todorokite, which has a large tunnel size but a low-specific capacitance. The reason is the presence of very short and strong hydrogen bonds, which stabilize the hydrated cations inside the cavity and impede the ion diffusion through the tunnels. Additionally, it was found the specific capacitance did not depend on the specific surface area, and the ionic conductivity exerted great influence on the electrochemical performance. The comparison suggests the charge storage in MnO2 is a fundamentally

**Figure 1.** Comparison of the specific capacitance, ionic conductivity, and surface area of various MnO2 nanostructures

, etc., which favors the energy storage according

, Mg+

the specific capacitance.

48 Supercapacitor Design and Applications

modate cations, such as K+

Faradaic process.

[7].

**4.2. Crystalline MnO2 powder electrodes**

, Na+ , Li+ , Ca+ Thin-film MnO2 electrodes have been massively explored for deepening fundamental studies and finding potential applications as microscale energy storage devices. In this regard, a number of MnO2 thin-film electrodes have been directly prepared on a current collector, such as metal foils, carbon cloth, etc., through anodic/cathodic electrodeposition. In terms of the charge storage mechanism, the most desirable morphology should be of three-dimensional and porous nanoarchitectures with good connection to a highly conductive substrate. Therefore, much effort has focused on the morphology-controlled growth of porous MnO2 nanostructures, with a purpose of obtaining more accessible electroactive sites and short ion diffusion pathways. This can be achieved by controlling the deposition conditions and using porous templates. For example, galvanostatic, or potentiostatic techniques were employed to fabricate porous fibrous network, nanorods, and nanosheets [25–27].

As for anodic electrodeposition, chemically stable metal (e.g., Ti or Ta foils, etc.) and carbon (graphites, carbon fabric, etc.) substrates should be employed to address the problem of anodic oxidation and dissolution of metal substrates (e.g., stainless steel). In sharp contrast, cathodic deposition is a good method to break this limit. Electrodeposition from a dilute NaMnO4 − solution (0.02 M) produced a typical surface morphology of fibrous film with a birnessite-type crystal structure [28]. Compared to the pure MnO2 powders in Sections 4.1 and 4.2, the MnO2 thin films can deliver much higher specific capacitance.

#### **4.4. Summary**

Pure MnO2 nanostructures with different crystal phases and morphologies can be easily prepared using solution-based methods. Unfortunately, the overall specific capacitances of the amorphous and crystalline MnO2 powders are mostly ranging from 100 to 250 F/g, which are far from satisfaction. According to the charge storage mechanisms, the Faradaic reactions of MnO2 only occur at the surface or subsurface within tens of nanometers. Thus, to maximize the high electrochemical utilization of MnO2, the thickness of MnO2 should be kept in a nanoscale range. This is confirmed by a high-specific capacitance of 300–800 F/g when MnO2 thin films were deposited on the conductive substrates. The specific capacitance can even reach the theoretical value when thickness of MnO2 film is only several nanometers. The reason can be rationalized by the nanostructured MnO2 shortening both ionic and electronic transport/ diffusion distances for faster reaction kinetics. However, the thin films are limited by their low mass loading, which cannot meet large scale practical applications.

## **5. Electrochemical performance of MnO2-based nanocomposites**

#### **5.1. MnO2-carbon nanocomposites**

Over the past decades, carbon nanostructures have received intensive attention in various communities of material science and chemistry. First, carbon nanostructures are the most effective electrode material candidates for EDLCs owing to the versatile size dimensionality, high surface area, good electrical conductivity, and strong structural stability. Nevertheless, carbon-based electrodes are limited by their low specific capacitance. To break this limitation, carbon structures often serve as favorable supports for MnO2 by making complementary functions. In this section, we emphasize on the composites with one-dimensional carbon nanotubes (CNTs) two-dimensional, graphene, and three-dimensional mesoporous carbon [29].

## *5.1.1. MnO2-CNTs*

CNTs exhibit fascinating physicochemical properties of high electrical conductivity, high surface area, good chemical stability, and strong mechanical strength [29]. CNTs are increasingly employed as substrates for anchoring pseudocapacitive MnO2, to form nanocomposites. The nanocomposites can utilize the large pseudocapacitance of MnO2 and the conducting and mechanical merits from CNTs. Moreover, the nanocomposites could inherit the well-developed CNTs networks with open mesopores. Consequently, the MnO2-CNTs nanocomposites are expected to exhibit superior electrochemical properties. The key issue to obtain highperformance MnO2/CNT nanocomposites is how to uniformly deposit nanostructured MnO2 onto CNTs. To ensure the uniformity, the CNT powders should be homogeneously dispersed in a solution by means of surface modification and/or ultrasonic treatments. The threedimensional CNT network in the suspension can combine MnO2 nanostructures (e.g., nanorods or nanoclusters) using various chemical or electrochemical deposition methods. The most favorable method is based on a spontaneously redox reaction of MnO4 − and carbon, enabling *in situ* deposition of MnO2 on CNTs [30–32]. **Figure 2a** shows a typical core-sheath hierarchy architecture constructing by MnO2 nanostructures surrounding the CNT surfaces [33]. Under assistances of some modified techniques, such as the hydrothermal method, the refluxing, and the microwave irradiation, a variety of MnO2/CNT nanocomposites with different mass loadings and morphologies (e.g., nanoflakes and nanorods) of MnO2 have been demonstrated to exhibit improved electrochemical performance.

In addition to the surface deposition, MnO2 nanoparticles can also be embedded into the welldefined CNT nanochannels, as shown in **Figure 2b** [34]. The confined MnO2 are likely to exist in a more reduced state, and accordingly, the pseudocapacitive performance of MnO2 can be improved by modifying the redox couples of Mn3+/Mn4+. The specific capacitance normalized by MnO2 is almost close to the theoretical value (1250 F/g). These intriguing results may provide new insights into the configuration of MnO2 with nanocarbon for developing highperformance electrode materials.

Engineering Nanostructured MnO2 for High Performance Supercapacitors http://dx.doi.org/10.5772/65008 51

**Figure 2.** TEM (Transmission Electron Microscopy) images of (a) MnO2-out-CNT and (b) MnO2-in-CNT nanocomposites.

#### *5.1.2. MnO2-graphene*

**5. Electrochemical performance of MnO2-based nanocomposites**

Over the past decades, carbon nanostructures have received intensive attention in various communities of material science and chemistry. First, carbon nanostructures are the most effective electrode material candidates for EDLCs owing to the versatile size dimensionality, high surface area, good electrical conductivity, and strong structural stability. Nevertheless, carbon-based electrodes are limited by their low specific capacitance. To break this limitation, carbon structures often serve as favorable supports for MnO2 by making complementary functions. In this section, we emphasize on the composites with one-dimensional carbon nanotubes (CNTs) two-dimensional, graphene, and three-dimensional mesoporous carbon

CNTs exhibit fascinating physicochemical properties of high electrical conductivity, high surface area, good chemical stability, and strong mechanical strength [29]. CNTs are increasingly employed as substrates for anchoring pseudocapacitive MnO2, to form nanocomposites. The nanocomposites can utilize the large pseudocapacitance of MnO2 and the conducting and mechanical merits from CNTs. Moreover, the nanocomposites could inherit the well-developed CNTs networks with open mesopores. Consequently, the MnO2-CNTs nanocomposites are expected to exhibit superior electrochemical properties. The key issue to obtain highperformance MnO2/CNT nanocomposites is how to uniformly deposit nanostructured MnO2 onto CNTs. To ensure the uniformity, the CNT powders should be homogeneously dispersed in a solution by means of surface modification and/or ultrasonic treatments. The threedimensional CNT network in the suspension can combine MnO2 nanostructures (e.g., nanorods or nanoclusters) using various chemical or electrochemical deposition methods. The most

*in situ* deposition of MnO2 on CNTs [30–32]. **Figure 2a** shows a typical core-sheath hierarchy architecture constructing by MnO2 nanostructures surrounding the CNT surfaces [33]. Under assistances of some modified techniques, such as the hydrothermal method, the refluxing, and the microwave irradiation, a variety of MnO2/CNT nanocomposites with different mass loadings and morphologies (e.g., nanoflakes and nanorods) of MnO2 have been demonstrated

In addition to the surface deposition, MnO2 nanoparticles can also be embedded into the welldefined CNT nanochannels, as shown in **Figure 2b** [34]. The confined MnO2 are likely to exist in a more reduced state, and accordingly, the pseudocapacitive performance of MnO2 can be improved by modifying the redox couples of Mn3+/Mn4+. The specific capacitance normalized by MnO2 is almost close to the theoretical value (1250 F/g). These intriguing results may provide new insights into the configuration of MnO2 with nanocarbon for developing high-

−

and carbon, enabling

favorable method is based on a spontaneously redox reaction of MnO4

to exhibit improved electrochemical performance.

performance electrode materials.

**5.1. MnO2-carbon nanocomposites**

50 Supercapacitor Design and Applications

[29].

*5.1.1. MnO2-CNTs*

Graphene is a hottest star material in recent years due to its attractive characteristics of high electrical conductivity, good mechanical flexibility, high theoretical surface area (2600 m2 /g), and high thermal and chemical stability [35]. To exploit the potential of graphene-MnO2 nanomaterials for supercapacitors, the main challenge is to prevent graphene nanosheets from restacking during material synthesis and cycling operation. MnO2 nanostructures can act as interlayer spacers to effectively suppress the graphene restacking, thus facilitating fast ion diffusion/transport within the electrode materials. Moreover, the electrical intimate interaction of MnO2 and graphene can boost interfacial charge transfer to ensure rapid redox reactions of MnO2. In addition, graphene is capable of playing a "flexible confinement" role to enwrap MnO2 nanostructures through preventing inner MnO2 from electrochemical dissolution. Therefore, a strong synergetic effects benefiting from the enhanced conductivity, increased interfacial area as well as reinforced structural stability can be yielded in MnO2-graphene nanocomposites. To date, a great number of studies are devoting to (i) addressing the synthesis complexity and scalability, (ii) tailoring MnO2 nanostructures with a desired morphology and mass loadings in between graphene nanosheets, and (iii) improving electrical and mechanical connections between graphene and MnO2.

There are two favorable approaches for the controlled fabrication of the target MnO2/graphene nanocomposites. The first one is based on an *in situ* redox reaction between KMnO4 and graphene or GO, as illustrated in **Figure 3a** [36]. The rich oxygen-containing functionalities of graphene layers ensure a high-MnO2 mass loading. Functionalization of graphene, such as nitrogen doping, has become a key-enabling technology to improve the activities of graphene by increasing its conductivity and surface active sites. The other approach is a solution-based chemical mixing of separate MnO2 nanostructures with graphene, i.e., the formation of MnO2 is independent on graphene. Based on the electrostatic interaction assembly method, MnO2 nanostructures can be uniformly incorporated and strong anchored in between the graphene nanosheets [22, 37]. **Figure 3b** and **c** exemplify a graphene-wrapped MnO2 nanocomposite by co-assembling positively charged honeycomb-like MnO2 nanospheres and negatively charged graphene nanosheets [22]. It is noted that the morphology, crystallinity, and mass loading of the nano-MnO2 in the nanocomposites have great influence on the performance. In addition, controlling oxygen content of graphene plays a critical role in the fabrication of high-performance nanocomposites. It is believed that a simple, low-cost, and eco-friendly method is more favorable for controllably and scale-up fabricating MnO2/graphene nanocomposites with high performance.

**Figure 3.** (a) Schematic illustration for redox deposition and charge storage process of MnO2/graphene nanocomposite. (b) TEM image of MnO2 honeycombs/graphene. (c) Schematic illustration for honeycomb-like MnO2 on graphene through electrostatic assembly.

#### *5.1.3. MnO2/carbon nanotube/graphene*

CNTs could serve as a spacer to prevent adjacent graphene nanosheets from restacking and as conductive networks to accelerate electron transport of hybrid carbon composites. Recently, a few studies have exploited MnO2/CNT/graphene nanocomposites for high performance supercapacitors. The highly porous and conductive CNT/graphene composites can offer fast electronic and ionic channels for reversible Faradic redox reaction of MnO2 nanostructures. MnO2-coated CNTs were intercalated in between graphene nanosheets to form hierarchical nanocomposites [38]. The core step is the creation of positively charged MnO2/CNT functionalized with PDDA, which drives an electrostatical self-assembly with highly negatively charged graphene nanosheets. In this architecture, MnO2/CNT is able to effectively prevent graphene nanosheets from severe agglomeration. In addition to solution-based chemical assembly, chemical vapor deposition (CVD) method was recently utilized to grow CNT spacers in between graphene nanosheets [39]. The highly conductive network and the porous architecture enable the nanocomposites to exhibit an improvement in the specific capacitance.

## *5.1.4. MnO2-mesoporous carbon*

nanostructures can be uniformly incorporated and strong anchored in between the graphene nanosheets [22, 37]. **Figure 3b** and **c** exemplify a graphene-wrapped MnO2 nanocomposite by co-assembling positively charged honeycomb-like MnO2 nanospheres and negatively charged graphene nanosheets [22]. It is noted that the morphology, crystallinity, and mass loading of the nano-MnO2 in the nanocomposites have great influence on the performance. In addition, controlling oxygen content of graphene plays a critical role in the fabrication of high-performance nanocomposites. It is believed that a simple, low-cost, and eco-friendly method is more favorable for controllably and scale-up fabricating MnO2/graphene nanocomposites with high

**Figure 3.** (a) Schematic illustration for redox deposition and charge storage process of MnO2/graphene nanocomposite. (b) TEM image of MnO2 honeycombs/graphene. (c) Schematic illustration for honeycomb-like MnO2 on graphene

CNTs could serve as a spacer to prevent adjacent graphene nanosheets from restacking and as conductive networks to accelerate electron transport of hybrid carbon composites. Recently, a few studies have exploited MnO2/CNT/graphene nanocomposites for high performance supercapacitors. The highly porous and conductive CNT/graphene composites can offer fast electronic and ionic channels for reversible Faradic redox reaction of MnO2 nanostructures. MnO2-coated CNTs were intercalated in between graphene nanosheets to form hierarchical nanocomposites [38]. The core step is the creation of positively charged MnO2/CNT function-

performance.

52 Supercapacitor Design and Applications

through electrostatic assembly.

*5.1.3. MnO2/carbon nanotube/graphene*

Compared to activated carbon, ordered mesoporous carbons (OMC) possess uniform mesopores of several nanometers in diameter. The presence of the ordered mesopores is beneficial for electrolyte wetting and rapid ionic motion, which could address the rate-limiting issue of supercapacitors. The favorable mesostructure also enables OMC to be an ideal host material for compositing MnO2 nanostructures. However, the encapsulated MnO2 may result in disappearance of uniform mesopores and generate micropores between the nanoparticles, which prevent the mass transfer of electrolyte ions and also the formation of a double-layer required for high performance supercapacitors [40]. Hence, controlled growth of the nanoparticles homogeneously within the mesopores is crucial.

**Figure 4.** (a) TEM images of MnO2/CMK-3. (b) TEM image of highly graphitic carbon-tipped MnOx/mesoporous carbon/MnOx hybrid nanowires.

**Figure 4a** exhibits a MnO2/OMC composite prepared using an *in situ* redox reaction between KMnO4 and the mesoporous carbon [41]. MnO2 nanostructures are uniformly incorporated into the mesoporous carbon wall and the ordered mesopore structures are well-preserved. It is anticipated that OMC hosts are capable of holding the advantage of providing fast ion transport pathways for a high-rate power delivery; however, the rate capability of the composites is scarcely evaluated. This may be attributed to the relatively low electrical conductivity of OMC because the carbonization temperature is generally no more than 1000°C. This factor would become a bottle neck that would kinetically limit the charge transfer process. In combination of the favorable mesoporous structure and enhanced electrical conductivity, a novel one-dimensional, highly graphitic carbon-tipped MnOx/mesoporous carbon/MnOx (MMCM) hybrid nanostructure was demonstrated as a high-performance electrode material [42]. The unique TEM image of as-prepared hybrid nanowires are illustrated in **Figure 4b**, which is different from traditional structures with MnO2 deposited on the surfaces of carbon materials. Such a highly graphitic carbon-tipped mesoporous carbon shell provides efficient channels for ion transport into the core MnOx and improves electrical conductivity for electron transfer. The fascinating results offer a new direction on the design of ideal electrode materials with a high-specific capacitance, an excellent rate capability, and a long-term cyclability.

#### **5.2. MnO2-conducting polymer nanocomposites**

Conducting polymers are another class of pseudocapacitive materials due to their fast and reversible doping/undoping kinetics, ease in preparation, environmental stability, and anticorrosion purposes [43]. The conducting polymers could provide good electronic conductivity for MnO2, and in turn, MnO2 offers a solid support and a percolated electrical conducting pathway by interlinking the polymer chains, thus improving charge exchange efficiency and stability during redox cycling. In combination of high pseudocapacitance of both components, a strong synergistic effect is expected to integrate into a composite of MnO2/conducting polymers.

#### *5.2.1. MnO2-polyainiline (PANi)*

Polyaniline (PANi) is one of the most used conducting polymers due to its high doping level, good electrical conductivity, and environmental stability. Theoretical specific capacitance of PANi reaches 750 F/g, which makes it being widely used to combine various MnO2 [56]. KMnO4 is the common oxidant to polymerize aniline monomers into PANi accompanying with an instantaneous formation of MnO2:

$$\text{KMnO}\_4 + \text{aniline} \rightarrow \text{MnO}\_2 + \text{PANN} \tag{7}$$

This method leads to a good contact at an inter-molecule level between each component. It is demonstrated mesoporous MnO2/PANi hollow spheres using an interfacial synthesis technique (**Figure 5a**) [16]. The coupling reaction was carried out at the organic/water interfacial region, favoring the self-assembly of the composites with high surface area, uniform pore-size distribution, and hierarchical architecture. In addition to the aniline monomers, PANi nanostructures were used as reactive templates to reduce KMnO4, with MnO2 nanostructures depositing onto the PANi surface. **Figure 5b** shows ultrathin MnO2 nanorods grown on surfaces of conducting polymer nanofibers (PANi, polypyrrole (PPy), and PEDOT (poly(3,4 ethylenedioxythiophene)) by simply soaking the nanofibers in a KMnO4 aqueous solution [43]. Furthermore, owing to the high chemical oxidation potential of MnO2 (1.23 V) in an acidic condition, the aniline monomers can be polymerized with a simultaneous reduction of MnO2 into soluble Mn2+ ions, which provides a rational chemical method for *in situ* dispersing MnO2 on PANi supports [44]. Owing to the reactive template nature, morphologies of the nanocomposites can be shaped by the pristine MnO2, and the PANi content can be controlled by the polymerization time or the reaction ratio of aniline/MnO2.

**Figure 5.** TEM images ofMnO2/PANi (a) hollow spheres and (b) nanofibers.

## *5.2.2. MnO2-polypyrrole (PPy)*

novel one-dimensional, highly graphitic carbon-tipped MnOx/mesoporous carbon/MnOx (MMCM) hybrid nanostructure was demonstrated as a high-performance electrode material [42]. The unique TEM image of as-prepared hybrid nanowires are illustrated in **Figure 4b**, which is different from traditional structures with MnO2 deposited on the surfaces of carbon materials. Such a highly graphitic carbon-tipped mesoporous carbon shell provides efficient channels for ion transport into the core MnOx and improves electrical conductivity for electron transfer. The fascinating results offer a new direction on the design of ideal electrode materials with a high-specific capacitance, an excellent rate capability, and a long-term cyclability.

Conducting polymers are another class of pseudocapacitive materials due to their fast and reversible doping/undoping kinetics, ease in preparation, environmental stability, and anticorrosion purposes [43]. The conducting polymers could provide good electronic conductivity for MnO2, and in turn, MnO2 offers a solid support and a percolated electrical conducting pathway by interlinking the polymer chains, thus improving charge exchange efficiency and stability during redox cycling. In combination of high pseudocapacitance of both components, a strong synergistic effect is expected to integrate into a composite of MnO2/conducting poly-

Polyaniline (PANi) is one of the most used conducting polymers due to its high doping level, good electrical conductivity, and environmental stability. Theoretical specific capacitance of PANi reaches 750 F/g, which makes it being widely used to combine various MnO2 [56]. KMnO4 is the common oxidant to polymerize aniline monomers into PANi accompanying with

This method leads to a good contact at an inter-molecule level between each component. It is demonstrated mesoporous MnO2/PANi hollow spheres using an interfacial synthesis technique (**Figure 5a**) [16]. The coupling reaction was carried out at the organic/water interfacial region, favoring the self-assembly of the composites with high surface area, uniform pore-size distribution, and hierarchical architecture. In addition to the aniline monomers, PANi nanostructures were used as reactive templates to reduce KMnO4, with MnO2 nanostructures depositing onto the PANi surface. **Figure 5b** shows ultrathin MnO2 nanorods grown on surfaces of conducting polymer nanofibers (PANi, polypyrrole (PPy), and PEDOT (poly(3,4 ethylenedioxythiophene)) by simply soaking the nanofibers in a KMnO4 aqueous solution [43]. Furthermore, owing to the high chemical oxidation potential of MnO2 (1.23 V) in an acidic condition, the aniline monomers can be polymerized with a simultaneous reduction of MnO2 into soluble Mn2+ ions, which provides a rational chemical method for *in situ* dispersing MnO2 on PANi supports [44]. Owing to the reactive template nature, morphologies of the

KMnO aniline MnO PANi 4 2 + ®+ (7)

**5.2. MnO2-conducting polymer nanocomposites**

mers.

*5.2.1. MnO2-polyainiline (PANi)*

54 Supercapacitor Design and Applications

an instantaneous formation of MnO2:

PPy is another low-cost conducting polymer that can be used as promising electrode materials for supercapacitors. It is noteworthy that, unlike PANi showing good supercapacitive behavior in an acidic electrolyte, PPy holds the capability to possess excellent electrochemical properties in neutral electrolytes, which are highly compatible with MnO2. KMnO4 is a proper oxidant to chemically polymerize pyrrole monomers into PPy in relation to the formation of MnO2 by reducing MnO4 − [45]. The dispersed MnO2 nanoparticles adhered to PPy chains increase the surface area and retard the structural deterioration of PPy backbones. As mentioned earlier, MnO2 nanostructures could serve as solid reactive templates for *in situ* polymerization of pyrrole monomers in an acidic environment, thus forming MnO2/PPy nanocomposites without using any surfactants and/or additional oxidants. To this end, unique one-dimensional coaxial MnO2/PPy nanotubes were developed [46, 47].

