**Meet the editors**

Rosalinda Inguanta is currently an assistant professor of Applied Physical Chemistry at the University of Palermo. She received her PhD degree in 2008 at the University of Palermo. Her research interests include the synthesis of nanostructured electrochemical devices for electrochemical sensors, batteries, solar cells, and electrolysers and the synthesis of biocoatings to prevent

the corrosion of medical devices in human body and the recovery of precious metals from electronic waste. She has developed a new method for the synthesis of nanostructures and thin film of different materials based on galvanic synthesis. In the field of semiconductors, she has studied the electrochemical deposition of different semiconductors, such as CIGS and CZTS; currently, her research team is employed on the study of perovskite thin film starting from electrodeposition of PbO2 layers.

Carmelo Sunseri is a full professor of Applied Thermodynamic at the University of Palermo. His main research interests are in the fields of thermodynamic and kinetic study of electrochemical reactions on both metallic and semiconducting electrodes, electrochemical synthesis and characterization of nanostructured materials, and photoelectrochemical investigation of semiconductor

surfaces. His research activity is focused on materials and systems of technological interest in view of their potential application in the following fields: electrochemical storage and conversion of energy (battery and fuel cells), photovoltaics, fabrication of thin-film devices, and electrochemical fabrication of porous devices.

Contents

**Preface VII**

**Metamaterials 3**

Victor Barsan

Anatoliy Opanasyuk

Massimo Innocenti

Joshi

**Semiconductor Thin Films 109**

**Section 1 General Methods for Semiconductors Growth 1**

Abhay Singh and Usha Philipose

**Conditions and Janus Nanorods 47**

**Section 2 Semiconductors for Energy Applications 67**

Chapter 1 **Investigation of the Nanostructured Semiconductor**

Aleksej Trofimov, Tatjana Gric and Ortwin Hess

Chapter 2 **Understanding the Mechanisms that Affect the Quality of**

Chapter 4 **Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method 69**

Chapter 5 **E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals 89**

**Electrochemically Grown Semiconducting Nanowires 25**

Oleksandr Dobrozhan, Denys Kurbatov, Petro Danilchenko and

Emanuele Salvietti, Andrea Giaccherini, Filippo Gambinossi, Maria Luisa Foresti, Maurizio Passaponti, Francesco Di Benedetto and

Sreekanth Mandati, Bulusu V. Sarada, Suhash R. Dey and Shrikant V.

Chapter 6 **Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2**

Chapter 3 **Semiconductor Quantum Wells with BenDaniel-Duke Boundary**

## Contents

#### **Preface XI**


#### Chapter 7 **The Electrochemical Performance of Deposited Manganese Oxide-Based Film as Electrode Material for Electrochemical Capacitor Application 133** Chan Pei Yi and Siti Rohana Majid

Preface

to any desired shape.

The discovery of semiconducting materials allowed for tremendous and important advance‐ ments in the field of photonics, photovoltaics, electronics, and thermoelectrics. Semiconduc‐ tors play a key role for the miniaturization of computers and manufacturing of computer

Over the years, many efforts have been devoted to increase the knowledge on the numerous and fascinating properties of semiconducting materials, not only on silicon and germanium, certainly the most famous and used semiconductors, but also on multielement semiconduc‐ tors such as chalcopyrite and kesterite. Furthermore, semiconducting plastics have attracted great attention in plastic light-emitting diodes (LEDs), which are flexible and can be molded

This book deals with the synthesis and characterization of different semiconducting materi‐ als with special attention to the possible technological applications. Fundamental issues are

Different deposition methods are discussed in this book, such as chemical bath deposition and spray pyrolysis for fabricating thin film and electrospinning process for obtaining semi‐ conducting nanofibers. Particular attention has been devoted to the use of electrochemical methods for synthesizing semiconducting materials of interest for the hi-tech industry. To make semiconductor-based devices more attractive, low-cost and easily scalable processes must be developed. In this context, electrodeposition is a valuable tool because it is cheap, simple, and quick in comparison to other fabrication methods. Good-quality large-area thin films and nanostructures can be obtained by this technique. In particular, template-based electrochemical synthesis is of advantage for fabricating semiconductors with nanostruc‐ tured morphology. A further advantage of the electrochemical route is the possibility to si‐ multaneously control both electrical conduction type and energy gap of the semiconductors by adjusting the deposition parameters. In the case of multielement semiconductors, it is also possible to control the composition of deposit by adjusting the nature and composition of the deposition bath. Apart from conventional electrochemical methods such as direct elec‐ troplating and pulse or pulse-reverse electroplating, electrochemical atomic layer deposition

> **Rosalinda Inguanta and Carmelo Sunseri** Applied Physical Chemistry Laboratory

> > University of Palermo

Palermo, Italy

Department of Industrial