## *5.2.3. MnO2-polythiophone (PTh)*

Lu and Zhou reported a one-pot interfacial synthesis to fabricate mesoporous MnO2/PTh nanocomposite having uniform submicron-sphere/nanosheet hierarchical structures [48]. The significant roles of PTh in the nanocomposite are emphasized in terms of their functions on enhancing the electrical conductivity and constraining the dissolution of MnO2 component.

#### **5.3. MnO2-carbon-conducting polymer ternary nanocomposites**

In order to fully exploit the advantages of different carbon and/or conducting polymers, MnO2 can be combined with the dual-supports to make a maximum use of its utilization. Carbon nanostructures (e.g., CNTs) can facilitate easy electrolyte accessibility and fast electron transport into the bulk electrode materials, meanwhile the conducting polymers can contribute to more charge storage and better inter-particle connectivity to MnO2. This ternary design approach is expected to show improved performance, since each component in the ternary nanocomposite serves its unique and desired functions to collectively optimize their electrochemical properties.

**Figure 6.** (a, b) Schematic illustration of synthesizing MnO2/conducting polymers on CNTs-PSS. (c) TEM image of MnO2/PEDOT/CNTs.

As for the MnO2-CNT-conducting polymer nanocomposites, **Figure 6a** presents a typical synthesis route [21]. The CNT surface was modified by wrapping negatively charged PSS, which electrostatically attract Mn2+ species and monomers (i.e., aniline, pyrrole, and EDOT). The large number of surface sites (–SO3 − ) have strong interactions of growing nucleus, which facilitate an ordered growth of nanostructured MnO2 and conducting polymers and further hinder the inter-particle agglomeration. The molecular level dispersion of MnO2 in the CNT networks and conducting polymer matrixes results in a strong synergistic interaction. Electrochemical polymerization can also be used to deposit nanometer-thick outer layer of PEDOT onto a coaxial MnO2/CNT composite [49]. Moreover, sonochemical processing was employed to wrap a water-soluble conducting polymer of PEDOT-PSS onto a binary composite of hierarchical MnO2 nanospheres/CNTs, as shown in **Figure 6b** and **c** [50]. PEDOT-PSS not only served as a dispersant to stabilize the composite suspension, thus facilitating the electrode fabrication without use of binders, but also offered a good inter-particle connectivity between MnO2 and CNTs.

#### **5.4. Summary**

Carbon nanostructures offer excellent substrates for enhancing the electrochemical performance of MnO2. The carbon can be dispersed homogeneously into a solution, facilitating a uniform deposition of MnO2 nanodeposits onto the carbon surface through a solution-based processing. The high surface area of carbon supports also provides a large number of anchoring sites for MnO2 formation and enlarges the contact interfaces between MnO2 and electrolyte. It is worthy to note that a porous carbon having well-ordered pore channels can permit fast electrolyte transport, meanwhile, a well-graphitic carbon with enhanced electrical conductivity can enable rapid charge collection/transfer under a high-rate operation. Conducting polymers can work in good synergy with MnO2 to capture a maximum electrochemical harvesting from the large pseudocapacitances. A rational design strategy is necessary to control the MnO2 dispersion into conducting polymers by tuning the synthesis conditions. For the *in situ* redox deposition, the mass loading of MnO2 is extremely limited because an overoxidation of conducting polymers would destroy their π-conjugated structure, and accordingly, substantially lose their conductivities and electroactivities. A better synthesis alternative is to use MnO2 nanostructures as reactive templates for polymerizing polymer monomers surrounding its outmost surfaces. The synchronous formation of MnO2 and conducting polymers in the chemical co-precipitation method facilitates a molecular level interaction with each other, but the electrochemical performance depends on the morphologies of the nanocomposites. Interfacial synthesis is an optional approach to prepare a composite with mesoporous structure.

Ternary nanocomposites are capable of maximizing the desirable functions of each component. The synergistic contribution results in great capacitance enhancement, high-rate delivery and better cycling performance. Furthermore, it is important to bridge MnO2 and carbon to electrolytes with a large interaction area by controlling the thickness of conducting polymer layers and their spatial distribution in the ternary structure.

## **6. Asymmetric supercapacitors based on MnO2 nanostructures**

#### **6.1. Fabrication principles**

transport into the bulk electrode materials, meanwhile the conducting polymers can contribute to more charge storage and better inter-particle connectivity to MnO2. This ternary design approach is expected to show improved performance, since each component in the ternary nanocomposite serves its unique and desired functions to collectively optimize their electro-

**Figure 6.** (a, b) Schematic illustration of synthesizing MnO2/conducting polymers on CNTs-PSS. (c) TEM image of

As for the MnO2-CNT-conducting polymer nanocomposites, **Figure 6a** presents a typical synthesis route [21]. The CNT surface was modified by wrapping negatively charged PSS, which electrostatically attract Mn2+ species and monomers (i.e., aniline, pyrrole, and EDOT).

facilitate an ordered growth of nanostructured MnO2 and conducting polymers and further hinder the inter-particle agglomeration. The molecular level dispersion of MnO2 in the CNT networks and conducting polymer matrixes results in a strong synergistic interaction. Electrochemical polymerization can also be used to deposit nanometer-thick outer layer of PEDOT onto a coaxial MnO2/CNT composite [49]. Moreover, sonochemical processing was employed to wrap a water-soluble conducting polymer of PEDOT-PSS onto a binary composite of hierarchical MnO2 nanospheres/CNTs, as shown in **Figure 6b** and **c** [50]. PEDOT-PSS not only served as a dispersant to stabilize the composite suspension, thus facilitating the electrode fabrication without use of binders, but also offered a good inter-particle connectivity be-

Carbon nanostructures offer excellent substrates for enhancing the electrochemical performance of MnO2. The carbon can be dispersed homogeneously into a solution, facilitating a

) have strong interactions of growing nucleus, which

−

chemical properties.

56 Supercapacitor Design and Applications

MnO2/PEDOT/CNTs.

tween MnO2 and CNTs.

**5.4. Summary**

The large number of surface sites (–SO3

To make an asymmetric supercapacitor that can operated stably in a wide potential range, the critical issue is to couple electrode materials with different high overpotentials for hydrogen or oxygen evolutions [51]. It is known that MnO2 nanostructures normally work in the electrochemical potential window of 0.0–1.0 V (vs. SCE). Therefore, the coupled electrode materials should have a potential window in a range far beyond 0.0–1.0 V. Since the first report of MnO2|| activated carbon asymmetric supercapacitor by Hong et al. in 2002 [52], carbon materials have attracted massive attention as coupled electrodes, because they can operate in a wide potential window from −1.0 to 0.1 V vs. SCE (Saturated Calomel Electrode) [17]. A wide operating potential window of 2.0 V is yielded in the assembled asymmetric supercapacitor in a Na2SO4 aqueous electrolyte. The asymmetric supercapacitor could even work in a working voltage as high as 2.3 V after elaborating the electrode materials (e.g., mass ratio) and cell assembly (e.g., oxygen expelling). Such a high-cell voltage is comparable to a commercially available symmetric supercapacitors using organic electrolyte (2.5 V) [51]. Consequently, the energy density of asymmetric supercapacitors is exceptionally high, reaching over 20 Wh/kg, almost one order of magnitude higher than that of the aqueous-based symmetric supercapacitor and also greater than that of organic-based SSCs (17–18 Wh/kg).

In addition to the carbon materials, pseudocapacitive materials, such as metal oxides and conducting polymers (e.g., PANi, PPy, and PEDOT), can be used as coupled electrodes for asymmetric supercapacitors. It is noted that the electrochemical oxidation/reduction potential is essentially associated with the work function of a metal oxide [53]. A large work function difference of two metal oxides provide an opportunity of enlarging operating voltage larger than the dissociation energy of an aqueous electrolyte because the water decomposition is kinetically limited by hydrogen and oxygen evolution reaction on the surface of metal oxides. Hence, a metal oxide with a largest work function difference from MnO2 is more appropriate for asymmetric supercapacitor assembly to maximizing cell voltage in neutral aqueous electrolytes.

Moreover, another important technological issue is to balance the charge (*Q*) that stored at positive and negative electrodes. The charge stored by each electrode is linearly proportional to the specific capacitance (*C*), the potential window (*ΔU*), and the mass (*m*) of an electrode material, i.e., Q = C × ΔU × m. In order to obtain charge balance, i.e., *Q+ = Q*−, the optimal mass ratio of the positive electrode (*m*+) against negative electrode (*m*−) can be estimated according to Eq. (8) [17]. It should be noted that the optimal mass ratio is responsible for an ideal capacitive behavior with a maximum cell voltage and a high-Coulombic efficiency.

$$\frac{m\_{+}}{m\_{-}} = \frac{C\_{-} \times \Delta U\_{-}}{C\_{+} \times \Delta U\_{+}} \tag{8}$$

#### **6.2. Asymmetric supercapacitor cells**

Carbon materials are of particular interest as negative electrodes for asymmetric supercapacitors due to their characteristics of good EDLC, a high-overpotential for hydrogen evolution, excellent electrical conductivity for high power delivery, and good chemical inertness for long cycle life. In this regard, various carbon materials have been explored including graphene, CNTs, activated carbon nanofibers (ACNFs), carbon sphere (CS), ordered mesoporous carbon, etc., to couple MnO2.

Graphene-based materials have been received intensive attention for asymmetric supercapacitors. First, graphene itself exhibits outstanding physiochemical and capacitive properties. Second, the ultrathin two-dimensional structure of graphene provides a large surface area for anchoring MnO2 nanostructures [54, 55]. In addition, graphene-based materials can also be compacted into free-standing electrodes, which endow asymmetric supercapacitor with flexibility and lightweight. The energy and power values would be more fascinating at a device level taking the additives (binders and conductors) and the current collectors into account. **Figure 7** demonstrated porous hybrid thin-film structures of MnO2/graphene and graphene/Ag, which are highly flexible, mechanically robust, and ultra-lightweight using the ordinal filtration assembly method [55]. The flexible device extended the operating potential window to 1.8 V and showed an energy density of 50.8 Wh/kg.

Engineering Nanostructured MnO2 for High Performance Supercapacitors http://dx.doi.org/10.5772/65008 59

**Figure 7.** Illustrations of (a) MnO2/graphene and graphene/Ag freestanding films and (b) the as-fabricated asymmetric supercapacitor devices.

CNTs and carbon nanofiber fabrics were demonstrated to serve excellent one-dimensional negative materials for energy storage devices due to their high electrical conductivity, good mechanical stability, and flexibility [56–60]. On one hand, MnO2 nanostructures can be uniformly deposited surround the one-dimensional carbons and the resulting nanocomposites showed pronounced electrochemical properties. On the other hand, the porous one-dimensional carbons possess good electric double-layer capacitance, thus becoming highly promising negative candidates for asymmetric supercapacitors. The asymmetric supercapacitors assembled by one-dimensional carbons as the negative electrodes and MnO2/carbon nanocomposites as the positive electrode can be cycled reversibly in a high-voltage region of 0–2.0 V and exhibited a superior energy density of 30–35 Wh/kg in a neutral aqueous Na2SO4 electrolyte. In particular, the asymmetric supercapacitors based on electrospun carbon nanofibers are rather promising to meet the ever-increasing demands of high energy/power densities because of the scalability, easy-fabrication, and low-cost of the electrospinning technique [56].

In addition to the carbon materials, metal oxides and conducting polymers, such as SnO2, MoO3 PANi, PPy, and PEDOT, were also explored as negative materials to couple MnO2 nanostructures [51, 61]. MoO3 would be the most favorable coupled pair of MnO2 because of their largest work function difference [53]. To this end, MoO3/graphene composite and MnO2/ graphene nanocomposite were prepared to serve as the negative electrode and the positive electrode for the fabrication of an asymmetric supercapacitor. The as-assembled device can operated stably in a wide voltage window of 2.0 V, and more significantly, exhibited a highspecific capacitance of 307 F/g and a high-energy density of 42.6 Wh/kg.

### **7. Summaries and perspectives**

almost one order of magnitude higher than that of the aqueous-based symmetric supercapa-

In addition to the carbon materials, pseudocapacitive materials, such as metal oxides and conducting polymers (e.g., PANi, PPy, and PEDOT), can be used as coupled electrodes for asymmetric supercapacitors. It is noted that the electrochemical oxidation/reduction potential is essentially associated with the work function of a metal oxide [53]. A large work function difference of two metal oxides provide an opportunity of enlarging operating voltage larger than the dissociation energy of an aqueous electrolyte because the water decomposition is kinetically limited by hydrogen and oxygen evolution reaction on the surface of metal oxides. Hence, a metal oxide with a largest work function difference from MnO2 is more appropriate for asymmetric supercapacitor assembly to maximizing cell voltage in neutral aqueous

Moreover, another important technological issue is to balance the charge (*Q*) that stored at positive and negative electrodes. The charge stored by each electrode is linearly proportional to the specific capacitance (*C*), the potential window (*ΔU*), and the mass (*m*) of an electrode material, i.e., Q = C × ΔU × m. In order to obtain charge balance, i.e., *Q+ = Q*−, the optimal mass ratio of the positive electrode (*m*+) against negative electrode (*m*−) can be estimated according to Eq. (8) [17]. It should be noted that the optimal mass ratio is responsible for an ideal

capacitive behavior with a maximum cell voltage and a high-Coulombic efficiency.

*mC U mC U* +- - -+ +

Carbon materials are of particular interest as negative electrodes for asymmetric supercapacitors due to their characteristics of good EDLC, a high-overpotential for hydrogen evolution, excellent electrical conductivity for high power delivery, and good chemical inertness for long cycle life. In this regard, various carbon materials have been explored including graphene, CNTs, activated carbon nanofibers (ACNFs), carbon sphere (CS), ordered mesoporous carbon,

Graphene-based materials have been received intensive attention for asymmetric supercapacitors. First, graphene itself exhibits outstanding physiochemical and capacitive properties. Second, the ultrathin two-dimensional structure of graphene provides a large surface area for anchoring MnO2 nanostructures [54, 55]. In addition, graphene-based materials can also be compacted into free-standing electrodes, which endow asymmetric supercapacitor with flexibility and lightweight. The energy and power values would be more fascinating at a device level taking the additives (binders and conductors) and the current collectors into account. **Figure 7** demonstrated porous hybrid thin-film structures of MnO2/graphene and graphene/Ag, which are highly flexible, mechanically robust, and ultra-lightweight using the ordinal filtration assembly method [55]. The flexible device extended the operating

potential window to 1.8 V and showed an energy density of 50.8 Wh/kg.

´ D <sup>=</sup> ´ D (8)

citor and also greater than that of organic-based SSCs (17–18 Wh/kg).

electrolytes.

58 Supercapacitor Design and Applications

**6.2. Asymmetric supercapacitor cells**

etc., to couple MnO2.

MnO2-based electrode materials hold a particular prospect for future supercapacitor applications although there are still some obstacles that need to be conquered. To better utilize the electrochemical performance of MnO2, it is necessary to create more electrochemically active sites or reduce the ion/electron transport distance by modifying surface chemistry and structure of MnO2. One of the most effective approaches is incorporating nanoscaled MnO2 into a highly porous and electronically conductive framework, such as carbon scaffolds and conducting polymers, to form a hybrid material. Benefiting from the synergistic contribution from each component in the composite, the consistent progress in the research and development of MnO2-based nanocomposites have achieved many breakthroughs in terms of charge storage mechanism, smart design strategies, and technological innovations. It should be noted that more advanced characterization tools and new methodologies as well systematic studies are still required to deepen the fundamental understanding of material chemistry and electrode/electrolyte interface. The energy and power densities of supercapacitors can be significantly improved by one magnitude through a simple asymmetric configuration method. Although many limiting factors, such as optimal coupled electrode, high-cost issue, complex fabrication process, moderate cycle life, etc., are required to be conquered urgently, it is believed that the ongoing research would bring MnO2-based asymmetric supercapacitors to an acceptable level for practical applications.

## **Acknowledgements**

The author acknowledge the financial supports of this work by the National Natural Science Foundation of China (51402236), the Natural Science Foundation of Shaanxi Province (2015JM5180), the Research Fund of the State Key Laboratory of Solidification Processing (NWPU), China (Grant No.: 123-QZ-2015), and the Key Laboratory of New Ceramic and Fine Processing (Tsinghua University, KF201607).

## **Author details**

Jian-Gan Wang

Address all correspondence to: wangjiangan@nwpu.edu.cn

State Key Laboratory of Solidification Processing, Center for Nano Energy Materials, School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an, China

## **References**

[1] Salanne M, Rotenberg B, Naoi K, Kaneko K, Taberna PL, Grey CP, Dunn B, Simon P. Efficient storage mechanisms for building better supercapacitors. Nature Energy. 2016; 1: 16070. doi:10.1038/NENERGY.2016.70

[2] Lee HY, Goodenough JB. Supercapacitor behavior with KCl electrolyte. Journal of Solid State Chemistry. 1999; 144: 220–223. doi:10.1006/jssc.1998.8128

electrochemical performance of MnO2, it is necessary to create more electrochemically active sites or reduce the ion/electron transport distance by modifying surface chemistry and structure of MnO2. One of the most effective approaches is incorporating nanoscaled MnO2 into a highly porous and electronically conductive framework, such as carbon scaffolds and conducting polymers, to form a hybrid material. Benefiting from the synergistic contribution from each component in the composite, the consistent progress in the research and development of MnO2-based nanocomposites have achieved many breakthroughs in terms of charge storage mechanism, smart design strategies, and technological innovations. It should be noted that more advanced characterization tools and new methodologies as well systematic studies are still required to deepen the fundamental understanding of material chemistry and electrode/electrolyte interface. The energy and power densities of supercapacitors can be significantly improved by one magnitude through a simple asymmetric configuration method. Although many limiting factors, such as optimal coupled electrode, high-cost issue, complex fabrication process, moderate cycle life, etc., are required to be conquered urgently, it is believed that the ongoing research would bring MnO2-based asymmetric supercapacitors to

The author acknowledge the financial supports of this work by the National Natural Science Foundation of China (51402236), the Natural Science Foundation of Shaanxi Province (2015JM5180), the Research Fund of the State Key Laboratory of Solidification Processing (NWPU), China (Grant No.: 123-QZ-2015), and the Key Laboratory of New Ceramic and Fine

State Key Laboratory of Solidification Processing, Center for Nano Energy Materials, School of Materials Science and Engineering, Northwestern Polytechnical University, Xi'an, China

[1] Salanne M, Rotenberg B, Naoi K, Kaneko K, Taberna PL, Grey CP, Dunn B, Simon P. Efficient storage mechanisms for building better supercapacitors. Nature Energy. 2016;

an acceptable level for practical applications.

Processing (Tsinghua University, KF201607).

Address all correspondence to: wangjiangan@nwpu.edu.cn

1: 16070. doi:10.1038/NENERGY.2016.70

**Acknowledgements**

60 Supercapacitor Design and Applications

**Author details**

Jian-Gan Wang

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#### **Nanostructured Metal Oxides-Based Electrode in Supercapacitor Applications Nanostructured Metal Oxides-Based Electrode in Supercapacitor Applications**

Zhenjun Qi, Shihao Huang, Adnan Younis, Dewei Chu and Sean Li Zhenjun Qi, Shihao Huang, Adnan Younis, Dewei Chu and Sean Li

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/65155

#### **Abstract**

To overcome the obstacle of low energy density, one of the most intensive approaches is the development of new materials for supercapacitor electrodes. Most explored materials today are carbon particle materials, which have high surface areas for charge storage. But in spite of these large specific surface areas, the charges physically stored on the carbon particles in porous electrode layers are unfortunately limited. Regarding advanced supercapacitor electrodes, metal oxides are considered the most promising material for the next generation of supercapacitors owing to their unique physical and chemical properties. In this chapter, the rational design and fabrication of metal oxide nanostructures for supercapacitor applications are addressed.

**Keywords:** supercapacitor, metal oxide, nanostructure, nanocomposite, graphene

## **1. Introduction**

In recent years, supercapacitors (or ultracapacitors) have attracted significant attention as a versatile solution to meet the increasing demands of energy storage because of their fast power energy delivery, long lifecycle, high power density and reasonably high energy density which are able to fill in the gap between the batteries and the conventional capacitors [1–3].

The comparison of specific power and specific energy for different electrical energy storage devices is shown in the Ragone plot (**Figure 1**) [4]. The data illustrate that supercapacitors are able to store more energy than conventional capacitors and deliver more power than batteries. Owing to the different energy storage mechanism from conventional capacitors, the specific

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

energy of supercapacitors can be thousands of times higher than it of conventional capacitors by forming an electric double layer at the interface of the electrode and electrolyte to store energy [5]. Thus, the high surface area of the electrode can be adequately utilized to collect amounts of positively and negatively charged ions from electrolyte when storing energy.

**Figure 1.** Specific power and energy for various electrical energy storage devices show in Ragone plot [4].

On the other hand, as is well-known, rechargeable batteries mainly depend on chemical reactions to charge and discharge which significantly restrict their lifetime [6]. Compared to batteries, the energy storage process of supercapacitors is based on electrostatic storage in the electrical double layer and reversible faradaic redox reactions by means of electron charge transfer on the surface of electrode. Thus, supercapacitors are expected to have a capability of faster charge–discharge under high current and a longer cycle life than batteries because no or negligibly few chemical reactions are involved [7].

Even though major progress has been yielded in the theoretical and practical research and development of supercapacitors, few disadvantages of supercapacitors, including low energy density and high production cost, have been identified as major challenges for the furtherance of supercapacitors technologies [8].

To overcome the obstacle of low energy density, one of the most intensive approaches is the development of new materials for supercapacitor electrodes. Most explored materials today are carbon particle-based materials, which have high surface areas for charge storage [3]. But in spite of these large specific surface areas, the charges physically stored on the carbon particles in porous electrode layers unfortunately limiting their electrochemical properties. Supercapacitors of this kind, called electrical double-layer supercapacitors (EDLS), have a limited specific capacitance and relative low energy density [9]. Supercapacitors with electrochemically active materials (polymers and metal oxides) as electrodes involving fast and reversible faradaic reactions on electrodes are called faradaic supercapacitors (FS). It has been demonstrated that faradaic or hybrid double-layer supercapacitors can yield much higher specific capacitance and energy density. Thus, regarding advanced supercapacitor materials, metal oxides are considered the most promising materials for the next generation of supercapacitors [10].

**Figure 2.** Mechanisms of (a) pseudocapacitance and (b) hybrid capacitance [13].

energy of supercapacitors can be thousands of times higher than it of conventional capacitors by forming an electric double layer at the interface of the electrode and electrolyte to store energy [5]. Thus, the high surface area of the electrode can be adequately utilized to collect amounts of positively and negatively charged ions from electrolyte when storing energy.

**Figure 1.** Specific power and energy for various electrical energy storage devices show in Ragone plot [4].

negligibly few chemical reactions are involved [7].

of supercapacitors technologies [8].

68 Supercapacitor Design and Applications

On the other hand, as is well-known, rechargeable batteries mainly depend on chemical reactions to charge and discharge which significantly restrict their lifetime [6]. Compared to batteries, the energy storage process of supercapacitors is based on electrostatic storage in the electrical double layer and reversible faradaic redox reactions by means of electron charge transfer on the surface of electrode. Thus, supercapacitors are expected to have a capability of faster charge–discharge under high current and a longer cycle life than batteries because no or

Even though major progress has been yielded in the theoretical and practical research and development of supercapacitors, few disadvantages of supercapacitors, including low energy density and high production cost, have been identified as major challenges for the furtherance

To overcome the obstacle of low energy density, one of the most intensive approaches is the development of new materials for supercapacitor electrodes. Most explored materials today are carbon particle-based materials, which have high surface areas for charge storage [3]. But The capacitance of a metal oxide-based supercapacitor is determined by two storage principles, one is double-layer capacitance and another one is pseudocapacitance. The mechanism of double-layer capacitance is shown in **Figure 2a**, when capacitor charged, electrostatic storage achieved by separation of charge in a double layer at the interface between the surface of a conductive electrode and an electrolyte. As a result, mirror image of charge distribution of ions in opposite polarity, called double-layer, is formed. When capacitor discharged, ions return and distribute randomly in the electrolyte. For pseudocapacitance, it stores electrical energy electrochemically by means of reversible faradaic redox reactions on the surface of suitable electrode materials in an electrochemical capacitor with an electric double-layer [11]. It can be seen in **Figure 2b**, pseudocapacitance is accompanied with an electron charge-transfer between electrolyte and electrode coming from a de-solvated and adsorbed ion whereby only one electron per charge unit is participating. This faradaic charge transfer originates via a very fast sequence of reversible redox, electrosorption or intercalation processes. The adsorbed ion has no chemical reaction with the atoms of the electrode. No chemical bonds arise, and only a charge-transfer takes place [12]. Even though both double-layer capacitance and pseudocapacitance contribute indivisibly to the total capacitance of metal oxide supercapacitors, the latter can be 10–100 times higher than the former.

In general, metal oxide-based supercapacitors are able to possess higher specific capacitance and energy density than carbon materials and conducting polymer materials [14]. A series of metal oxides with high theoretical performances has been studied, such as RuO2, MnO2, NiO, Co3O4, V2O5, CuO and Fe3O4. **Table 1** [15] lists the theoretical capacitance of some typical metal oxides as well as the charge storage reactions. However, the practical supercpacitive properties of these metal oxides are far behind their theoretical values due to their low conductivities, poor long-term stability, low surface areas and porosity.


**Table 1.** Pseudocapacitance and conductivity of selected metal oxides.

In this chapter, several important factors affecting the electrochemical properties of metal oxide-based electrodes are discussed firstly. Then various methods to fabricate nanostructured metal oxide electrode are summarized. Finally, advanced metal oxide-carbon composite electrodes are further described.

## **2. Factors affecting the performance of metal oxide-based electrodes**

#### **2.1. Crystallinity**

The degree of crystallinity is one of pivotal factors affecting the pseudocapacitance of metal oxide materials. In general, an amorphous structure exhibits superior electrochemical performance than a well-crystallized structure due to the former can make the fast, continuous and reversible faradaic reaction take place not only on the surface but also inside of metal oxide particles [16]. This is because the amorphous structure with a highly porous morphology is benefit for ion accessibility and cation diffusion. In addition, these porous structures also result in a higher specific surface area which can support more redox reactions to enhance the specific capacitance. Nevertheless, it is well-known that the poorly crystallized metal oxide simultaneously leads to a lower electrical conductivity limiting its pseudocapacitance. Thus, it is necessary and a great challenge to explore the appropriate crystallinity with optimal conductivity and ionic transport.

One method is to improve the electric conductivity of the amorphous metal oxide. It has been reported that annealing can significantly affect the electrical conductivity. For example, in [17], MnOx annealed at 200°C exhibited a higher specific capacitance at high scan rate than those without treated. Another effective approach is to optimize the structure of crystallized metal oxide in order to provide appropriate tunnels for the intercalation of cations.

#### **2.2. Crystal structure**

In general, metal oxide-based supercapacitors are able to possess higher specific capacitance and energy density than carbon materials and conducting polymer materials [14]. A series of metal oxides with high theoretical performances has been studied, such as RuO2, MnO2, NiO, Co3O4, V2O5, CuO and Fe3O4. **Table 1** [15] lists the theoretical capacitance of some typical metal oxides as well as the charge storage reactions. However, the practical supercpacitive properties of these metal oxides are far behind their theoretical values due to their low conductivities,

**capacitance (F g−1)**

= NiOOH + e− 2584 0.01 to 0.32

1380 10−5 to 10−6

2120 10−4 to 10−2

3560 10−4 to 10−2

**Conductivity (S cm−1)**

poor long-term stability, low surface areas and porosity.

(M could be H+

(M could be H+

V2O5 + 4M+

NiO + OH−

Co3O4 + OH−

CoOOH + OH−

**Table 1.** Pseudocapacitance and conductivity of selected metal oxides.

MnO2 Na2SO4 MnO2 + M+

70 Supercapacitor Design and Applications

Na2SO4

NaOH

NaOH

electrodes are further described.

V2O5 NaCl,

NiO KOH,

Co3O4 KOH,

**2.1. Crystallinity**

tivity and ionic transport.

**Oxide Electrolyte Charge storage reaction Theoretical**

, Li+ , Na+ , K+ )

, Li+ , Na+ , K+ )

= MMnO2

= M2V2O5

+ H2O = 3CoOOH + e−

= CoO2 + H2O + e−

In this chapter, several important factors affecting the electrochemical properties of metal oxide-based electrodes are discussed firstly. Then various methods to fabricate nanostructured metal oxide electrode are summarized. Finally, advanced metal oxide-carbon composite

**2. Factors affecting the performance of metal oxide-based electrodes**

The degree of crystallinity is one of pivotal factors affecting the pseudocapacitance of metal oxide materials. In general, an amorphous structure exhibits superior electrochemical performance than a well-crystallized structure due to the former can make the fast, continuous and reversible faradaic reaction take place not only on the surface but also inside of metal oxide particles [16]. This is because the amorphous structure with a highly porous morphology is benefit for ion accessibility and cation diffusion. In addition, these porous structures also result in a higher specific surface area which can support more redox reactions to enhance the specific capacitance. Nevertheless, it is well-known that the poorly crystallized metal oxide simultaneously leads to a lower electrical conductivity limiting its pseudocapacitance. Thus, it is necessary and a great challenge to explore the appropriate crystallinity with optimal conduc-

+ e−

+ 4e−

The crystal structure has a significant influence on the pseudocapacitance of metal oxide because it plays a crucial role in determining the cations intercalation. For instance, crystallized manganese oxide has different crystalline structures, including α-, β-, γ-, δ- and λ-MnO2 shown in **Figure 3** [18]. It can be seen thatα-MnO2 forms 1D (2 × 2) and (1 × 1) tunnels; β-MnO2 forms a 1D (1 × 1) tunnel; γ-MnO2 is consist of 1D (1 × 2) and (1 × 1) tunnels; δ-MnO2 is a 2D layered structure; and λ-MnO2 is a three-dimensional (3D) spinel structure, respectively [11, 19]. It is reported by Brousse et al. [20] that α-MnO2 with a large tunnel size exhibited a relatively high specific capacitance of 110 F g−1 due to K+ cations could easily insert the tunnels. On the contrary, β-MnO2 with a narrow tunnel size which is smaller than K+ cations inhibited the diffusion process and leaded to a low specific capaci-

**Figure 3.** Crystal structures of α-, β-, γ-, δ- and λ-MnO2 [18].

tance value of only 110 μF cm−2. The result indicated that the limited performance was manly obtained on the surface of manganese oxide. Furthermore, the 2D δ-MnO2 with an interlayer separation of around 0.7 nm also obtained good capacitance value of 236 F g−1 which could be contributed to the sufficiently large layer space for a high rate insertion/ extraction of K+ cations.

Thus, the rational selection of crystal structures can effectively accelerate the charge storage process of metal oxide and improve its electrochemical properties.

#### **2.3. Specific surface area**

The pseuocapacitance of metal oxide depends on redox reactions which mainly take place on the surface area. Thus, the specific surface area is one of the most important factors for metal oxide-based supercapacitor applications. In general, the higher specific surface area can result in the higher the specific capacitance due to more active sites are capable of providing multiple redox reactions [21].

Obviously, to explore the higher specific surface area is an effective approach to achieve better capacitive performance of metal oxide. Up to now, amounts of attempts have been made on metal oxides including decreasing the size of their particles, optimizing their morphologies and combining them with carbon materials which have high specific surface areas.

#### **2.4. Morphology**

The morphology of metal oxide is a crucial factor which closely relates to the specific surface area, the diffusion pathway, surface to volume ratios and therefore the supercapacitive performance. Thus, considerable efforts have been focused on various metal oxides with different morphologies such as nanowires, nanorods, nanotubes, nanoflowers, hollow spheres, nanopillar array and porous thin films.

Several morphologies of metal oxide and their supercapacitive performance are discussed in this part. The first morphology is one-dimensional nanostructured metal oxides which generally enhance the specific capacitance through offering short diffusion path lengths for both ions and electrons as well as a large specific surface area. Gao et al. [22] successfully synthesized Co3O4 nanowires on nickel foam via template-free method shown in **Figure 4a**. The nanowires, with diameters around 250 nm and the lengths up to around 15 μm, displayed a maximum specific capacitance of 746 F g−1 at a current density of 5 mA cm−2. In addition, Lu et al. [23] reported a slim (<20 nm) NiO nanorod structure (**Figure 4b**) had an ultrahigh specific capacitance of 2018 F g−1 (80% of the theoretical value) at a current density of 2.27 A g−1 and high power density of 1536 F g−1 at 22.7 A g−1. Generally, the diameter plays an important role in one-dimensional nanostructure. The smaller diameter can offer larger specific surface area and more active sites leading to a better specific capacitance. It is also reported that the porous nanotube structure of MnO2 could not only enhance the specific capacitance, but also improve the stability of electrode due to accommodating large volume charges during the charge-discharge cycle [24]. The second metal oxide structure should be noted is hollow spheres. Various metal oxides with hollow sphere morphologies have been successfully synthesized recently to pursue a high loosely mesoporous structure, large specific surface area and fast ion/electron transfer. For example, Yan et al. [25] fabricated hierarchically porous NiO hollow spheres composed of nanoflakes shown in **Figure 4c** with thicknesses of ∼10 nm via a powerful chemical bath deposition method. The specific capacitance of NiO hollow spheres can remain 346 F g−1 at 1 A g−1 after 2000 cycles indicating an excellent supercapacitive performance. Another example is [26] that Co3O4 hollow spheres prepared by a facile carbonaceous microsphere templated synthesis as shown in **Figure 4d**. The as-obtained Co3O4 hollow spheres are composed of nanoparticles and possess a high surface area of 60 m2 g−1 owing to their mesoporous structure. Such a unique hollow-sphere architecture can greatly contributed to the comparatively high capacitance and excellent cycling stability. The third type of metal oxide morphology is three-dimensional porous structure. It can be seen in **Figure 4e** that 3D highly ordered nanoporous CuO with interconnected bimodal nanopores were fabricated by Moosavifard and coworkers [27]. This morphology offered a large specific surface area of 149 m2 g−1 displayed an excellent specific capacitance of 431 F g−1 at 3.5 mA cm−2 due to the 3D porous structure providing facilitated ion transport, short ion and electron diffusion pathways and more active sites for electrochemical reactions. Moreover, 3D-nanonet hollow structured Co3O4,

tance value of only 110 μF cm−2. The result indicated that the limited performance was manly obtained on the surface of manganese oxide. Furthermore, the 2D δ-MnO2 with an interlayer separation of around 0.7 nm also obtained good capacitance value of 236 F g−1 which could be contributed to the sufficiently large layer space for a high rate insertion/

Thus, the rational selection of crystal structures can effectively accelerate the charge storage

The pseuocapacitance of metal oxide depends on redox reactions which mainly take place on the surface area. Thus, the specific surface area is one of the most important factors for metal oxide-based supercapacitor applications. In general, the higher specific surface area can result in the higher the specific capacitance due to more active sites are capable of providing multiple

Obviously, to explore the higher specific surface area is an effective approach to achieve better capacitive performance of metal oxide. Up to now, amounts of attempts have been made on metal oxides including decreasing the size of their particles, optimizing their morphologies

The morphology of metal oxide is a crucial factor which closely relates to the specific surface area, the diffusion pathway, surface to volume ratios and therefore the supercapacitive performance. Thus, considerable efforts have been focused on various metal oxides with different morphologies such as nanowires, nanorods, nanotubes, nanoflowers, hollow spheres,

Several morphologies of metal oxide and their supercapacitive performance are discussed in this part. The first morphology is one-dimensional nanostructured metal oxides which generally enhance the specific capacitance through offering short diffusion path lengths for both ions and electrons as well as a large specific surface area. Gao et al. [22] successfully synthesized Co3O4 nanowires on nickel foam via template-free method shown in **Figure 4a**. The nanowires, with diameters around 250 nm and the lengths up to around 15 μm, displayed a maximum specific capacitance of 746 F g−1 at a current density of 5 mA cm−2. In addition, Lu et al. [23] reported a slim (<20 nm) NiO nanorod structure (**Figure 4b**) had an ultrahigh specific capacitance of 2018 F g−1 (80% of the theoretical value) at a current density of 2.27 A g−1 and high power density of 1536 F g−1 at 22.7 A g−1. Generally, the diameter plays an important role in one-dimensional nanostructure. The smaller diameter can offer larger specific surface area and more active sites leading to a better specific capacitance. It is also reported that the porous nanotube structure of MnO2 could not only enhance the specific capacitance, but also improve the stability of electrode due to accommodating large volume charges during the charge-discharge cycle [24]. The second metal oxide structure should be noted is hollow spheres. Various metal oxides with hollow sphere morphologies

and combining them with carbon materials which have high specific surface areas.

process of metal oxide and improve its electrochemical properties.

extraction of K+

**2.3. Specific surface area**

72 Supercapacitor Design and Applications

redox reactions [21].

**2.4. Morphology**

cations.

nanopillar array and porous thin films.

**Figure 4.** (a) Co3O4 nanowires [22]; (b) NiO nanorods [23]; (c) NiO hollow spheres [25]; (d) Co3O4 hollow spheres [26]; (e) 3D highly nanoporous CuO [27]; and (f) 3D nanonet hollow structured Co3O4 [28].

shown in **Figure 4f**, exhibited a maximum specific capacitance of 820 F g−1 at a scan rate of 5 mV s−1 and remained 90.2% of its initial capacitance after 1000 cycles [28]. The above results have proven that three-dimensional porous structured metal oxide is promising for supercapacitor applications.

As mentioned above, the electrochemical properties of metal oxide-based supercapacitors are largely affected by the morphologies. The rational design of morphologies with high specific surface area and porous nanostructures is necessary for the development of metal oxide-based electrodes.

#### **2.5. Conductivity**

It is well-known that the electronic conductivity of electrode materials is a vital factor to affect their high performance in supercapacitor applications. Unfortunately, the electronic conductivity of metal oxide materials is generally poor which largely limits ion and electron transfer. When the charge/discharge rate increases, the low conductivity would result in the localized charge/discharge process in a limited volume near the current collector, leading to low specific capacitance and low rate capability. For example, due to the low electronic conductivity (∼10−5 to 10−6 S cm−1), the realistic specific capacitance of manganese oxide is usually up to 350 F g−1 far behind its theoretic value of 1370 F g−1 [15, 29]. The same condition also takes place on other metal oxide materials such as NiO, Co3O4 and V2O5 [30, 31].

Thus, it is urgent to improve the conductivity of metal oxide in order to an ideal supercapacitive performance. One effective approach is doping metal elements into metal oxide and compositing metal oxide with high conductive carbon materials and conducting polymer.

#### **2.6. Mass loading**

The quantity of active materials loading on substrates can affect the specific capacitance, power and energy performance of metal oxide electrodes. On one hand, the large mass loading can cause longer transport paths for the diffusion of protons and an increase of thickness of metal oxide thin films resulting in lower electrical conductivity, limited access of electrolyte ions and higher series resistance. As a result, only partial active material on the surface of electrode film takes part in the charge storage leading to a lower specific capacitance of metal oxide. It is reported by Yang et al. [32] the specific capacitance of MnO2 thin films decreased from 203 to 155 F g−1 when the mass loading increased from 6 to 18 mg cm−2. On the other hand, a high mass loading is needed for high power and energy density, which thus makes the application of a light and durable supercapacitor possible [11]. Therefore, it is still a challenge for metal oxide to achieve both high mass loading and excellent specific capacitance.

## **3. Typical synthesis methods of metal oxide nanostructures**

Nanostructured metal oxide materials have been intensively investigated due to their superior supercapacitive performance. The factors mentioned above are all related to the metal oxide fabrication processes and parameters. The main synthesis techniques exploited include hydrothermal, electrodeposition, sol–gel, microwave assisted as well as template assisted methods.

#### **3.1. Hydrothermal method**

shown in **Figure 4f**, exhibited a maximum specific capacitance of 820 F g−1 at a scan rate of 5 mV s−1 and remained 90.2% of its initial capacitance after 1000 cycles [28]. The above results have proven that three-dimensional porous structured metal oxide is promising for

As mentioned above, the electrochemical properties of metal oxide-based supercapacitors are largely affected by the morphologies. The rational design of morphologies with high specific surface area and porous nanostructures is necessary for the development of metal oxide-based

It is well-known that the electronic conductivity of electrode materials is a vital factor to affect their high performance in supercapacitor applications. Unfortunately, the electronic conductivity of metal oxide materials is generally poor which largely limits ion and electron transfer. When the charge/discharge rate increases, the low conductivity would result in the localized charge/discharge process in a limited volume near the current collector, leading to low specific capacitance and low rate capability. For example, due to the low electronic conductivity (∼10−5 to 10−6 S cm−1), the realistic specific capacitance of manganese oxide is usually up to 350 F g−1 far behind its theoretic value of 1370 F g−1 [15, 29]. The same condition also takes place on

Thus, it is urgent to improve the conductivity of metal oxide in order to an ideal supercapacitive performance. One effective approach is doping metal elements into metal oxide and compositing metal oxide with high conductive carbon materials and conducting polymer.

The quantity of active materials loading on substrates can affect the specific capacitance, power and energy performance of metal oxide electrodes. On one hand, the large mass loading can cause longer transport paths for the diffusion of protons and an increase of thickness of metal oxide thin films resulting in lower electrical conductivity, limited access of electrolyte ions and higher series resistance. As a result, only partial active material on the surface of electrode film takes part in the charge storage leading to a lower specific capacitance of metal oxide. It is reported by Yang et al. [32] the specific capacitance of MnO2 thin films decreased from 203 to 155 F g−1 when the mass loading increased from 6 to 18 mg cm−2. On the other hand, a high mass loading is needed for high power and energy density, which thus makes the application of a light and durable supercapacitor possible [11]. Therefore, it is still a challenge for metal

Nanostructured metal oxide materials have been intensively investigated due to their superior supercapacitive performance. The factors mentioned above are all related to the metal oxide

oxide to achieve both high mass loading and excellent specific capacitance.

**3. Typical synthesis methods of metal oxide nanostructures**

other metal oxide materials such as NiO, Co3O4 and V2O5 [30, 31].

supercapacitor applications.

74 Supercapacitor Design and Applications

electrodes.

**2.5. Conductivity**

**2.6. Mass loading**

Hydrothermal synthesis is well-known as one of the most outstanding approaches to prepare nanoparticles due to a serious of advantages such as fine powder (nanoscale), high purity, good dispersion, uniform, narrow distribution, without agglomeration, good crystal form and shape controllability. In a hydrothermal process, crystal grows by chemical reactions taking place at high temperature and pressure conditions in a sealed pressure vessel with water as solvent. Under hydrothermal conditions, water can act as a chemical component and participate in the reactions. Moreover, the solvent is not only a mineralizing agent but also a pressure medium. By the control of physical and chemical factors, the formation and modification of nanostructured metal oxide can be achieved. Up to now, hydrothermal method has been successfully used to synthesize metal oxide with various nanostructures, such as nanowires, nanorods, nanoflowers, nanospheres, nanosheets, nanotubes and so on.

Purushothaman and his group [33] successfully prepared NiO nanoparticles via the hydrothermal method, using SDS as a surfactant. The different temperatures (120, 140, 160 and 180°C) in hydrothermal processes have been studied to optimize the morphology and electrochemical properties of NiO. The high degree of phase purity of the NiO particles with nanosizes of 8–16 nm have been achieved under all selected temperatures. However, the morphologies and electrochemical properties were different. At 120 and 140°C, the assembly of nanosheets is slow and, hence, they probably assemble into microspheres under the assistance of a surfactant. A decrease in the surface tension with increasing preparation temperature results in weak electrostatic interaction. The reduced surface tension lowers the aggregation, enabling the formation of microspheres with well-resolved nanosheets at 160°C. The initial nucleation and growth rate will be faster at 180°C. The faster nucleation hinders the assembly of anion surfactant and cation, resulting in the formation of nanorod assembled thicker plates. In case of supercapacitance, the sample prepared at 120°C exhibited a specific capacitance of 871 F g−1, while the value of sample formed at 140°C was 925 F g−1. The maximum specific capacitance of 989 F g−1 was obtained at 160°C while the specific capacitance showed a reduced value of 496 F g−1 at 180°C. The formation of a nanosheet-like structure seem to have facilitated the ion exchange process by reducing the diffusion lengths for the electrolyte, yielding a superior redox process in the sample prepared at 160°C, which exhibited the best specific capacitance. The specific capacitance was lower at elevated temperatures (180°C) might because the crystallites of larger size were formed, and they limited the paths available for ion transport. Moreover, Xia et al. [34] synthesized hollow Co3O4 nanowire arrays through a facile hydrothermal method. The Co3O4 nanowires have an average diameter of 200 nm, and the hollow centers have a diameter of 25 nm. In addition, a hierarchically porous can be found in nanowires allowing easier electrolyte penetration. Such a novel structure with porous walls and hollow center possesses more sites for ions to enter and allows facile ion diffusion at high current density leading to superior specific capacitances of 599 F g−1 at a current density of 2 A g−1 and 439 F g−1 at 40 A g−1.

#### **3.2. Electrodeposition methods**

Electrochemical deposition is a simple, fast, nonpolluting and facile technique, thus becomes one of most commonly used approaches to prepare metal oxide thin films. An electrochemical synthesis is achieved by a series of procedures that electron transfer between two or more electrodes separated by electrolyte making the occurrence of oxidation or reduction in the electrode–electrolyte interface which finally results in thin films deposited on electrode substrates. Electrochemical deposition can be divided into two different methods: anodic deposition and cathodic deposition. For example, Aghazadeh [35] prepared nanostructured Co3O4 via a simple cathodic electrodeposition method. The porous Co3O4 nanoplates displayed the average pore diameter and the surface area of 4.75 nm and 208.5 m2 g−1, respectively. A good specific capacitance as high as 393.6 F g−1 at the constant current density of 1 A g−1 and an excellent capacity retention (96.5% after 500 charge-discharge cycles) was obtained. Deng et al. [36] have reported that the nanoarchitectured CuO electrodes with a 3D hierarchically porous structure were prepared by an anodic electrodeposition method. An exceptionally large specific capacitance of 880 and 800 F g−1 was obtained at scan rates of 10 and 200 mV s−1.

According to the different modes of external power supplying, three main electrochemical deposition techniques including potentiostatic, galvanostatic and pulse period methods have been used by researchers. These different deposition routes with different applied current, potential and time have a crucial impact on the surface morphologies and crystal structures of metal oxide thin films. For example, Lee et al. [37] electrodeposited manganese oxide using three different modes: constant potential (CP) at 1 V for 900 s, pulse potential (PP) at 1 and 0 V with 0.5 s/0.5 s on-off time for 10,000 s, and pulse reverse potential (PRP) at 1 V and −1 V with 0.5 s/0.5 s interval time for 10,000 s. The different deposition times are applied to obtain similar mass loading. The results demonstrated that the different electrodeposition methods have a significant influence on the morphologies of manganese oxide. A traditional bulk film composed of 100 nm particles was prepared by CP mode and exhibited relatively low specific capacitance of 184 F g−1 at a scan rate of 10 mV s−1. In case of PP mode, nanostructured MnOx with porous flower petals morphology was obtained and showed a higher specific capacitance of 227.7 F g−1. The highest specific capacitance of 448 F g−1 was achieved by PRP mode due to the formation of nanorods with the average diameter of 20 nm which can supply higher specific surface area and faster ion transfer.

#### **3.3. Sol–gel method**

The sol–gel technique also attracts significant attention for the synthesis of nanostructured metal oxides because it offers controllable purity, composition, homogeneity of the products. Kim et al. [38] have reported that NiO nanostructures with three distinct morphologies were fabricated by a sol–gel method. The nanoflower structure was created in hexamethylene tetramine (HMTA) solution, while the nanoslice (diameters of 300–530 nm) was prepared in ammonium hydroxide (NH4OH) solution. The smaller nanoparticles with a diameter of around 50 nm were obtained when the reaction process took place in a strong basic LiOH. In addition, their morphology-dependent supercapacitor properties were exploited. Compared to the nanoslice and nanoparticle-shaped NiO, the nanoflowershaped NiO showed the best supercapacitor properties (480 F g−1 at 0.5 A g−1) despite it having the lowest specific surface area. This is because that the flower-shaped nanostructure has the unique three-dimensional (3D) networks which can provide longer diffusion paths and the highest pore volume (0.66 cm3 g−1) which offers advantages in contact with and transport of the electrolyte.

Yu and coworkers [39] successfully prepared the three-dimensional (3D) network mesoporous nanostructured α-MnO2 (MN-α-MnO2) powders using an inexpensive glucose–permanganate sol–gel method at room temperature and under ambient pressure. The MN-α-MnO2 exhibited high specific surface areas (ca. 220 m2 g−1) and narrow pore size distributions (5.6 nm) resulting in a good specific capacitance of 264 F g−1 after 1000 charge-discharge cycles.

#### **3.4. Microwave-assisted method**

current density leading to superior specific capacitances of 599 F g−1 at a current density of 2

Electrochemical deposition is a simple, fast, nonpolluting and facile technique, thus becomes one of most commonly used approaches to prepare metal oxide thin films. An electrochemical synthesis is achieved by a series of procedures that electron transfer between two or more electrodes separated by electrolyte making the occurrence of oxidation or reduction in the electrode–electrolyte interface which finally results in thin films deposited on electrode substrates. Electrochemical deposition can be divided into two different methods: anodic deposition and cathodic deposition. For example, Aghazadeh [35] prepared nanostructured Co3O4 via a simple cathodic electrodeposition method. The porous Co3O4 nanoplates displayed the average pore diameter and the surface area of 4.75 nm and 208.5 m2 g−1, respectively. A good specific capacitance as high as 393.6 F g−1 at the constant current density of 1 A g−1 and an excellent capacity retention (96.5% after 500 charge-discharge cycles) was obtained. Deng et al. [36] have reported that the nanoarchitectured CuO electrodes with a 3D hierarchically porous structure were prepared by an anodic electrodeposition method. An exceptionally large specific capacitance of 880 and 800 F g−1 was obtained at scan rates of 10 and 200 mV s−1.

According to the different modes of external power supplying, three main electrochemical deposition techniques including potentiostatic, galvanostatic and pulse period methods have been used by researchers. These different deposition routes with different applied current, potential and time have a crucial impact on the surface morphologies and crystal structures of metal oxide thin films. For example, Lee et al. [37] electrodeposited manganese oxide using three different modes: constant potential (CP) at 1 V for 900 s, pulse potential (PP) at 1 and 0 V with 0.5 s/0.5 s on-off time for 10,000 s, and pulse reverse potential (PRP) at 1 V and −1 V with 0.5 s/0.5 s interval time for 10,000 s. The different deposition times are applied to obtain similar mass loading. The results demonstrated that the different electrodeposition methods have a significant influence on the morphologies of manganese oxide. A traditional bulk film composed of 100 nm particles was prepared by CP mode and exhibited relatively low specific capacitance of 184 F g−1 at a scan rate of 10 mV s−1. In case of PP mode, nanostructured MnOx with porous flower petals morphology was obtained and showed a higher specific capacitance of 227.7 F g−1. The highest specific capacitance of 448 F g−1 was achieved by PRP mode due to the formation of nanorods with the average diameter of 20

The sol–gel technique also attracts significant attention for the synthesis of nanostructured metal oxides because it offers controllable purity, composition, homogeneity of the products. Kim et al. [38] have reported that NiO nanostructures with three distinct morphologies were fabricated by a sol–gel method. The nanoflower structure was created in hexamethylene tetramine (HMTA) solution, while the nanoslice (diameters of 300–530 nm) was prepared in ammonium hydroxide (NH4OH) solution. The smaller nanoparticles with

nm which can supply higher specific surface area and faster ion transfer.

A g−1 and 439 F g−1 at 40 A g−1.

76 Supercapacitor Design and Applications

**3.2. Electrodeposition methods**

**3.3. Sol–gel method**

Recently, microwave-assisted method has drawn large attention in the synthesis of oxide materials for supercapacitors application. Compared to conventional oil bath or hydrothermal heating, microwave heating can reduce the reaction time often by orders of magnitude, reduce the manufacturing cost and enhance product yield. An inverted temperature gradient takes place during the microwave-assisted process, and a rapid dielectric heating is generated internally within the material due to applied microwave radiation with a commonly used frequency of 0.3–2.45 GHz. Microwave-assisted synthesis has been adopted to prepare metal oxides with highly uniform nanostructures [40–42].

For example, Zhang et al. [43] have successfully synthesized γ-MnO2 nanoparticles and α-MnO2 urchin-like nanostructures by the microwave-assisted reflux as short as 5 min under neutral and acidic conditions, respectively. The γ-MnO2 nanoparticles showed a smaller particle size, a higher specific surface area and a larger pore volume than those of α-MnO2 urchin-like nanostructures resulting in a higher capacitance of 311 F g−1 at a current density of 0.2 A g−1. The specific capacitance retention and coulombic efficiency after 5000 cycles at 1 A g −1 were about 93% and almost 100% for γ-MnO2 nanoparticles, respectively.

Cao and coworkers [44] prepared flower-like NiO hollow nanosphere precursors via an efficient gas/liquid interfacial microwave-assisted process and were then transformed to NiO by simple calcinations. The wall of the sphere is composed of twisted NiO nanosheets that intercalated with each other. Such hollow structure is different from widely reported flowerlike nanostructures with solid cores. These flower-like NiO hollow nanospheres have high surface area of 176 m2 g−1. Electrochemical properties show a high specific capacitance of 585 F g−1 at a discharge current of 5 A g−1 and excellent cycling stability.

#### **3.5. Template-assisted method**

Hard templates are those materials which are either used as scaffolds for the deposition or employed not only as shape defined templates, but also as chemical reagents that react with other chemicals to produce desired nanomaterials. In the development of various metal oxide nanostructures, the hard template method is widely used and coupled with other methods, such as electrochemical deposition, solvothermal/hydrothermal and sol–gel methods. There are quite a lot of hard templates have been used for the synthesis of metal oxide nanostructures, such as porous anodic aluminum oxide (AAO), polycarbonate membranes (PC), carbon spheres, porous carbon, SiO2 spheres, mesoporous silica and naturally existing diatomite.

The anodic aluminum oxide (AAO) film is also one of the attractive templates since it possesses very regular and highly anisotropic porous structures with pore diameters ranging from below 10 to 200 nm, pore length from 1 to 50 mm and pore densities in the range of 109–1011 cm−2 [45]. The pores have been found to be uniform and nearly parallel, which is useful for the synthesis of one-dimensional metal oxide nanostructures, affording short ion diffusion paths and fast kinetics during the electrochemical reactions. Using AAO as the template, Dar et al. [46] synthesized NiO nanotubes via electrochemical deposition and nanorods after 25 min annealing at 450°C. Due to a suitable combination of nanocrystalline grain size and the high surface area akin to the tubular structure, NiO nanotube exhibits an excellent supercapacitive performance with a maximum specific capacitance of 2093 F g−1 which approaches the theoretical value of NiO (2584 F g−1). In contrast, the NiO nanorod structure is characterized by lower performance (797 F g−1). Furthermore, both NiO nanotube and nanorod show high stabilities with almost no alteration to performance after 500 cycles at high current densities of 125 and 80 A g−1. It has also been reported by Xu et al. [47] that Co3O4 nanotubes were successfully prepared via the AAO template method. The Co3O4 nanotubes have an average diameter of 300 nm and thickness of 50 nm which mainly controlled by the pore size of the AAO template. A good specific capacitance of 574 F g−1 was also obtained at a current density of 0.1 A g−1. However, the high cost of AAO templates limits their large-scale application for the production of well-organized metal oxide nanostructured electrodes.

In addition to the AAO template, carbon-based materials with different structures were also developed to prepare metal oxides. For instance, Du et al. [26] reported that Co3O4 hollow spheres composed of numerous small nanocrystals were prepared via one-pot hydrothermal carbonization and calcination method with carbon spheres as templates. The specific capacitance is 470 F g−1 at a current density of 1 A g−1, and no obvious capacitance decrease was observed over 1000 cycles of charge and discharge. Moreover, Yao et al. [48] synthesized nanostructured hierarchical mesoporous ribbon-like NiO via a hard-template method combining the calcination process. The mesoporous carbon was used as a hard template to control the structure growth and pore size distribution. A large surface area (147 m2 g−1) and high pore volume (0.2 cm3 g−1) were achieved when the molar ratio of Ni/C was 2/5. Notably, the outstanding pseudocapacitive performance was obtained with a high specific capacitance of 1260 F g−1 at a 1 A g−1 and only 5% deterioration of the initial capacitance after 5000 cycles.

## **4. Metal oxide-carbon composite electrodes**

other chemicals to produce desired nanomaterials. In the development of various metal oxide nanostructures, the hard template method is widely used and coupled with other methods, such as electrochemical deposition, solvothermal/hydrothermal and sol–gel methods. There are quite a lot of hard templates have been used for the synthesis of metal oxide nanostructures, such as porous anodic aluminum oxide (AAO), polycarbonate membranes (PC), carbon spheres, porous carbon, SiO2 spheres, mesoporous silica and naturally existing dia-

The anodic aluminum oxide (AAO) film is also one of the attractive templates since it possesses very regular and highly anisotropic porous structures with pore diameters ranging from below 10 to 200 nm, pore length from 1 to 50 mm and pore densities in the range of 109–1011 cm−2 [45]. The pores have been found to be uniform and nearly parallel, which is useful for the synthesis of one-dimensional metal oxide nanostructures, affording short ion diffusion paths and fast kinetics during the electrochemical reactions. Using AAO as the template, Dar et al. [46] synthesized NiO nanotubes via electrochemical deposition and nanorods after 25 min annealing at 450°C. Due to a suitable combination of nanocrystalline grain size and the high surface area akin to the tubular structure, NiO nanotube exhibits an excellent supercapacitive performance with a maximum specific capacitance of 2093 F g−1 which approaches the theoretical value of NiO (2584 F g−1). In contrast, the NiO nanorod structure is characterized by lower performance (797 F g−1). Furthermore, both NiO nanotube and nanorod show high stabilities with almost no alteration to performance after 500 cycles at high current densities of 125 and 80 A g−1. It has also been reported by Xu et al. [47] that Co3O4 nanotubes were successfully prepared via the AAO template method. The Co3O4 nanotubes have an average diameter of 300 nm and thickness of 50 nm which mainly controlled by the pore size of the AAO template. A good specific capacitance of 574 F g−1 was also obtained at a current density of 0.1 A g−1. However, the high cost of AAO templates limits their large-scale application for the production of well-organized metal oxide nano-

In addition to the AAO template, carbon-based materials with different structures were also developed to prepare metal oxides. For instance, Du et al. [26] reported that Co3O4 hollow spheres composed of numerous small nanocrystals were prepared via one-pot hydrothermal carbonization and calcination method with carbon spheres as templates. The specific capacitance is 470 F g−1 at a current density of 1 A g−1, and no obvious capacitance decrease was observed over 1000 cycles of charge and discharge. Moreover, Yao et al. [48] synthesized nanostructured hierarchical mesoporous ribbon-like NiO via a hard-template method combining the calcination process. The mesoporous carbon was used as a hard template to

control the structure growth and pore size distribution. A large surface area (147 m2

the outstanding pseudocapacitive performance was obtained with a high specific capacitance of 1260 F g−1 at a 1 A g−1 and only 5% deterioration of the initial capacitance after 5000

g−1) were achieved when the molar ratio of Ni/C was 2/5. Notably,

g−1) and

tomite.

78 Supercapacitor Design and Applications

structured electrodes.

high pore volume (0.2 cm3

cycles.

Preparing metal oxide-carbon composites is one of the most effective approaches to improve the supercapacitive performance of metal oxide electrodes. In such composite structures, the carbon materials with large specific surface area and high electric conductivity can provide the channels for charge transfer and benefit to the rate capability. Among a series of carbon materials, carbon nanotubes (CNTs), carbon nanofibers (CNFs), graphene and carbon nanofoams have been mostly studied to combine with metal oxide.

#### **4.1. Metal oxide-carbon nanotubes (CNTs) composite**

CNTs have outstanding pore structure, high electrical conductivity, and good mechanical and thermal stability which make them one of the most widely used carbon materials for supporting metal oxide. Thus, CNTs have been coupled with various metal oxides such as NiO, Co3O4, V2O5, MnO2, SnO2 and CuO to form the metal oxide–CNTs composite electrodes. It has been reported in 2005 that Lee et al. [49] synthesized NiO/CNTs nanocomposite via a hydrothermal method and explored the influences of CNT network existing in NiO. Compared to bare NiO, NiO/CNTs nanocomposite electrode exhibited a more rectangular shape in the CV curve and a smaller IR loss indicating a better supercapacitive performance. The specific capacitance increased from 122 to 160 F g−1 at a scan rate of 2 mV s−1 with the presence of 10% CNTs. The optimized properties owe to that CNTs can effectively improve the electrical conductivity of NiO and supply more active sites for redox reaction of NiO by increasing its specific surface area. After that, Lin et al. [50] prepared mesoporous sphere NiO nanostructures dispersing on the surface of CNTs and the maximum specific capacitance of 1329 F g−1 was observed at a very high current density of 84 A g−1. Gund et al. [51] fabricated highly flexible electrode with NiO/MWCNTs nanohybrid thin films on stainless steel substrate with an excellent specific capacitance of 1727 F g−1 at a current density of 5 mA cm−2 and good stability (91% retention after 2000 cycles).

The advantages of the metal oxide-CNTs composite were further demonstrated. Cheng et al. synthesized nanocomposites of V2O5 nanowires and interpenetrating CNTs via a hydrothermal process. When the nanocomposite contained 33 wt% of the CNTs, the V2O5-CNTs showed the best specific capacitance of 530 F g−1 which was significantly higher than the V2O5 nanowires (146 F g−1). The improved conductivity and the increased specific surface area (from 83 to 125 m2 g−1) were considered to be responsible for the better properties. Moreover, Wang et al. [52] designed a Co3O4@MWCNT nanocable using multiwall carbon nanotubes (MWCNTs) as the core cable. Compared to the pristine Co3O4 which has a low specific capacitance less than 130 F g−1, the prepared Co3O4@MWCNT nanocable exhibits a better performance with a specific capacitance of 590 F g−1 at 15 A g−1 and 510 F g−1 even at 100 A g−1. Furthermore, many efforts have been made on the preparation of MnO2-CNTs nanocomposites in order to improve the supercapacitive performance of MnO2. For example, MnO2-CNTs composites were prepared through a modified one-pot reaction process by Li and coworkers [53]. This cross-linked MnO2 nanoflakes-CNTs structure showed a good specific capacitance of 201 F g−1 and remarkable cycle stability (no obvious decay after 10,000 cycles). It has also been reported by Chen et al. [54] that MnO2 nanoparticles were introduced into the inner wall of CNT channels by a wet-chemistry method. The result of electrochemical tests shows that the composite has a much higher specific capacitance of 225 F g−1 than MnO2 with a value of 13 F g−1.

#### **4.2. Metal oxide-carbon nanofibers (CNFs) composite**

CNFs are very attractive as the support for metal oxides in the composite electrodes due to their conductive networks with appropriate pore channels. Typically the metal oxides are coated on the CNFs surface to form a core–shell structure in which the CNFs can serves as the physical backbone support and offer the channel for efficient electron and ion transportation. Zhi and coworkers [55] reported the synthesis of CNFs/MnO2 nanocomposite with a coaxial cable structure as shown in **Figure 5a**. The CNFs with a diameter of 200 nm coated by 4-nm-thick MnO2 nanowhiskers sheath giving a high specific surface area could be seen in **Figure 5b**. The nanocomposite electrode showed a good specific capacitance of 311 F g−1 for the whole electrode and 900 F g−1 for the MnO2 shell at a scan rate of 2 mV s−1. In addition, **Figure 5c** and **d** indicated that this CNFs/MnO2 nanocomposite also exhibited good cycling stability (2.4% loss after 1000 cycles), high energy density (80.2 Wh kg−1) and high power density (57.7 kW kg−1). Moreover, a Fe3O4/CNFs nanocomposite was designed by Fu et al. [56] through a solvent thermal reaction. Compared to the low specific capacitance (4 F g−1) of pure Fe3O4, the calculated specific capacitance of Fe3O4/ CNFs nanocomposite is as high as 127 F g−1 which indicates a better supercapacitive performance. The CNFs have not only improved the electronic/ionic conductivity of Fe3O4, but also prevented the aggregation of Fe3O4 nanosheets. CuO has also been used to fabricate composite with CNFs in order to improve its supercapacitive performance. As reported by Moosavifard [57], one-dimensional hierarchical hybrid CuO nanorod arrays-CNFs composite has been prepared via a solution method and an annealing treatment. The CuO nanorods with a length of around 300 nm and a diameter of around 15 nm are grown uniformly surrounding the CNFs. It should be noted that empty space existed among adjacent nanorods indicating a hierarchical array structure. The unique nanocomposite structure contributed to a high capacitance of 398 F g−1.

In addition, carbon fibers can also form a paper. Ghosh et al. [58] prepared carbon nanofiber paper (CFP) with fiber diameters ranging from 100 to 300 nm by electrospinning the polyacrylonitrile (PAN) precursor. The CFP showed good conductivity (0.1 S cm−1), high porosity and large surface area of 700 m2 g−1 indicating its potential to be a promising material for supporting metal oxide. Using this carbon nanofiber paper as substrate, 3-nm-thick V2O5 was obtained by electrodeposition method. The V2O5-CFP composite exhibited a total specific capacitance of 214 F g−1. Recently, Yang et al. [59] have reported a CFP/Co3O4 paper electrode with an excellent specific capacitance of 1124 F g−1 at a high current density of 25.34 A g−1 in the NaOH electrolyte. The composite also displayed a remarkable electrochemical stability with around 94.4% retention after 5000 charge-discharge cycles. The outstanding supercapacitive performance was attributed to the unique 1D nanonet structure of the electrodes and the improved electronic conductivity as well as ion diffusion by CFP.

#### **4.3. Metal oxide-graphene composite**

10,000 cycles). It has also been reported by Chen et al. [54] that MnO2 nanoparticles were introduced into the inner wall of CNT channels by a wet-chemistry method. The result of electrochemical tests shows that the composite has a much higher specific capacitance of

CNFs are very attractive as the support for metal oxides in the composite electrodes due to their conductive networks with appropriate pore channels. Typically the metal oxides are coated on the CNFs surface to form a core–shell structure in which the CNFs can serves as the physical backbone support and offer the channel for efficient electron and ion transportation. Zhi and coworkers [55] reported the synthesis of CNFs/MnO2 nanocomposite with a coaxial cable structure as shown in **Figure 5a**. The CNFs with a diameter of 200 nm coated by 4-nm-thick MnO2 nanowhiskers sheath giving a high specific surface area could be seen in **Figure 5b**. The nanocomposite electrode showed a good specific capacitance of 311 F g−1 for the whole electrode and 900 F g−1 for the MnO2 shell at a scan rate of 2 mV s−1. In addition, **Figure 5c** and **d** indicated that this CNFs/MnO2 nanocomposite also exhibited good cycling stability (2.4% loss after 1000 cycles), high energy density (80.2 Wh kg−1) and high power density (57.7 kW kg−1). Moreover, a Fe3O4/CNFs nanocomposite was designed by Fu et al. [56] through a solvent thermal reaction. Compared to the low specific capacitance (4 F g−1) of pure Fe3O4, the calculated specific capacitance of Fe3O4/ CNFs nanocomposite is as high as 127 F g−1 which indicates a better supercapacitive performance. The CNFs have not only improved the electronic/ionic conductivity of Fe3O4, but also prevented the aggregation of Fe3O4 nanosheets. CuO has also been used to fabricate composite with CNFs in order to improve its supercapacitive performance. As reported by Moosavifard [57], one-dimensional hierarchical hybrid CuO nanorod arrays-CNFs composite has been prepared via a solution method and an annealing treatment. The CuO nanorods with a length of around 300 nm and a diameter of around 15 nm are grown uniformly surrounding the CNFs. It should be noted that empty space existed among adjacent nanorods indicating a hierarchical array structure. The unique nanocomposite struc-

In addition, carbon fibers can also form a paper. Ghosh et al. [58] prepared carbon nanofiber paper (CFP) with fiber diameters ranging from 100 to 300 nm by electrospinning the polyacrylonitrile (PAN) precursor. The CFP showed good conductivity (0.1 S cm−1), high porosity and large surface area of 700 m2 g−1 indicating its potential to be a promising material for supporting metal oxide. Using this carbon nanofiber paper as substrate, 3-nm-thick V2O5 was obtained by electrodeposition method. The V2O5-CFP composite exhibited a total specific capacitance of 214 F g−1. Recently, Yang et al. [59] have reported a CFP/Co3O4 paper electrode with an excellent specific capacitance of 1124 F g−1 at a high current density of 25.34 A g−1 in the NaOH electrolyte. The composite also displayed a remarkable electrochemical stability with around 94.4% retention after 5000 charge-discharge cycles. The outstanding supercapacitive performance was attributed to the unique 1D nanonet structure of the electrodes and the

225 F g−1 than MnO2 with a value of 13 F g−1.

80 Supercapacitor Design and Applications

**4.2. Metal oxide-carbon nanofibers (CNFs) composite**

ture contributed to a high capacitance of 398 F g−1.

improved electronic conductivity as well as ion diffusion by CFP.

Since a mechanically exfoliated graphene monolayer was first observed and characterized in 2004, considerable research has been carried out in supercapacitor applications due to its large theoretical specific area (2630 m2 g−1), high electrical conductivity (104 S cm−2), abundant raw material resource and good electrochemical stability [60]. Therefore, graphene with these fascinating properties is expected as the potential supporting materials to improve the performance of metal oxide-based supercapacitors. Besides graphene, graphene oxide (GO) and reduced graphene oxide (rGO) have also attracted considerable attention due their unique physical and chemical properties. Until now, various metal oxides such as NiO, MnO2, CuO, V2O5, Co3O4 and TiO2 have been coupled with graphene materials to form supercapacitor electrodes for superior performance.

**Figure 5.** (a) TEM images of a single CNF@MnO2 nanostructure and (b) the MnO2 porous shell with nanowhiskers; (c) stability of the CNF@MnO2 electrodes; (d) Ragone plots of the CNF@MnO2 electrodes [55].

For example, Ge et al. [61] reported the preparation of 3D flower-like NiO and graphene sheets composite via incorporating a facile hydrothermal process with a thermal treatment process. The resultant composite exhibits a specific capacitance of 346 F g−1 (1.5 A g−1), a good rate performance and cycle stability in 2 M KOH. It should be noted that NiO in the composite could provide a specific capacitance as high as 778.7 F g−1, which far exceeded the bare NiO of only 220 F g−1. The magnitude of equivalent series resistances (ESR) are 0.71, 0.99 and 0.85 Ω cm−2 for graphene sheet, NiO and composite, respectively, indicating that the conductivity of the composite is improved by the presence of graphene sheets which could contribute to the superior supercapacitive performance of NiO. Wu et al. [62] successfully synthesized NiO particles/graphene oxide (GO) nanosheets composites. Compared with pure NiO and graphene, the composite electrode showed highest current response during CV scanning, reflecting the high capacitance value. The specific capacitance of the NiO/GO composite electrode (460 F g−1) is much higher than those of the bare graphene oxide electrode (13 F g−1) and NiO (40 F g−1) electrode at a current density of 10 A g−1. In addition, the capacitance retention of NiO/GO composite can remain nearly 100% after 3000 cycles which indicates excellent cycle-life stability. Moreover, Co3O4/graphene nanosheet (GNS) composite has been synthesized via a microwave-assisted method by Yan et al. [63]. The Co3O4 nanoparticles with a small size of 3–5 nm were uniformly distributed on the surface of graphene sheets. The electrochemical properties of composite were observed with a good specific capacitance of 243.2 F g−1 and an outstanding stability (only 4.3% loss after 2000 cycles).

Recently, the uniform rod-like V2O5 nanocrystals have been fabricated on the surface of reduced graphene oxide (rGO) to form the V2O5-rGO nanocomposites as the supercapacitor electrode [64]. The V2O5 nanoparticles on rGO were prepared by hydrolysing vanadium oxytripropoxide (VOTP) in ethanol solution with the existence of GO. For a comparative study, the pure V2O5 was also prepared in the same condition but without GO. The results have shown that the electrochemical performance of the V2O5-rGO nanocomposites with an excellent specific capacitance of 537 F g−1 at a current density of 1 A g−1 was much better than pure V2O5 which had a relatively low value of 202 F g−1. After 1000 charge-discharge cycles, the composite electrode could retain 84% of its initial capacitance while only 30% was retained for pure V2O5 indicating that V2O5-rGO nanocomposite electrode had a better electrochemical stability. In addition, the higher power and energy densities were also obtained in the composite. The synergistic effect of V2O5 nanorods and rGO has been considered to be responsible for the better supercapacitive performance. Firstly, the conductivity can be improved due to the presence of rGO. Secondly, the nanocomposites possess a larger surface area (49.16 m2 g−1) than that of pure V2O5 (37.57 m2 g−1). Thirdly, the rGO sheets can inhibit the disintegration of V2O5 and buffer the strain aroused by the volume expansion during the charging and discharging processes. Finally, the strong adhesion between V2O5 nanorods and rGO sheets may facilitate fast electron transfer through the highly conductive rGO sheets. Similarly, Xiang et al. [65] fabricated the rGO–TiO2 nanobelt composite by a hydrothermal processing in the ethanol solution. When the rGO/TiO2 mass ratio was 7:3, the composite obtained the best specific capacitance of 200 F g−1 which far exceeded pure TiO2 nanobelt (17 F g−1) and rGO (40 F g−1) in the Na2SO4 electrolyte at a scan rate of 2 mV s−1.

#### **4.4. Metal oxide-carbon foams composite**

In addition to high specific capacitance, high energy density and power density are also desirable for metal oxide-based supercapacitors. In general, increasing the mass loading can effectively store more energy and power. However, it is a challenge to load a large amount of materials on electrode without undermining the electrochemical performance. One promising approach is to form metal oxide-carbon nanofoams composite as electrodes. Three-dimensional (3D) carbon nanofoams with a through-connected pore network possess large specific surface areas allowing high metal oxide loading. Moreover, the high electric conductivity of carbon nanofoams can improve the electrochemical properties of metal oxide. Chen et al. [66] fabricated nanostructured MnO2 on CNT foams (or sponges) and the flower-like MnO2 nanoparticles were uniformly deposited on the skeleton of CNT sponges. An outstanding specific capacitance of 1270 F g−1 close to the theoretical value has been obtained and only 4% of degradation after 10,000 cycles at a charge-discharge current density of 5 A g−1. Furthermore, the specific power and energy of this composite are high with values of 63 kW kg−1 and 31 Wh kg−1, respectively. Dong and co-workers deposited Co3O4 nanowires on the 3D graphene foam in [67]. It can be seen in the figure, the graphene skeleton is fully and uniformly covered by the network of Co3O4 nanowires together provide a large accessible surface area. The composite electrode exhibited a high specific capacitance of ∼1100 F g−1 after 500 cycles at a current density of 10 A g−1 and stayed stable afterward indicating a good cycling stability.

## **5. Conclusion**

which could contribute to the superior supercapacitive performance of NiO. Wu et al. [62] successfully synthesized NiO particles/graphene oxide (GO) nanosheets composites. Compared with pure NiO and graphene, the composite electrode showed highest current response during CV scanning, reflecting the high capacitance value. The specific capacitance of the NiO/GO composite electrode (460 F g−1) is much higher than those of the bare graphene oxide electrode (13 F g−1) and NiO (40 F g−1) electrode at a current density of 10 A g−1. In addition, the capacitance retention of NiO/GO composite can remain nearly 100% after 3000 cycles which indicates excellent cycle-life stability. Moreover, Co3O4/graphene nanosheet (GNS) composite has been synthesized via a microwave-assisted method by Yan et al. [63]. The Co3O4 nanoparticles with a small size of 3–5 nm were uniformly distributed on the surface of graphene sheets. The electrochemical properties of composite were observed with a good specific capacitance of 243.2 F g−1 and an outstanding stability

Recently, the uniform rod-like V2O5 nanocrystals have been fabricated on the surface of reduced graphene oxide (rGO) to form the V2O5-rGO nanocomposites as the supercapacitor electrode [64]. The V2O5 nanoparticles on rGO were prepared by hydrolysing vanadium oxytripropoxide (VOTP) in ethanol solution with the existence of GO. For a comparative study, the pure V2O5 was also prepared in the same condition but without GO. The results have shown that the electrochemical performance of the V2O5-rGO nanocomposites with an excellent specific capacitance of 537 F g−1 at a current density of 1 A g−1 was much better than pure V2O5 which had a relatively low value of 202 F g−1. After 1000 charge-discharge cycles, the composite electrode could retain 84% of its initial capacitance while only 30% was retained for pure V2O5 indicating that V2O5-rGO nanocomposite electrode had a better electrochemical stability. In addition, the higher power and energy densities were also obtained in the composite. The synergistic effect of V2O5 nanorods and rGO has been considered to be responsible for the better supercapacitive performance. Firstly, the conductivity can be improved due to the presence of rGO. Secondly, the nanocomposites possess a larger surface area (49.16 m2 g−1) than that of pure V2O5 (37.57 m2 g−1). Thirdly, the rGO sheets can inhibit the disintegration of V2O5 and buffer the strain aroused by the volume expansion during the charging and discharging processes. Finally, the strong adhesion between V2O5 nanorods and rGO sheets may facilitate fast electron transfer through the highly conductive rGO sheets. Similarly, Xiang et al. [65] fabricated the rGO–TiO2 nanobelt composite by a hydrothermal processing in the ethanol solution. When the rGO/TiO2 mass ratio was 7:3, the composite obtained the best specific capacitance of 200 F g−1 which far exceeded pure TiO2 nanobelt (17 F g−1) and rGO (40 F

In addition to high specific capacitance, high energy density and power density are also desirable for metal oxide-based supercapacitors. In general, increasing the mass loading can effectively store more energy and power. However, it is a challenge to load a large amount of materials on electrode without undermining the electrochemical performance. One promising

(only 4.3% loss after 2000 cycles).

82 Supercapacitor Design and Applications

g−1) in the Na2SO4 electrolyte at a scan rate of 2 mV s−1.

**4.4. Metal oxide-carbon foams composite**

This chapter presents a relatively general understanding of the correlation between the composition, microstructure and electrochemical behaviors of metal oxide nanostructuresbased electrodes for the applications of supercapacitors. The current possibility of controlled growth and self-assembly represents an important step toward the design and tuning of metal oxide nanocrystals, and also it will be a significant step to the applications of metal oxides as ideal electrodes in high performance electrochemical energy storage devices.

## **Author details**

Zhenjun Qi\* , Shihao Huang, Adnan Younis, Dewei Chu and Sean Li

\*Address all correspondence to: zhenjun.qi@student.unsw.edu.au

School of Materials Science and Engineering, University of New South Wales, Sydney, NSW, Australia

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Fadhel El Kamel Fadhel El Kamel

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Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/65110

#### **Abstract**

In the electric double-layer capacitors (EDLCs), a large amount of electrical energy can be stored in the double layer by reversible accumulation of ions onto the active electrode material. In these devices, mobile charge carriers can accumulate (or deplete) near the electrode/electrolyte interface resulting in a space charge layer. So, the appropriate combination of space charge layer and large effective surface of the electrodes constitutes a significant factor to get high specific capacitance. Here, we incorporated protons in BaTiO3 films during a low-temperature deposition process. Drastic changes occurred on both chemical and electrical properties of the films when H2 was added to the sputtering gas. It is well known that protons are very mobile species even at low temperature. Therefore, upon the application of a sufficiently high electric field, positively charged protons move toward the cathode with an activation energy around 0.6 eV and pileup to form a capacitive double layer of several μF/cm2 which enhances the dielectric permittivity of the film.

**Keywords:** hydrogenated BaTiO3 films, double layer, all-solid-state supercapacitors

## **1. Introduction**

Owing to the progress in thin-film technology and materials engineering, electronic chips now integrate several functions on the same area, especially in wireless sensor networks, portable equipments, and other microsystems. As a result, microenergy sources need to be developed in order to drive such integrated electronic devices or to provide power during the temporary failure of the primary power sources. Two major performance criteria for any electrical energy storage device are required, energy and power densities. The first is defined as the amount of

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

energy stored per unit mass (Wh/kg) or per unit volume (Wh/m3 ) in the device. The latter is a measurement of how fast energy is extracted from, or transferred into, the device per unit mass (W/kg) or per unit volume (W/m3 ). The two criteria are particularly important where there is an excessive requirement for portability. Therefore, supercapacitors and Li-ion batteries should not necessarily be seen as competitors, because their charge storage mechanisms and thus their characteristics are different. Batteries convert chemical energy directly to electrical energy, in which charge is generated by redox reaction at electrodes and voltage is established between cellterminals depending on the chemical species and their concentrations. In previous batteries, this energy transformation is irreversible, but in novel ones, the chemical reaction is reversible, and thus the battery can be charged by supplying electrical energy to the cell.

**Figure 1.** Specific power vs. specific energy (*Ragone plot*) for various electrical energy storage devices. Time constant values of each device are obtained by dividing the energy density by the power density.

Generally, the capacitance obtained with conventional capacitors finds its origin in the electronic, ionic, and dipolar polarization occurring in the bulk, whereas the essential reasons for electrical energy storage in supercapacitors are achieved by both the heavy load of electrode material per unit area and the relatively large specific capacitance of the electrolyte material. Supercapacitors can principally be classified as either electric double-layer capaci-

tors (EDLCs) or pseudocapacitors. The energy storage mechanism in the former devices relies on the separation of charges at the electrode/electrolyte interface, whereas in the latter, a faradic process occurs in addition to a simple charge separation. This feature explains why the charge storage capacity of pseudocapacitors is typically larger than that of EDLCs.

energy stored per unit mass (Wh/kg) or per unit volume (Wh/m3

mass (W/kg) or per unit volume (W/m3

90 Supercapacitor Design and Applications

the cell.

measurement of how fast energy is extracted from, or transferred into, the device per unit

there is an excessive requirement for portability. Therefore, supercapacitors and Li-ion batteries should not necessarily be seen as competitors, because their charge storage mechanisms and thus their characteristics are different. Batteries convert chemical energy directly to electrical energy, in which charge is generated by redox reaction at electrodes and voltage is established between cellterminals depending on the chemical species and their concentrations. In previous batteries, this energy transformation is irreversible, but in novel ones, the chemical reaction is reversible, and thus the battery can be charged by supplying electrical energy to

**Figure 1.** Specific power vs. specific energy (*Ragone plot*) for various electrical energy storage devices. Time constant

Generally, the capacitance obtained with conventional capacitors finds its origin in the electronic, ionic, and dipolar polarization occurring in the bulk, whereas the essential reasons for electrical energy storage in supercapacitors are achieved by both the heavy load of electrode material per unit area and the relatively large specific capacitance of the electrolyte material. Supercapacitors can principally be classified as either electric double-layer capaci-

values of each device are obtained by dividing the energy density by the power density.

) in the device. The latter is a

). The two criteria are particularly important where

The availability of the stored charge will always be faster for supercapacitor (*surface storage*) than for a Li-ion battery (*bulk storage*), with a larger stored energy for the latter [1]. Both devices must be used in their respective time-constant domains (**Figure 1**).

Using a Li-ion battery for repeated high-power delivery/uptake applications for a short duration (less than 10 s) will quickly degrade the cycle life of the system. The only way to avoid this is to oversize the battery, increasing the cost and volume. In the same way, using supercapacitors for power delivery longer than 10 s requires oversizing.

Supercapacitors [2, 3], until now consisting of liquid-state electrolytes, have been widely regarded as energy storage devices for several electronic systems. Hence, they can ensure this power request since they have high-power density, which can be supplied in a very short time. Pseudocapacitive metal oxides (RuO2) [4, 5], conducting polymers [6], and carbon-based materials such as carbon nanotubes [7], graphite oxides [8], onion-like carbon [9], and activated carbon [10] have been widely reported in literature as electrode materials especially for liquidstate supercapacitors with interdigital fingers [4, 9, 11] and roll-like [12] and sandwich [8] shapes. To date, these supercapacitors are principally based on liquid-state electrolytes, such as aqueous or organic solutions, and display high capacitance values (from 1 to 100 mF/cm2 ), according to the electrode nature and its geometry, the number of interdigitated microelectrodes in the device, the electrolyte, and so on. However, these devices cannot be used at high temperatures, because the aqueous or organic electrolytes undergo decomposition. In addition, they exhibit leakage current and ionic conductivity which vary significantly with temperature and frequency.

Although solid electrolytes are characterized by lower ionic conductivity compared to their liquid counterparts, we can overcome this drawback by decreasing the thickness of solid electrolytes [13] in order to reduce the diffusion path of charged ionic species or by increasing the effective electrode surface by using porous materials [14]. On the other hand, with solidstate electrolyte, supercapacitors have wide operational temperature range and negligible leakage current. Hydrogel-polymer electrolyte has been reported by Kaempgen et al. [15] for the all-solid-state supercapacitors. The estimated cell capacitance was around 1 mF/cm2 , but the operating voltage range is limited to 1 V, making them nonfunctional for most applications. Yoon et al. [5] have investigated the use of a pseudocapacitive amorphous RuO2 electrode with a LixPOyNz (LiPON) solid electrolyte for the elaboration of a thin-film supercapacitor. They reported a specific capacitance of about 3.5 mF/cm2 ×μm and a high operating voltage range. However, the fast capacitance degradation observed after several charge/discharge cycles derives from the lower ion mobility of Li+ ions in the LiPON electrolyte than that of H+ and OH− ions in the liquid electrolyte. This suggests that solid electrolyte with high protonic conductivity holds promise for thin-film supercapacitor with high capacitance and high cycle life. Recently, solid-state electrolytes such as hydrated lithium fluoride [2], yttria-stabilized zirconia [3], and Ta2O5:H [16] are examples of ionic conductors investigated for such a purpose. The deposition of solid-state supercapacitors reported up to now involves high-temperature processing [5, 17] or humidified environment [2]. This constitutes a serious limitation to integrate supercapacitors in electronic chips. As an alternative, here protons were incorporated in barium titanate (BaTiO3) films during a low-temperature deposition process. The main purpose of this study was focused on related electrical defects. We emphasize that drastic changes occur on both chemical and electrical properties of the films when H2 is added to the sputtering gas. The electric double-layer capacitance can reach values up to several μF/cm2 .

## **2. Experimental details**

Barium titanate films were grown by rf magnetron sputtering process on gold, copper, and carbon nanowalls/Pt-coated silicon substrates (*previously coated with a chromium adhesion layer*). Film thickness is around 1 μm as measured by a Veeco Dektak 6 M surface profiler and confirmed by a scanning electron microscopy. During deposition the substrates were water cooled (*room temperature deposition*). Standard sputtering was performed with pure argon. Here, hydrogen was incorporated in the BaTiO3 layers during their growth by introducing H2 in the sputtering gas [H2/(H2 + Ar) varies from 0 to 30 %]. The stoichiometry of the sputtering gas (H2 + Ar) was controlled using respective mass flows for each gas (Ar *and* H2). Since the films are grown at low temperature, they are found to be amorphous, as detected by X-ray diffraction.

Electrical and dielectric measurements were performed, in a dark-shielded cell filled with dry nitrogen, on M/*a*-BaTiO3:H/M (M = Au, Cu) planar capacitors, where gold and copper electrodes (1.77 mm2 *area*) were evaporated through a shadow mask on the front side of the deposited films. Dielectric properties of these devices were studied as a function of temperature and frequency using a Novocontrol impedance analyzer (*with an ac test voltage of* 100 mV).

The capacitance (*C*), conductance (*G*), and dissipation factor (tan*δ* = *G*/*Cω*) were measured from 10−1 to 106 Hz. Known the geometrical factor (*top electrode diameter and film thickness*), the dielectric constant *εʹ* and conductivity *σ* can be determined, and results will be presented as *εʹ*(*f*), *tanδ*(*f*), and *σ*(*f*) plots. Current-voltage measurements were performed using a Keithley 6517A electrometer. Temperature variation was controlled by a Linkam hot stage.

## **3. Results and discussion**

Films grown under pure argon are transparent [18]. This transparency diminishes widely with the addition of hydrogen in the sputtering gas. The color of hydrogenated films tends gradually to yellow and then to brown, and finally they darken for hydrogen mixing ratio (HMR) around 25 %. This dark color indicated a high oxygen deficiency [19, 20]. In order to better understand this color change, optical absorbance spectrum was recorded on the hydrogenated barium titanate films at room temperature using a Shimadzu (UV-3101 PC) UV-Vis-NIR scanning spectrophotometer in diffuse reflection mode. Amorphous barium titanate is reported as a direct band gap material. So, it is evident that in the shorter wavelength region, the absorption coefficient α follows a power law:

The deposition of solid-state supercapacitors reported up to now involves high-temperature processing [5, 17] or humidified environment [2]. This constitutes a serious limitation to integrate supercapacitors in electronic chips. As an alternative, here protons were incorporated in barium titanate (BaTiO3) films during a low-temperature deposition process. The main purpose of this study was focused on related electrical defects. We emphasize that drastic changes occur on both chemical and electrical properties of the films when H2 is added to the sputtering gas. The electric double-layer capacitance can reach values up to several μF/cm2

Barium titanate films were grown by rf magnetron sputtering process on gold, copper, and carbon nanowalls/Pt-coated silicon substrates (*previously coated with a chromium adhesion layer*). Film thickness is around 1 μm as measured by a Veeco Dektak 6 M surface profiler and confirmed by a scanning electron microscopy. During deposition the substrates were water cooled (*room temperature deposition*). Standard sputtering was performed with pure argon. Here, hydrogen was incorporated in the BaTiO3 layers during their growth by introducing H2 in the sputtering gas [H2/(H2 + Ar) varies from 0 to 30 %]. The stoichiometry of the sputtering gas (H2 + Ar) was controlled using respective mass flows for each gas (Ar *and* H2). Since the films are grown at low temperature, they are found to be amorphous, as detected by X-ray diffrac-

Electrical and dielectric measurements were performed, in a dark-shielded cell filled with dry nitrogen, on M/*a*-BaTiO3:H/M (M = Au, Cu) planar capacitors, where gold and copper electrodes (1.77 mm2 *area*) were evaporated through a shadow mask on the front side of the deposited films. Dielectric properties of these devices were studied as a function of temperature and frequency using a Novocontrol impedance analyzer (*with an ac test voltage of* 100 mV). The capacitance (*C*), conductance (*G*), and dissipation factor (tan*δ* = *G*/*Cω*) were measured from

dielectric constant *εʹ* and conductivity *σ* can be determined, and results will be presented as *εʹ*(*f*), *tanδ*(*f*), and *σ*(*f*) plots. Current-voltage measurements were performed using a Keithley

Films grown under pure argon are transparent [18]. This transparency diminishes widely with the addition of hydrogen in the sputtering gas. The color of hydrogenated films tends gradually to yellow and then to brown, and finally they darken for hydrogen mixing ratio (HMR) around 25 %. This dark color indicated a high oxygen deficiency [19, 20]. In order to better understand this color change, optical absorbance spectrum was recorded on the hydrogenated barium titanate films at room temperature using a Shimadzu (UV-3101 PC) UV-Vis-NIR scanning spectrophotometer in diffuse reflection mode. Amorphous barium

6517A electrometer. Temperature variation was controlled by a Linkam hot stage.

Hz. Known the geometrical factor (*top electrode diameter and film thickness*), the

**2. Experimental details**

92 Supercapacitor Design and Applications

tion.

10−1 to 106

**3. Results and discussion**

.

$$
\alpha \hbar \nu = \mathsf{B} \sqrt{\hbar \nu - \mathsf{E}\_0} \tag{1}
$$

where *hν* is the energy of the incident photon, *B* is the absorption edge width parameter, and *E*0 is the optical band gap. Tauc plots [(*αhν*)2 vs. *hν*] for BaTiO3:H films are shown in **Figure 2**.

**Figure 2.** (a) UV-Vis absorbance spectra of hydrogenated barium titanate films deposited under different HMRs. Optical band gap energy was extracted by linear fit calculation. (b) Optical band gap energy as a function of HMR in the sputtering gas.

The linear fit of the straight-line portion of the data indicates a direct band gap with an energy value varying from 3.76 to 3.40 eV when HMR varies from 0 to 30 %. The optical absorption in the UV region is mainly attributed to the electron transition from the valence band maximum to the conduction band minimum. So, a plausible explication for the band gap narrowing can be the existence of several defect levels, oxygen vacancies, or other kinds of punctual defects, within the band gap.

Evidences for hydrogen incorporation were previously given by the Fourier transform infrared spectroscopy and X-ray photoelectron spectroscopy [21]. It is worth noting that the hydrogenated films display a large density of hydroxide compared to the standard films (*grown under pure argon*), while the density of oxide remains constant in both films. The incorporation of hydrogen in the films during the deposition process can explain this feature. As it is already established [22–24], the hydrogen ionizes to give an electron (*donor defect*) and a proton (H+ ) which combines with oxygen of the perovskite lattice to form hydroxide groups.

**Figure 3** shows an overview on the frequency spectra associated to the capacitance density *C*(*f*) (**Figure 3(a)**) and the conductivity *σ*(f) (**Figure 3(b)**) performed at room temperature on *a*-BaTiO3:H films grown under different hydrogen mixing ratios (HMRs) in the sputtering gas. We note that as hydrogen was introduced with a content exceeding 10 %, films display (**Figure 3(a)**) hundred times higher permittivity (~2500) than that measured on films grown under pure argon (~20). The double layer is clearly evidenced, especially at low temperatures by a dispersive behavior of the capacitance accompanied by a relaxation peak (LT Relax) in the loss measurements (**Figure 4(a)**).

**Figure 3.** Frequency dependence of the capacitance density (a) and the conductivity (b) measured at room temperature on barium titanate layers sputtered under different hydrogen mixing ratios ranging from 0 to 30 %. The inset shows the variation of real permittivity vs. the hydrogen mixing ratios.

**Figure 4.** Temperature dependence of the loss factor tan*δ*(*f*), measured on films grown under 25 % HMR. Data curve shows two relaxation peaks, which appear at low (LT Relax). (a)) and high (HT Relax. (b)) temperatures.

We believe that the dielectric losses are mainly the result of the ion migration and/or trap release. At high frequencies (*f* > 10 kHz), *C* relaxes toward the bulk contribution, where mobile species cannot reach the electrodes under the ac test voltage (electrode polarization is basically a slow relaxation mechanism). As we can see in **Figure 3(a)**, the bulk capacitance of the hydrogenated films tends to the value measured when films were grown under pure argon.

The cutoff frequency (*f* = 10 kHz) represents the speed of the capacitor [2] which constitutes an important parameter for both fundamental studies and technological applications. Furthermore, conductivity is also affected by the addition of hydrogen in the sputtering gas. At low frequency, it increases by up to eight orders of magnitude. The *σ*(*f*) characteristic shifts toward higher frequencies with the appearance of two plateaus, corresponding to the bulk and interfacial conductivities, that appear at high and low frequencies, respectively. This feature was previously [25] explained through the electrode polarization mechanism [26] which arises from proton accumulation at the cathode over a Debye length and gives rise to a large capacitance [25]. In the low-frequency domain, measured capacitance was completely determined by the charge stored in the double layer. As a result, electric double-layer capacitors have been widely regarded as energy storage devices.

We note that as hydrogen was introduced with a content exceeding 10 %, films display (**Figure 3(a)**) hundred times higher permittivity (~2500) than that measured on films grown under pure argon (~20). The double layer is clearly evidenced, especially at low temperatures by a dispersive behavior of the capacitance accompanied by a relaxation peak (LT Relax) in

**Figure 3.** Frequency dependence of the capacitance density (a) and the conductivity (b) measured at room temperature on barium titanate layers sputtered under different hydrogen mixing ratios ranging from 0 to 30 %. The inset shows

**Figure 4.** Temperature dependence of the loss factor tan*δ*(*f*), measured on films grown under 25 % HMR. Data curve

We believe that the dielectric losses are mainly the result of the ion migration and/or trap release. At high frequencies (*f* > 10 kHz), *C* relaxes toward the bulk contribution, where mobile species cannot reach the electrodes under the ac test voltage (electrode polarization is basically a slow relaxation mechanism). As we can see in **Figure 3(a)**, the bulk capacitance of the hydrogenated films tends to the value measured when films were grown under pure argon.

shows two relaxation peaks, which appear at low (LT Relax). (a)) and high (HT Relax. (b)) temperatures.

the loss measurements (**Figure 4(a)**).

94 Supercapacitor Design and Applications

the variation of real permittivity vs. the hydrogen mixing ratios.

**Figure 5.** Selected literature data for the *C*(*f*) characteristics measured on different capacitors. (a) BaTiO3 crystal (2 mm *thick*) sandwiched between two smooth metal electrodes. Protons are incorporated as mobile carriers by hydrogen charge for 8 days using electrolysis of water. (b) Yttria-stabilized zirconia (0.6 mm *thick*) sandwiched between two porous Pt/YSZ composite layers. Electric double layer is formed by the accumulation of oxygen ions. (c–d) Cu/ LiF(0.32 μm *thick*)/Cu-based supercapacitor. Under the electric field, two charge carriers (Li+ and F− ) are accumulated at the interfaces. Measurements were carried out under two humid environments, 34 % RH (c) and 80 % RH (d). (e) *a*-BaTiO3:H (1 μm *thick*) sandwiched between two smooth Au electrodes (*the present work*). Electric double layer is formed by the accumulation of mobile proton carriers. (f) Solid-state supercapacitor from single-walled carbon nanotube arrays coated with Al2O3.

In general, electric double-layer capacitors (EDLCs) exhibit the property that a large electrical energy can be electrostatically stored in the double layer by reversible accumulation of ions (*from the electrolyte*) onto the active electrode material that is electrochemically stable. In these devices mobile charge carriers can accumulate or deplete near the electrode/electrolyte interface resulting in a space charge layer. So, the appropriate combination of space charge layer and large effective surface of the electrodes constitutes a significant factor to get high specific capacitance. **Figure 5** shows selected literature data for the *C*(*f*) characteristics measured on different capacitors.

The most common devices at present use carbon-based active materials (*or other porous composite materials*) with high surface area as electrodes in EDLCs. In the present study, we used smooth metallic electrodes (*low surface area*) in order to study the effect of electrolyte without any contribution of the electrode nature in addition to carbon nanowall-coated platinum (CNW/Pt) as a high specific surface electrode. **Figure 6** displays the *C*(*f*), *ε*(*f*), and *σ*(*f*) characteristics of Au/*a*-BaTiO3/Au (a), Au/*a*-BaTiO3:H/Au (c), and Au/*a*-BaTiO3:H/Cu (d) devices.

**Figure 6.** *C*(*f*), *ε*(*f*), and *σ*(*f*) characteristics of (a) Au/*a*-BaTiO3/Au, (b) Au/*a*-BaTiO3/*a*-BaTiO3:H(20%H2)/*a*-BaTiO3/Au, (c) Au/*a*-BaTiO3:H(25%H2)/Au, (d) Au/*a*-BaTiO3:H(25%H2)/Cu, and (e) Au/*a*-BaTiO3:H(25%H2)/CNW/Pt devices. Even with low effective surface of electrodes, we can get acceptable values of capacitance that is higher than 2 μF/cm2 , which leads to a real permittivity in the 103 –104 range.

We showed that even with low effective surface of electrodes, we can get acceptable values of capacitance that is higher than 2 μF/cm2 , which leads to a real permittivity in the 103 –104 range. By increasing the surface area of the electrodes, we anticipate higher specific capacitance for our devices as observed with the Au/*a*-BaTiO3:H/CNW/Pt (e) device (*C* = 10 μF/cm2 *and εʹ* = 104 *at low frequency*).

To study the temperature dependence of the conduction mechanisms taking place in hydrogenated films, we plotted the relaxation frequencies *f*0(*T*), bulk *σ*b(*T*), and interfacial *σ*<sup>i</sup> (*T*) conductivities, extracted from the relaxation peak position (**Figure 4**) and high- and lowfrequency semicircles in the impedance diagram (**Figure 7**), respectively, as a function of 1000/ *T* (*Arrhenius plot*, **Figure 8**).

specific capacitance. **Figure 5** shows selected literature data for the *C*(*f*) characteristics meas-

The most common devices at present use carbon-based active materials (*or other porous composite materials*) with high surface area as electrodes in EDLCs. In the present study, we used smooth metallic electrodes (*low surface area*) in order to study the effect of electrolyte without any contribution of the electrode nature in addition to carbon nanowall-coated platinum (CNW/Pt) as a high specific surface electrode. **Figure 6** displays the *C*(*f*), *ε*(*f*), and *σ*(*f*) characteristics of Au/*a*-BaTiO3/Au (a), Au/*a*-BaTiO3:H/Au (c), and Au/*a*-BaTiO3:H/Cu (d)

**Figure 6.** *C*(*f*), *ε*(*f*), and *σ*(*f*) characteristics of (a) Au/*a*-BaTiO3/Au, (b) Au/*a*-BaTiO3/*a*-BaTiO3:H(20%H2)/*a*-BaTiO3/Au, (c) Au/*a*-BaTiO3:H(25%H2)/Au, (d) Au/*a*-BaTiO3:H(25%H2)/Cu, and (e) Au/*a*-BaTiO3:H(25%H2)/CNW/Pt devices. Even with low effective surface of electrodes, we can get acceptable values of capacitance that is higher than 2 μF/cm2

We showed that even with low effective surface of electrodes, we can get acceptable values of

By increasing the surface area of the electrodes, we anticipate higher specific capacitance for our devices as observed with the Au/*a*-BaTiO3:H/CNW/Pt (e) device (*C* = 10 μF/cm2 *and εʹ* =

To study the temperature dependence of the conduction mechanisms taking place in hydrogenated films, we plotted the relaxation frequencies *f*0(*T*), bulk *σ*b(*T*), and interfacial *σ*<sup>i</sup>

, which leads to a real permittivity in the 103

, which

(*T*)

–104 range.

ured on different capacitors.

96 Supercapacitor Design and Applications

leads to a real permittivity in the 103

104 *at low frequency*).

capacitance that is higher than 2 μF/cm2

–104 range.

devices.

**Figure 7.** Complex impedance spectra (*Z*ʺ–*Zʹ*) carried out, at different temperatures, on hydrogenated barium titanate films. Samples were grown under 25 % HMR in the sputtering gas.

**Figure 8.** Relaxation frequency, bulk, and interfacial conductivities as a function of 1000/*T* (*Arrhenius diagram*).

Measurements were carried out on samples grown under 25 % HMR in the sputtering gas. Both characteristics show the Arrhenius-type dependence over the whole temperature range, excluding the low-temperature (*T* < −100 °C) interfacial conductivity. As shown in **Figure 8**, we can extract three activation energies which correspond to three different conduction processes.

I–V characteristics were recorded on hydrogenated films at different temperatures (from −150 to −50 °C) in order to explain the conduction mechanism in dc regime. DC bias was applied to the bottom electrode, and the current was measured after 60 s for stabilization concern. Experimental data measured on the Au/*a*-BaTiO3:H (17% HMR)/Au device are shown in **Figure 9(a)**.

**Figure 9.** (a) Temperature dependence of the leakage current (*J*–*E characteristics*). (b) *J*–*E* characteristics replotted according to the Poole-Frenkel model.

It is clearly seen that the leakage current density exhibits a Poole-Frenkel (PF)-type behavior, which implies that the conduction mechanism can be described by a thermally stimulated emission from a discrete set of traps. The PF model predicts that the current density can be expressed as:

$$J\_{\rm PF} = qN\_{\rm c} \mu E \exp\left(\frac{-q\left(\phi\_{\rm t} - \sqrt{qE \wedge \pi \varepsilon\_{\rm opt}}\right)}{\mathbf{k}\_{\rm B}T}\right) \tag{2}$$

where *E* is the applied field, kB is the Boltzmann constant, *μ* is the electronic mobility, *ε*opt is the optical permittivity, *q* is the elementary charge, *q*φt is the trap level, and *N*c is the effective density of states in the conduction band assumed to be around 1018 cm−3 [27]. As stated above (Poole-Frenkel *model*), log(*J*/*E*) vs. *E*1/2 should be linear (**Figure 9(b)**). Then, we guess that the leakage current arises from the carriers release from shallow trap levels localized within the band gap. In order to determine this energy level, we plotted the current density vs. 1/*T* at a fixed electric field. *J* vs. 1/*T* curves (*not shown here*) display a linear behavior with a negative slope which suggests that the conduction process is thermally stimulated. Activation energy and carrier mobility vary, respectively, from 0.12 and 4.30 × 10−7 to 0.05 eV and 2.00 × 10−6 cm2 / Vs when the electric field varies from 0.2 to 3.0 MV/m. Such energy values correspond to an electric-field–dependent effective trap depth [ϕ0 − Δϕ(*E*)]. Based on the slowness of the PF mechanism, these mobility values are in a good agreement with the electronic conduction. In the literature, electronic mobility shows a large discrepancy and is found to vary from 10−3 [28– 32] to 10−10 cm2 /Vs [27]. To determine the real trap's depth ϕ0 (*at E* = 0 V/m), we plotted the effective trap depths [ϕ0 − Δϕ(*E*)] as a function of *E*1/2 (**Figure 10**).

Measurements were carried out on samples grown under 25 % HMR in the sputtering gas. Both characteristics show the Arrhenius-type dependence over the whole temperature range, excluding the low-temperature (*T* < −100 °C) interfacial conductivity. As shown in **Figure 8**, we can extract three activation energies which correspond to three different conduction

I–V characteristics were recorded on hydrogenated films at different temperatures (from −150 to −50 °C) in order to explain the conduction mechanism in dc regime. DC bias was applied to the bottom electrode, and the current was measured after 60 s for stabilization concern. Experimental data measured on the Au/*a*-BaTiO3:H (17% HMR)/Au device are shown in

**Figure 9.** (a) Temperature dependence of the leakage current (*J*–*E characteristics*). (b) *J*–*E* characteristics replotted ac-

It is clearly seen that the leakage current density exhibits a Poole-Frenkel (PF)-type behavior, which implies that the conduction mechanism can be described by a thermally stimulated emission from a discrete set of traps. The PF model predicts that the current density can be

where *E* is the applied field, kB is the Boltzmann constant, *μ* is the electronic mobility, *ε*opt is the

density of states in the conduction band assumed to be around 1018 cm−3 [27]. As stated above

(2)

is the trap level, and *N*c is the effective

processes.

98 Supercapacitor Design and Applications

**Figure 9(a)**.

cording to the Poole-Frenkel model.

optical permittivity, *q* is the elementary charge, *q*φt

expressed as:

We found that the leakage current can be explained by the carrier release from shallow trap level ϕ0 localized at around 0.15 eV below the conduction band. It is believed that positively charged protons can provide such energetic levels. It is interesting to note that the same amount of energy is required to activate the LT relaxation, bulk, and interfacial (*at T* < 120 °C) conduction processes (**Figure 8**).

**Figure 10.** Effective trap depth as a function of *E*1/2. Leakage currents can be explained by carrier detrapping from shallow trap level ϕ0 localized at around 0.15 eV below the conduction band.

In amorphous materials, charged defects surrounding the proton tilt markedly the hydroxide group toward one of the neighboring oxygen ions. Oxygen vacancies constitute the main charged defects since the hydrogenated films were deposited under a reduced atmosphere. They effectively act as positively charged defects that repel the proton. This feature allows the formation of a rather weak directional interaction (*hydrogen bond*) between the proton and one of the adjacent oxygen, O[1]-H~O[2] [33]. Such a site could be considered as a trap for protons. In this approach, electron issued from the ionization of hydrogen is weakly bonded, and it can be easily activated to the conduction band, from a shallow donor level estimated at around 0.15 eV, to increase the dielectric loss and the leakage current [34].

The hydrogen bond breakage can be responsible for a large proton conduction process in oxide materials. Several reports (*Ref*. [35] *and references therein*) adopted a model for O-H oscillators localized on regular oxygen sites and undergoing a stretching mode alongside the "O-O" direction. This situation allows small proton oscillations (*displacements to less than atomic spacing*). Due to their weakness, hydrogen bonds require an activation energy around 0.22 eV [36] to be dissociated even at low temperature. This value agrees well with the one (0.26 eV) determined previously [37] using dielectric measurements at temperature ranging from −25 to −75 °C (*not shown here*). Then, the dielectric response rising at that range can be ascribed to the dissociation of hydrogen bonds and consequently to a localized migration (*or oscillation*) of the protons alongside the hydroxide bond. Such process was previously emphasized and discussed by Weber et al. [35] in ionic conductor oxides.

Additionally, the interfacial conduction process considered at temperature higher than 120 °C for hydrogenated films was thermally activated with around 0.6 eV. It is whispered that the proton diffusion within oxide materials requires almost the same amount of energy. Actually, the reported activation energy should be strongly affected by the amorphous state of the material, but it agrees well with predicted experimental [38, 39] and theoretical [38–40] values (0.4–0.6 eV) for the best proton conductors such as Y-BaZrO3 or BaCeO3 [38, 39, 41]. Based on the experimental observation and on the light of literature mentioned above, the conduction mechanism is predominately of the Grotthuss type [42], that is, involving proton transfer (*hopping*) from the hydroxide defects to a nearest neighboring oxygen ion.

In order to better illustrate that the huge increase of the dielectric constant (**Figure 3**) arises from the accumulation of protons at the metal-BaTiO3:H interface (*double layer*), a tri-layer stack (BaTiO3/BaTiO3:H/BaTiO3) was grown by inserting the hydrogen-doped layer between two intrinsic thin layers. In that way, we can separate protons (*incorporated in hydrogenated layer*) from the metal electrode by an intrinsic (*proton free*) layer. The intrinsic layers (BaTiO3) were grown during 15 min under pure argon gas, whereas the hydrogenated layer (BaTiO3:H) was deposited during 90 min under Ar/H2 gas mixture (80% Ar + 20 % H2) without breaking the deposition process. **Figure 11** reports the *εʹ*(*f*) and *σ*(*f*) characteristics measured at room temperature on barium titanate films deposited under different atmospheres (**Figure 11(a)**) and in different MIM structures (**Figure 11(b)**).

In **Figure 11(a)** gold was used as metal electrode, and curves labeled (A) were conducted on a sample grown under pure argon gas (*reference*), curves labeled (B) were carried out on a sample grown under 20 % HMR, and curves labeled (C) were done on a tri-layer stack (BaTiO3/ BaTiO3:H/BaTiO3). Open symbols denote the dielectric constant *εʹ* and filled symbols represent the conductivity *σ*. In **Figure 11(b)** measurements were carried out on samples grown under 20 % HMR, where gold [*curves labeled* (B)] and copper [*curves labeled* (D)] were used as metal electrodes. Further evidences that the observed frequency dispersion of the dielectric constant is related to the electrodes (*and not to the bulk*) are given. At low frequencies, we show that when an intrinsic layer between the hydrogenated bulk film and metallic electrodes was added, both the conductivity and the dielectric constant decrease while remaining higher compared to values measured on Ar-deposited films. This confirms that the low-frequency behavior is related to electrode effects. We previously [25] discussed the dispersive behavior of the *εʹ*(*f*) characteristic and attributed it to the proton migration. Protons are indeed very mobile species even at low temperature. Therefore, we can assume that upon the application of a sufficiently high electric field, the positively charged protons move toward the cathode and pile up to form a concentration gradient (*capacitive double layer*). The proton migration leads to an increase of the ionic conductivity and creates a capacitive double layer under the cathode which enhances the dielectric constant. In the case of tri-layer stack, the interfacial Ar-deposited layers react as *dead layers*, which reduce the capacitance of the MIM structure as observed at low frequencies.

They effectively act as positively charged defects that repel the proton. This feature allows the formation of a rather weak directional interaction (*hydrogen bond*) between the proton and one of the adjacent oxygen, O[1]-H~O[2] [33]. Such a site could be considered as a trap for protons. In this approach, electron issued from the ionization of hydrogen is weakly bonded, and it can be easily activated to the conduction band, from a shallow donor level estimated at around

The hydrogen bond breakage can be responsible for a large proton conduction process in oxide materials. Several reports (*Ref*. [35] *and references therein*) adopted a model for O-H oscillators localized on regular oxygen sites and undergoing a stretching mode alongside the "O-O" direction. This situation allows small proton oscillations (*displacements to less than atomic spacing*). Due to their weakness, hydrogen bonds require an activation energy around 0.22 eV [36] to be dissociated even at low temperature. This value agrees well with the one (0.26 eV) determined previously [37] using dielectric measurements at temperature ranging from −25 to −75 °C (*not shown here*). Then, the dielectric response rising at that range can be ascribed to the dissociation of hydrogen bonds and consequently to a localized migration (*or oscillation*) of the protons alongside the hydroxide bond. Such process was previously emphasized and

Additionally, the interfacial conduction process considered at temperature higher than 120 °C for hydrogenated films was thermally activated with around 0.6 eV. It is whispered that the proton diffusion within oxide materials requires almost the same amount of energy. Actually, the reported activation energy should be strongly affected by the amorphous state of the material, but it agrees well with predicted experimental [38, 39] and theoretical [38–40] values (0.4–0.6 eV) for the best proton conductors such as Y-BaZrO3 or BaCeO3 [38, 39, 41]. Based on the experimental observation and on the light of literature mentioned above, the conduction mechanism is predominately of the Grotthuss type [42], that is, involving proton transfer

In order to better illustrate that the huge increase of the dielectric constant (**Figure 3**) arises from the accumulation of protons at the metal-BaTiO3:H interface (*double layer*), a tri-layer stack (BaTiO3/BaTiO3:H/BaTiO3) was grown by inserting the hydrogen-doped layer between two intrinsic thin layers. In that way, we can separate protons (*incorporated in hydrogenated layer*) from the metal electrode by an intrinsic (*proton free*) layer. The intrinsic layers (BaTiO3) were grown during 15 min under pure argon gas, whereas the hydrogenated layer (BaTiO3:H) was deposited during 90 min under Ar/H2 gas mixture (80% Ar + 20 % H2) without breaking the deposition process. **Figure 11** reports the *εʹ*(*f*) and *σ*(*f*) characteristics measured at room temperature on barium titanate films deposited under different atmospheres (**Figure 11(a)**)

In **Figure 11(a)** gold was used as metal electrode, and curves labeled (A) were conducted on a sample grown under pure argon gas (*reference*), curves labeled (B) were carried out on a sample grown under 20 % HMR, and curves labeled (C) were done on a tri-layer stack (BaTiO3/ BaTiO3:H/BaTiO3). Open symbols denote the dielectric constant *εʹ* and filled symbols represent the conductivity *σ*. In **Figure 11(b)** measurements were carried out on samples grown under 20 % HMR, where gold [*curves labeled* (B)] and copper [*curves labeled* (D)] were used as metal

(*hopping*) from the hydroxide defects to a nearest neighboring oxygen ion.

0.15 eV, to increase the dielectric loss and the leakage current [34].

100 Supercapacitor Design and Applications

discussed by Weber et al. [35] in ionic conductor oxides.

and in different MIM structures (**Figure 11(b)**).

**Figure 11.** Frequency dependence of the dielectric constant and the conductivity. Measurements were carried out at room temperature on BaTiO3-based MIM capacitors grown under different (a) atmospheres and in different (b) MIM (M = Au, Cu) structures. Curves labeled (A) were conducted on a sample grown under pure argon gas (*reference*), curves labeled (B, D) were carried out on a sample grown under 20 % HMR, and curves labeled (C) were done on a trilayer stack (BaTiO3/BaTiO3:H/BaTiO3). Open symbols denote the dielectric constant and filled symbols represent the conductivity.

To further check the influence of electrodes, the nature of the metal contact was varied (**Figure 11(b)**). Gold (Au/BaTiO3:H/Au) and copper (Cu/BaTiO3:H/Cu) were used as metal electrodes. It is seen that replacing the Au electrode by Cu is sufficient to modify the conductivity, especially in the low-frequency domain, where the conductivity measured on the Cu/ BaTiO3:H/Cu decreases. This clearly shows that the dielectric response is strongly dependent on the electrode nature. Under an electric field, positively charged protons are drifted toward the cathode where, ideally (*if ohmic contact*), they are neutralized by electronic transfer through the electrode/electrolyte interface. In the case of bad charge transfer (*blocking contact*), mobile charges accumulate at the electrode/electrolyte interface, and a space charge region builds up at the electrode boundaries. This feature agrees with decreasing conductivity.

Finally, the last feature which can be observed in **Figure 8** is related to the HT relaxation process (*T* > 120 °C) which can be thermally activated by around 1 eV. This value is close to the usual activation energy reported for oxygen vacancy migration in titanates [43, 44]. These defects, resulting from the deposition under reduced atmosphere, are indeed very mobile species in titanates [43–45]. Therefore, we can assume that upon the application of a low-frequency electric field at sufficiently high temperature, the oxygen vacancies move toward the cathode and pile up at the interface to form a concentration gradient (*space charge*). This feature highly competes with the double layer previously formed by protons which leads to an increase of the electrical conduction with elapsed time.

## **4. Conclusion**

Metal/*a*-BaTiO3:H(1 μm)/metal devices report a specific capacitance around 2 μF/cm2 , steady in the frequency range from 0.1 to 104 Hz. This feature is suitable for ac circuits. On the other hand, these devices have low-temperature performance and are capable of delivering energy down to − 100 °C with minimal effect on efficiency. Moreover, they can bear high voltages per cell. Compared to conventional capacitor technologies, the studied devices possess orders of magnitude higher energy density. This performance can be rather enhanced by using porous carbon electrodes to achieve a high surface area. In addition, these devices display a low internal resistance, hence providing acceptable power density capability compared to batteries. Finally, the deposition process seems to be compatible with easy incorporation of all-solidstate supercapacitors with other electronic devices. In addition, we can stack supercapacitors vertically to enhance the specific capacitance.

## **Author details**

Fadhel El Kamel

Address all correspondence to: elkf2000@yahoo.fr

LaPhyMNE, University of Gabès, Gabès, Tunisia

## **References**


Finally, the last feature which can be observed in **Figure 8** is related to the HT relaxation process (*T* > 120 °C) which can be thermally activated by around 1 eV. This value is close to the usual activation energy reported for oxygen vacancy migration in titanates [43, 44]. These defects, resulting from the deposition under reduced atmosphere, are indeed very mobile species in titanates [43–45]. Therefore, we can assume that upon the application of a low-frequency electric field at sufficiently high temperature, the oxygen vacancies move toward the cathode and pile up at the interface to form a concentration gradient (*space charge*). This feature highly competes with the double layer previously formed by protons which leads to an increase of

Metal/*a*-BaTiO3:H(1 μm)/metal devices report a specific capacitance around 2 μF/cm2

hand, these devices have low-temperature performance and are capable of delivering energy

cell. Compared to conventional capacitor technologies, the studied devices possess orders of magnitude higher energy density. This performance can be rather enhanced by using porous carbon electrodes to achieve a high surface area. In addition, these devices display a low internal resistance, hence providing acceptable power density capability compared to batteries. Finally, the deposition process seems to be compatible with easy incorporation of all-solidstate supercapacitors with other electronic devices. In addition, we can stack supercapacitors

[1] B. E. Conway. Electrochemical Supercapacitors: Scientific, Fundamentals and Technological Applications. New York: Kluwer Academic/Plenum Publishers, 1999.

100 °C with minimal effect on efficiency. Moreover, they can bear high voltages per

Hz. This feature is suitable for ac circuits. On the other

, steady

the electrical conduction with elapsed time.

in the frequency range from 0.1 to 104

vertically to enhance the specific capacitance.

Address all correspondence to: elkf2000@yahoo.fr

LaPhyMNE, University of Gabès, Gabès, Tunisia

[2] L. Ma et al. Appl. Phys. Lett. 2005;87:123503.

[3] M. G. H. Hendriks et al. J. Appl. Phys. 2001;90:5303.

**4. Conclusion**

102 Supercapacitor Design and Applications

**Author details**

Fadhel El Kamel

**References**

down to −


#### **Oxidation-Ultrasound Process on Removing Potassium Ions from Activated Carbon for Improving Electrochemical Properties of Supercapacitor Oxidation-Ultrasound Process on Removing Potassium Ions from Activated Carbon for Improving Electrochemical Properties of Supercapacitor**

Kang Sun Kang Sun

[32] J. P. Boyeaux et al. J. Phys. C. 1979;12:545.

[35] G. Weber et al. Phys. Rev. B. 1986;34:8406.

[38] W. Münch et al. Solid State Ionics. 1999;125:39.

[41] K. D. Kreuer. Solid State Ionics. 1999;125:285.

[43] F. El Kamel et al. J. Appl. Phys. 2006;99:094107.

[44] R. Waser. J. Am. Ceram. Soc. 1989;72:2234. [45] R. Waser. J. Am. Ceram. Soc. 1991;74:1934.

[42] K. D. Kreuer. Chem. Mater. 1996;8:610.

[39] E. Matsushita et al. Solid State Ionics. 1997;97:45. [40] E. Matsushita et al. Solid State Ionics. 1999;125:31.

1970.

[34] L. Dobaczewski et al. Phys. Rev. B. 2004;69:245207.

[37] F. El Kamel et al. J. Vac. Sci. Technol. A. 2012;30:04D110-1.

[36] J. O. M. Bockris and A. K. N. Reddy. Modern Electrochemistry. New York: Plenum;

[33] L. Pejov. Chem. Phys. Lett. 2003;11:376.

104 Supercapacitor Design and Applications

Additional information is available at the end of the chapter Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/64809

#### **Abstract**

To produce pure activated carbon (AC) with low potassium ions (K+ ) content for supercapaci‐ tor, coconut‐shell AC activated by KOH was treated with a novel oxidation‐ultrasound process on the basis of hydrochloric acid (HCl) washing. The electrochemical performances of the ACs as supercapacitor electrodes were characterized by cyclic voltammetry (CV), galvanostatic charge‐discharge (GC), and electrochemical impedance spectroscopy (EIS). Results showed that the obtained AC, which was washed with 1.0 wt% HCl solution for 120 min and subsequently treated with 0.6 wt% H2O2 solution at 60°C in an ultrasonic oscillator for 8 h, possessed a K+ content of 46 mg/kg, much lower than that of 417 mg/kg of the AC without oxidation‐ ultrasound treatment. Furthermore, a large specific surface area and pore volume of 3460 m2 /g and 1.869 cm3 /g, respectively, were obtained for AC after oxidation‐ultrasound treatment. A high specific capacitance of 306 F/g at the current density of 1 A/g in 1 M H2SO4 electrolyte was observed for the prepared AC, indicating it has good electrochemical performances, and remained at 294 F/g with a capacitance retention of 96% after 3000 cycles, indicating excellent stability and capacitive behavior of the AC electrode for supercapacitor.

**Keywords:** activated carbon, potassium ions, oxidation‐ultrasound, specific capaci‐ tance, supercapacitor

## **1. Introduction**

As a kind of new energy storage device between battery and conventional capacitor, super‐ capacitor has drawn much attention recently due to its high power density and long cycle life [1–4]. Besides, it was assumed to be one of the most promising energy storage devices for

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

various applications such as brake energy recovery systems, hybrid electrical vehicles, digital telecommunication systems [5, 6], etc. It is known that the surface area, pore size, and purity texture have significant influence on the performances of supercapacitor for the electrode material [7]. So far, AC has been widely applied in electrode material for supercapacitor applications, owing to its advantages of good chemical stability and conductivity, large surface area, low cost [8–11], etc. However, the AC products show a high K+ content due to the process limitations of KOH activation, as a result, the already‐formed AC pores were blocked, which had a negative effect on the specific capacitance and cycle life of supercapacitor. Therefore, it is necessary to remove the K+ from the AC, aiming to greatly improve the electrochemical performances of supercapacitor.

In general, the K+ existed in the AC often spreads in the pores, and HCl washing cannot efficiently remove it. Although ultrasound oscillation was effective to disperse K+ into solu‐ tion, oxidation process is able to weaken the force between AC pore walls and K+ , resulting in deeply purification of the AC samples. Therefore, this study aims to make use of a novel oxidation‐ultrasound process to reduce the K+ content of ACs. The effect of K+ content on AC properties will be reflected by the physical and electrochemical performances.

## **2. Experimental**

#### **2.1. Preparation of AC sample**

In this study, coconut shell was used to prepare AC by KOH activation. First of all, the coconut shell was subjected to be carbonized at 600°C for 1 h and was crushed to obtain a particle size of about 1 mm. Then 50 wt% KOH solution with the KOH/C weight ratio of 4:1 was mixed with the coconut shell carbide in a stainless steel reactor. In addition, the mixture was activated in a muffle furnace in which the mixture was first pretreated at 350°C for 2 h under air atmosphere, then heated up to 800°C at a heating rate of 10°C min−1, and maintained at 800°C for 1 h under sealed condition. After cooled down to the room temperature, the resultant samples were subsequently washed with distilled water until a pH of approximately 6 was obtained. Finally, the prepared AC samples were dried at 120°C for 10 h under vacuum condition.

#### **2.2. Removal of the K+ from AC sample**

There are two steps included in the removal process of the K+ . In the first step, 2.0 g as‐ob‐ tained AC was impregnated with 100 ml 0.3, 0.5, 0.7, 1.0, and 1.2 wt% of HCl solution, re‐ spectively. Then the mixture was stirred at 80°C for 30–150 min in an electro‐thermostatic water cabinet. The oxidation‐ultrasound process was used to purify the HCl‐washed AC deeply in the second step. In an ultrasonic oscillator, 100 ml different weight ratios 0, 0.2, 0.4, 0.6, 0.8, and 1.0 wt% of H2O2 solution were mixed with 2.0 g HCl‐washed AC for 2–10 h at 10–80°C. The final treated samples were dried at 120°C under vacuum condition over‐ night. After that, the resulting AC samples were named AC‐W, AC‐W‐H, and AC‐W‐H‐U, corresponding to the water, HCl, and oxidation‐ultrasound treatment on AC, respectively.

#### **2.3. Characterizations**

various applications such as brake energy recovery systems, hybrid electrical vehicles, digital telecommunication systems [5, 6], etc. It is known that the surface area, pore size, and purity texture have significant influence on the performances of supercapacitor for the electrode material [7]. So far, AC has been widely applied in electrode material for supercapacitor applications, owing to its advantages of good chemical stability and conductivity, large surface area, low cost [8–11], etc. However, the AC products show a high K+ content due to the process limitations of KOH activation, as a result, the already‐formed AC pores were blocked, which had a negative effect on the specific capacitance and cycle life of supercapacitor. Therefore, it is necessary to remove the K+ from the AC, aiming to greatly improve the electrochemical

In general, the K+ existed in the AC often spreads in the pores, and HCl washing cannot efficiently remove it. Although ultrasound oscillation was effective to disperse K+ into solu‐

in deeply purification of the AC samples. Therefore, this study aims to make use of a novel oxidation‐ultrasound process to reduce the K+ content of ACs. The effect of K+ content on

In this study, coconut shell was used to prepare AC by KOH activation. First of all, the coconut shell was subjected to be carbonized at 600°C for 1 h and was crushed to obtain a particle size of about 1 mm. Then 50 wt% KOH solution with the KOH/C weight ratio of 4:1 was mixed with the coconut shell carbide in a stainless steel reactor. In addition, the mixture was activated in a muffle furnace in which the mixture was first pretreated at 350°C for 2 h under air atmosphere, then heated up to 800°C at a heating rate of 10°C min−1, and maintained at 800°C for 1 h under sealed condition. After cooled down to the room temperature, the resultant samples were subsequently washed with distilled water until a pH of approximately 6 was obtained. Finally, the prepared AC samples were dried at 120°C for 10 h under vacuum

tained AC was impregnated with 100 ml 0.3, 0.5, 0.7, 1.0, and 1.2 wt% of HCl solution, re‐ spectively. Then the mixture was stirred at 80°C for 30–150 min in an electro‐thermostatic water cabinet. The oxidation‐ultrasound process was used to purify the HCl‐washed AC deeply in the second step. In an ultrasonic oscillator, 100 ml different weight ratios 0, 0.2, 0.4, 0.6, 0.8, and 1.0 wt% of H2O2 solution were mixed with 2.0 g HCl‐washed AC for 2–10 h at 10–80°C. The final treated samples were dried at 120°C under vacuum condition over‐ night. After that, the resulting AC samples were named AC‐W, AC‐W‐H, and AC‐W‐H‐U, corresponding to the water, HCl, and oxidation‐ultrasound treatment on AC, respectively.

, resulting

. In the first step, 2.0 g as‐ob‐

tion, oxidation process is able to weaken the force between AC pore walls and K+

AC properties will be reflected by the physical and electrochemical performances.

performances of supercapacitor.

106 Supercapacitor Design and Applications

**2. Experimental**

condition.

**2.1. Preparation of AC sample**

**2.2. Removal of the K+ from AC sample**

There are two steps included in the removal process of the K+

Inductively coupled plasma mass spectrometry (ICP‐MS) was used to measure the K+ con‐ tent of ACs. Scanning electron microscopy (SEM, S‐3400, Hitachi) was used to exhibit the surface morphologies of the samples. Nitrogen (77 K) adsorption was carried out using a Micromeritics ASAP 2020 analyzer to detect the porous properties of ACs. The specific sur‐ face area values, total pore volume, and the pore size distribution of activated carbons were calculated by the BET method, according to the adsorbed amount and density functional theory (DFT), respectively.

#### **2.4. Electrochemical characterization**

The capacitive performance measurements were carried out at room temperature in a three‐ electrode system on an electrochemical workstation (Bio‐Logic, France) in 1 M H2SO4 electro‐ lyte. The active AC materials, acetylene black, and 5 wt% PTFE with a weight ratio of 8.5:1:0.5 were fully mixed and grounded in a mortar. Then the mixture was pressed onto a Ni‐foam current collector to form the working electrode. And Ag/AgCl and platinum foil were employed as the reference and counter electrode, respectively. The capacitive behavior of the AC electrode was characterized by CV, GCD, and electrochemical impedance spectroscopy (EIS) measurements. CV was performed from 0.01 to 0.9 V at various scan rates in a range of 5–200 mV s−1, GC curve was measured from 0.01 to 0.9 V at different current densities ranging from 0.1 to 5 Ag−1, and EIS was recorded at the frequency from 100 kHz to 10 MHz.

## **3. Results and discussion**

#### **3.1. Influencing factors on removal process of K+**

#### *3.1.1. Effect of HCl treatment on the K+ content of AC*

In this paper, HCl solution is first used to treat with the AC‐W samples to remove the K+ . The effects of the treatment time and mass fraction of HCl solution are investigated in detail. It can be seen in **Figure 1(a)** that the relationship between the mass fraction of HCl solution and K+ content for 90 min. And the figure also shows the K+ content exhibits a conspicuous decrease with the increasing mass fraction of HCl solution from 0 to 1.0 wt%, and then it presents a rising trend. The excessive acid radical ions results in K+ removed from AC difficulty during electric double layer process, which is formed on the surface of K+ . Thus, the optimal mass fraction of HCl solution is 1.0 wt%.

It shows the relationship between K+ content and the treatment time of HCl solution at the mass fraction of 1.0 wt% in **Figure 1(b)**. The content of K+ decreases from 1242 to 417 mg/kg significantly, while the treatment time increases from 30 min to 2 h, and it does not change anymore after 2 h. Therefore, treatment time of 120 min and 1.0 wt% HCl solution are chosen as the optimal treatment conditions for the whole experiments, and the corresponding AC sample is named AC‐W‐H1%/120min.

**Figure 1.** Effect of mass fraction and treatment time of HCl solution on the K+ content of AC.

#### *3.1.2. Effect of oxidation-ultrasound process on the K+ content of AC*

The AC‐W‐H1%/120min is then subjected to oxidation‐ultrasound process to deeply remove the K+ . The influence of the mass fraction of H2O2 solution, ultrasound temperature, and ultra‐ sound time on the K+ content of ACs is studied. The obtained results are shown in **Table 1** and **Figure 2**. The mass fraction of H2O2 solution is in the range of 0–1.0 wt% at an ultra‐ sound time of 6 h and an ultrasound temperature of 60°C. **Table 1** indicates that the K+ con‐ tent of AC without H2O2 treatment is 256 mg/kg, substantially lower than that of the AC‐W‐ H1%/120min with 417 mg/kg, which should have close relationships with the ultrasound effect that the ultrasonic wave can accelerate the mass transfer process of them between the solid‐ liquid phase and promote some uncarbonized substances inside the AC pores to be dis‐ persed into the solution along with the K+ . It can be seen from **Figure 2(a)** that the K+ content of AC decreases markedly from 256 to 61 mg/kg, and thereafter it remains substantially un‐ changed when the mass fraction of H2O2 solution increased from 0 to 0.6 wt%. Therefore, it can be stated that H2O2 treatment shows a great effect on producing ACs with low K+ con‐ tent. On the one hand, as an oxidizing agent, H2O2 solution can reduce the force between the K+ and AC pore walls, leading to it transferred to the solution easily. On the other hand, the K+ in the AC can coexist with some organic groups in the form of chemical bond, which can be destroyed by the presence of H2O2 solution, making it possible for the K+ of ACs to be removed effectively [12, 13].


**Table 1.** Results of mass fraction of H2O2 solution, ultrasound temperature and ultrasound time on the K+ content of AC.

Oxidation-Ultrasound Process on Removing Potassium Ions from Activated Carbon for Improving Electrochemical... http://dx.doi.org/10.5772/64809 109

**Figure 1.** Effect of mass fraction and treatment time of HCl solution on the K+

*3.1.2. Effect of oxidation-ultrasound process on the K+*

108 Supercapacitor Design and Applications

persed into the solution along with the K+

removed effectively [12, 13].

K+

K+

K+

K+

K+

K+

AC.

content of AC.

. It can be seen from **Figure 2(a)** that the K+

con‐

content

con‐

content of

 *content of AC*

The AC‐W‐H1%/120min is then subjected to oxidation‐ultrasound process to deeply remove the

tent of AC without H2O2 treatment is 256 mg/kg, substantially lower than that of the AC‐W‐ H1%/120min with 417 mg/kg, which should have close relationships with the ultrasound effect that the ultrasonic wave can accelerate the mass transfer process of them between the solid‐ liquid phase and promote some uncarbonized substances inside the AC pores to be dis‐

of AC decreases markedly from 256 to 61 mg/kg, and thereafter it remains substantially un‐ changed when the mass fraction of H2O2 solution increased from 0 to 0.6 wt%. Therefore, it can be stated that H2O2 treatment shows a great effect on producing ACs with low K+

tent. On the one hand, as an oxidizing agent, H2O2 solution can reduce the force between the

and AC pore walls, leading to it transferred to the solution easily. On the other hand, the

 in the AC can coexist with some organic groups in the form of chemical bond, which can be destroyed by the presence of H2O2 solution, making it possible for the K+ of ACs to be

Mass fraction of H2O2/wt% 0 0.2 0.4 0.6 0.8 1.0

 content/(mg/kg) 256 184 92 61 64 75 Ultrasonic temperature/°C 10 20 40 60 70 80

 content/(mg/kg) 157 124 96 61 74 91 Ultrasonic time/h 2 4 6 8 10 12

content/(mg/kg) 148 119 61 46 41 44

**Table 1.** Results of mass fraction of H2O2 solution, ultrasound temperature and ultrasound time on the K+

. The influence of the mass fraction of H2O2 solution, ultrasound temperature, and ultra‐ sound time on the K+ content of ACs is studied. The obtained results are shown in **Table 1** and **Figure 2**. The mass fraction of H2O2 solution is in the range of 0–1.0 wt% at an ultra‐ sound time of 6 h and an ultrasound temperature of 60°C. **Table 1** indicates that the K+

**Figure 2.** Effect of mass fraction of (a) H2O2 solution, (b) ultrasound temperature and (c) ultrasound time on the K+ content of AC.

**Figure 2(b)** exhibits the relationship between K+ content and the ultrasound temperature in the range of 10~80°C at an ultrasound time of 6 h, using 0.6 wt% H2O2 solution. It was observed that the K+ content of AC decreased from 157 to 61 mg/kg when the temperature increased from 10 to 60°C, but increased while the temperature was further increased from 60 to 80°C. This might be attributed to the fact that the energy and activity of the K+ were improved with the increase of temperature in ultrasonic field and the movement of K+ into the solution was accelerated. **Figure 2(c)** shows the effect of ultrasound time on the K+ content of AC, using 0.6 wt% H2O2 solution at the ultrasound temperature of 60°C. It is clear that the K+ content drops from 148 to 46 mg/kg as the ultrasound time increases from 2 to 8 h, and it scarcely changes with a continuing increase of ultrasound time. Therefore, we consider 0.6 wt% H2O2 solution, ultrasound temperature of 60°C and ultrasound time of 8 h as the optimum experi‐ ment conditions to deeply remove the K+ from AC, and the corresponding AC is marked as AC‐W‐H1%/120min‐U0.6%/60°C/8h.

#### **3.2. Sample characterization**

#### *3.2.1. Morphology characterization*

**Figure 3**(**a**–**c**) shows the surface morphologies of the AC‐W, AC‐W‐H1%/120min, and AC‐W‐ H1%/120min‐U0.6%/60°C/8h by SEM at 1000× magnification. It can be seen clearly that the mor‐ phologies of the AC samples display an obvious change on the surface. The AC‐W has large quantities of uncarbonized substances presented on the surface. The AC‐W‐H1%/120min has relatively flat and smooth surface compared to the AC‐W, but there are still some residues on the surface. Meanwhile, as shown in **Figure 3(c)**, the surface of the AC‐W‐ H1%/120min‐U0.6%/60°C/8h was very clean and smooth, confirming that the oxidation‐ultrasound process has a significant effect on the purification of AC, which is consistent with the aforementioned results of the K+ content.

**Figure 3.** SEM images of (a) AC‐W, (b) AC‐W‐H1%/120min and (c) AC‐W‐H1%/120min‐U0.6%/60°C/8h.

#### *3.2.2. Porous texture characterization*

It was reported that the specific surface area and pore size of AC were critical factors for electrochemical capacitor applications [14], hence, the N2 adsorption‐desorption isotherms and corresponding pore size distribution curves of the as‐prepared ACs were analyzed as shown in **Figure 4**. **Table 2** gives the textural characteristics and specific capacitance of the samples. It was noticed that the presented isotherms were type I for all samples according to IUPAC classification [15, 16]. The N2 adsorption volume of the AC‐W achieved saturation at low relative pressure of 0.1 and the lowest isotherm was observed, suggesting that a great number of activating agents (KOH) are still blocking in the AC pores, especially the mesopores. The AC‐W‐H1%/120min displays significantly higher isotherm than the AC‐W and its N2 adsorption volume shows a gradual increase until the relative pressure of 0.4, implying that most of the KOH remained in the AC pores are removed by HCl washing. As for the AC‐W‐H1%/120min‐U0.6%/ 60°C/8h. The existence of a hysteresis loop in the desorption branch at the relative pressure of 0.47 was observed obviously. And the isotherm with a highest adsorbed volume indicated that the deep removal of the K+ along with other residues from the AC pores by the oxidation‐ ultrasound process greatly increased the mesopore volumes.

**3.2. Sample characterization**

110 Supercapacitor Design and Applications

*3.2.1. Morphology characterization*

aforementioned results of the K+ content.

**Figure 3.** SEM images of (a) AC‐W, (b) AC‐W‐H1%/120min and (c) AC‐W‐H1%/120min‐U0.6%/60°C/8h.

It was reported that the specific surface area and pore size of AC were critical factors for electrochemical capacitor applications [14], hence, the N2 adsorption‐desorption isotherms and corresponding pore size distribution curves of the as‐prepared ACs were analyzed as shown in **Figure 4**. **Table 2** gives the textural characteristics and specific capacitance of the samples. It was noticed that the presented isotherms were type I for all samples according to IUPAC

*3.2.2. Porous texture characterization*

**Figure 3**(**a**–**c**) shows the surface morphologies of the AC‐W, AC‐W‐H1%/120min, and AC‐W‐ H1%/120min‐U0.6%/60°C/8h by SEM at 1000× magnification. It can be seen clearly that the mor‐ phologies of the AC samples display an obvious change on the surface. The AC‐W has large quantities of uncarbonized substances presented on the surface. The AC‐W‐H1%/120min has relatively flat and smooth surface compared to the AC‐W, but there are still some residues on the surface. Meanwhile, as shown in **Figure 3(c)**, the surface of the AC‐W‐ H1%/120min‐U0.6%/60°C/8h was very clean and smooth, confirming that the oxidation‐ultrasound process has a significant effect on the purification of AC, which is consistent with the

**Figure 4.** (a) N2 adsorption‐desorption isotherms and (b) the corresponding pore size distributions of the as‐prepared samples.


*Note*: SBET: specific surface area; Vt : total pore volume; Vmeso: mesopore volume; Vmicro: micropore volume; Dp 1 : average pore size; Cg: specific gravimetric capacitance; yield: the independent result of every step for removing K+ 2 .

3 **Table 2.** Textural characteristics and specific capacitance of the as‐prepared samples.

4 **Figure 4(b)** gives the pore size distributions of the as‐prepared ACs. Results showed that the pore size of AC‐W was in the range of 0.6–2 nm, while that of AC‐W‐H1%/120min‐U0.6%/60°C/8h 5 and AC‐W‐H1%/120min 6 ranged from 2 to 4 nm and 0.5 to 2.0 nm, respectively. It is obvious that the 7 purification of the AC‐W with HCl and oxidation‐ultrasound treatment in turn gradually made 8 mesopores (2–4 nm) become the dominant type of pores.

9 It can be seen from **Table 2** that the AC‐W had an average pore size of 1.781 nm a mesopore volume of 0.311 cm3 /g, while the AC‐W‐H1%/120min and AC‐W‐H1%/120min 10 ‐U0.6%/60°C/8h achieve the mesopore volume of 1.297 and 1.615 cm3 11 /g, and the average pore size of 2.086 and 2.228 nm, 12 respectively. Moreover, the AC‐W‐H1%/120min has significantly larger surface area and total pore volume than the AC‐W, increasing respectively from 2198 to 3178 m2 13 /g and 1.110 to 1.686 cm3 /g, and the AC‐W‐H1%/120min 14 ‐U0.6%/60°C/8h possesses the largest surface area and pore volume of 3460 m2 /g and 1.869 cm3 /g. It is evident that the removal of K+ 15 from AC samples 16 creates some new microspores and obviously promotes the widening of existing micropores 17 into mesopores at the same time, which can be accessible to electrolyte ions for electrical double‐layer formation [17]. Thus, the large surface area, high pore volume, and low K+ 18 con‐ 19 tent of AC samples make them good candidates for electrode materials.

#### 20 **3.3. Electrochemical characterization**

21 CV is used in determination of the electrochemical performances of as‐prepared samples. The typical CV results of AC‐W, AC‐W‐H1%/120min and AC‐W‐H1%/120min 22 ‐U0.6%/60°C/8h at the scan rate of 23 10 mV/s with the potential range of −0.2~0.8 V are shown in **Figure 5(a)**. Results showed that 24 all the AC electrodes displayed an approximately quasi‐rectangular voltammogram shape, 25 which was the characteristic of electrochemical double‐layer capacitance [18–20]. Furthermore, the CV curve of AC‐W‐H1%/120min 26 ‐U0.6%/60°C/8h electrode exhibits a bigger current response and larger area of rectangle than that of the AC‐W and AC‐W‐H1%/120min 27 electrode, demonstrating an obvious increase in capacity during the removal process of K+ 28 . **Figure 5(b)** shows the CV curves of the AC‐W‐H1%/120min 29 ‐U0.6%/60°C/8h electrode at different scan rates of 5–50 mV/s. The CV 30 curve eventually becomes titled with the increase of the scan rate, but still maintains a 31 rectangular‐like shape even at 50 mV/s, implying a small resistance in the accessible pores and an excellent capacitive behavior [21, 22] of the AC‐W‐H1%/120min‐U0.6%/60°C/8h 32 electrode.

Oxidation-Ultrasound Process on Removing Potassium Ions from Activated Carbon for Improving Electrochemical... <sup>9</sup> Oxidation-Ultrasound Process on Removing Potassium Ions from Activated Carbon for Improving Electrochemical... http://dx.doi.org/10.5772/64809 

**Samples SBET**

8 Supercapacitor Design and Applications Supercapacitor Design and Applications

*Note*: SBET: specific surface area; Vt

volume of 0.311 cm3

volume of 3460 m2

**3.3. Electrochemical characterization**

1.686 cm3

AC‐W‐H1%/120min‐ U0.6%/60°C/8h

**(m2 /g)** **Vt (cm3 /g)**

**Table 2.** Textural characteristics and specific capacitance of the as‐prepared samples.

mesopores (2–4 nm) become the dominant type of pores.

/g and 1.869 cm3

tent of AC samples make them good candidates for electrode materials.

**Vmeso (cm3 /g)**

AC‐W 2198 1.110 0.311 0.799 1.781 147 1.2 — AC‐W‐H1%/120min 3178 1.686 1.297 0.389 2.086 255 0.1 94.3

: total pore volume; Vmeso: mesopore volume; Vmicro: micropore volume; Dp : average pore size; Cg: specific gravimetric capacitance; yield: the independent result of every step for removing K+ .

 **Figure 4(b)** gives the pore size distributions of the as‐prepared ACs. Results showed that the pore size of AC‐W was in the range of 0.6–2 nm, while that of AC‐W‐H1%/120min‐U0.6%/60°C/8h and AC‐W‐H1%/120min ranged from 2 to 4 nm and 0.5 to 2.0 nm, respectively. It is obvious that the purification of the AC‐W with HCl and oxidation‐ultrasound treatment in turn gradually made

It can be seen from **Table 2** that the AC‐W had an average pore size of 1.781 nm a mesopore

/g, while the AC‐W‐H1%/120min and AC‐W‐H1%/120min ‐U0.6%/60°C/8h achieve the mesopore volume of 1.297 and 1.615 cm3 /g, and the average pore size of 2.086 and 2.228 nm, respectively. Moreover, the AC‐W‐H1%/120min has significantly larger surface area and total pore volume than the AC‐W, increasing respectively from 2198 to 3178 m2 /g and 1.110 to

/g, and the AC‐W‐H1%/120min ‐U0.6%/60°C/8h possesses the largest surface area and pore

/g. It is evident that the removal of K+ from AC samples creates some new microspores and obviously promotes the widening of existing micropores into mesopores at the same time, which can be accessible to electrolyte ions for electrical double‐layer formation [17]. Thus, the large surface area, high pore volume, and low K+ con‐

 CV is used in determination of the electrochemical performances of as‐prepared samples. The typical CV results of AC‐W, AC‐W‐H1%/120min and AC‐W‐H1%/120min ‐U0.6%/60°C/8h at the scan rate of 10 mV/s with the potential range of −0.2~0.8 V are shown in **Figure 5(a)**. Results showed that all the AC electrodes displayed an approximately quasi‐rectangular voltammogram shape, which was the characteristic of electrochemical double‐layer capacitance [18–20]. Furthermore, the CV curve of AC‐W‐H1%/120min ‐U0.6%/60°C/8h electrode exhibits a bigger current response and larger area of rectangle than that of the AC‐W and AC‐W‐H1%/120min electrode, demonstrating an obvious increase in capacity during the removal process of K+ . **Figure 5(b)** shows the CV curves of the AC‐W‐H1%/120min ‐U0.6%/60°C/8h electrode at different scan rates of 5–50 mV/s. The CV curve eventually becomes titled with the increase of the scan rate, but still maintains a rectangular‐like shape even at 50 mV/s, implying a small resistance in the accessible pores and

an excellent capacitive behavior [21, 22] of the AC‐W‐H1%/120min‐U0.6%/60°C/8h electrode.

**Vmicro (cm3 /g)**

1.869 1.615 0.254 2.228 306 0.01 97.5

**Dp (nm)** **Cg (F/g)** **Ash (wt.%)** **Yield (%)**

**Figure 5.** Cyclic voltammograms of the samples: (a) all AC electrodes at a scan rate of 10 mV/s and (b) AC‐W‐H1%/120min ‐ U0.6%/60°C/8h with different scan rates.

 GC measurements have also been conducted to investigate the electrochemical performances of as‐prepared samples. **Figure 6(a)** shows the GC curves of the AC‐W, AC‐W‐H1%/120min and AC‐W‐H1%/120min‐U0.6%/60°C/8h at the current density of 1 A/g with the potential range of −0.2–0.8 V. It can be seen that all the samples present a virtually linear shape and isosceles triangle curve, indicating a good reversibility and typically capacitive behavior of the AC electrodes [23, 24]. According to the GC curves, the AC‐W‐H1%/120min ‐U0.6%/60°C/8h has longer discharging time compared to the AC‐W and AC‐W‐H1%/120min , implying larger specific capacitance, which proved that the deep removal of K+ from AC samples had a significant effect on the electro‐ chemical performance of AC electrode.

**Figure 6.** Galvanostatic charge/discharge curves of the samples: (a) all activated carbon electrodes at a current density of 1 A/g and (b) AC‐W‐H1%/120min‐U0.6%/60°C/8h with different current densities.

**Figure 6(b)** represents the GC curves of the AC‐W‐H1%/120min‐U0.6%/60°C/8h measured in a current density range of 0.5–5 A/g. It was obvious that the discharging time dropped as the current density increased, while the curve was constant as typical triangle shape even at a high loading current density of 5 A/g, revealing that the AC‐W‐H1%/120min‐U0.6%/60°C/8h as electrode material is promising for a high performance supercapacitor. In addition, the specific capacitance of the AC electrodes can be calculated from the charge‐discharge curves based on the following equation [25, 26]:

Oxidation-Ultrasound Process on Removing Potassium Ions from Activated Carbon for Improving Electrochemical... http://dx.doi.org/10.5772/64809 115

$$C\_g = 2\frac{I \cdot \Delta t}{m \cdot \Delta V} \tag{1}$$

where *C*g is the specific gravimetric capacitance (F/g), *I* is the current loaded, *∆t* is the discharge time (s), *∆V* is the potential change during the discharge process, and *m* (g) represents the mass of the AC. As shown in **Table 2**, the AC‐W‐H1%/120min‐U0.6%/60°C/8h has the highest specific gravi‐ metric capacitance of 306 F/g, which is more than twice as large as that of the AC‐W (147 F/g), and increases by about 20% compared with the AC‐W‐H1%/120min (255 F/g). The enhanced specific capacitance of AC samples can be ascribed to the enhancement of the surface area, effective pore volume and purity, as confirmed by the above results of K+ content and N2 adsorption‐desorption isotherms.

**Figure 7.** Cyclic performances of the samples at a current density of 1 A/g.

**Figure 6.** Galvanostatic charge/discharge curves of the samples: (a) all activated carbon electrodes at a current density

**Figure 6(b)** represents the GC curves of the AC‐W‐H1%/120min‐U0.6%/60°C/8h measured in a current density range of 0.5–5 A/g. It was obvious that the discharging time dropped as the current density increased, while the curve was constant as typical triangle shape even at a high loading current density of 5 A/g, revealing that the AC‐W‐H1%/120min‐U0.6%/60°C/8h as electrode material is promising for a high performance supercapacitor. In addition, the specific capacitance of the AC electrodes can be calculated from the charge‐discharge curves based on the following

of 1 A/g and (b) AC‐W‐H1%/120min‐U0.6%/60°C/8h with different current densities.

equation [25, 26]:

114 Supercapacitor Design and Applications

**Figure 7** shows the cyclic stability of AC‐W, AC‐W‐H1%/120min and AC‐W‐H1%/120min‐U0.6%/60°C/8h electrodes detected by galvanostatic charge‐discharge at the current density of 1 A/g for 3000 cycles. The specific capacitance of AC‐W decreased obviously from 147 to 125 F/g after 1000 cycles and was maintained at about 116 F/g with the capacitance retention of 79% after 3000 cycles. The specific capacitance of AC‐W‐H1%/120min is relatively stable and achieves capacitance retention of 91% after 3000 cycles. The specific capacitance for AC‐W‐H1%/120min‐U0.6%/60°C/8h is found to be 306 F/g in the first cycle and 294 F/g after 3000 cycles with a coulombic efficiency of 96%, revealing the excellent stability and reversibility of the AC‐W‐H1%/120min‐U0.6%/60°C/8h electrode.

**Figure 8.** Nyquist plots of the as‐prepared samples (inset: enlarged high‐frequency region of Nyquist plots).

EIS measurements are employed to obtain impedance performance information. Typical Nyquist plots for AC‐W, AC‐W‐H1%/120min and AC‐W‐H1%/120min‐U0.6%/60°C/8h electrodes are presented in **Figure 8**. Results revealed that a semicircle in the high‐frequency region and a straight line in the low‐frequency region were observed in all the three plots [27]. And the plot of AC‐W‐H1%/120min‐U0.6%/60°C/8h sample displays a relatively smaller semicircle than that of AC‐ W and AC‐ W‐H1%/120min samples in the high‐frequency region, indicating a lower charge transfer resistance. This may be ascribed to the larger mesopore volume and lower K+ content of AC‐W‐H1%/120min‐U0.6%/60°C/8h, which can facilitate the rapid diffusion of electrolyte ions into the pores of electrode materials. In addition, the AC‐W‐H1%/120min‐U0.6%/60°C/8h shows more vertical line leaning to imaginary axis in the low‐frequency region [28], suggesting better capacitive behavior than AC‐W and AC‐W‐H1%/120min.

#### **4. Conclusions**

In this work, ACs with high performance for supercapacitor are prepared from coconut shell by KOH activation, using a novel oxidation‐ultrasound process to deeply remove the K+ of AC. The experiment results demonstrate that the AC samples washed with 1.0 wt% HCl solution for 120 min and subsequently treated with 0.6 wt% H2O2 solution at 60°C in an ultrasonic oscillator for 8 h possess a very low K+ content of 46 mg/kg. Compared to the AC‐W and AC‐W‐H1%/120min, the AC‐W‐H1%/120min‐U0.6%/60°C/8h exhibits a larger surface area and pore volume of 3460 m2 /g and 1.869 cm3 /g. As the electrode material for electrochemical application, the AC‐W‐H1%/120min‐U0.6%/60°C/8h showed a high specific capacitance of 306 F/g with a coulombic efficiency of 96% after 3000 cycles. The oxidation‐ultrasound process has a great potential to produce ACs with high performance for supercapacitor applications.

## **Acknowledgements**

This research is financially supported by the projects in Forestry Public Benefit Research Sector (Grant No. 201404610) and the National Natural Science Foundation of China (Grant No. 31400510).

## **Author details**

Kang Sun

Address all correspondence to: sunkang0226@163.com

Institute of Chemical Industry of Forest Products, CAF, Nanjing, China

## **References**

content

content of 46 mg/kg. Compared to the

/g. As the electrode material for electrochemical

**Figure 8.** Nyquist plots of the as‐prepared samples (inset: enlarged high‐frequency region of Nyquist plots).

transfer resistance. This may be ascribed to the larger mesopore volume and lower K+

behavior than AC‐W and AC‐W‐H1%/120min.

an ultrasonic oscillator for 8 h possess a very low K+

/g and 1.869 cm3

**4. Conclusions**

116 Supercapacitor Design and Applications

pore volume of 3460 m2

EIS measurements are employed to obtain impedance performance information. Typical Nyquist plots for AC‐W, AC‐W‐H1%/120min and AC‐W‐H1%/120min‐U0.6%/60°C/8h electrodes are presented in **Figure 8**. Results revealed that a semicircle in the high‐frequency region and a straight line in the low‐frequency region were observed in all the three plots [27]. And the plot of AC‐W‐H1%/120min‐U0.6%/60°C/8h sample displays a relatively smaller semicircle than that of AC‐ W and AC‐ W‐H1%/120min samples in the high‐frequency region, indicating a lower charge

of AC‐W‐H1%/120min‐U0.6%/60°C/8h, which can facilitate the rapid diffusion of electrolyte ions into the pores of electrode materials. In addition, the AC‐W‐H1%/120min‐U0.6%/60°C/8h shows more vertical line leaning to imaginary axis in the low‐frequency region [28], suggesting better capacitive

In this work, ACs with high performance for supercapacitor are prepared from coconut shell by KOH activation, using a novel oxidation‐ultrasound process to deeply remove the K+ of AC. The experiment results demonstrate that the AC samples washed with 1.0 wt% HCl solution for 120 min and subsequently treated with 0.6 wt% H2O2 solution at 60°C in

AC‐W and AC‐W‐H1%/120min, the AC‐W‐H1%/120min‐U0.6%/60°C/8h exhibits a larger surface area and

application, the AC‐W‐H1%/120min‐U0.6%/60°C/8h showed a high specific capacitance of 306 F/g with a coulombic efficiency of 96% after 3000 cycles. The oxidation‐ultrasound process has a great potential to produce ACs with high performance for supercapacitor applications.

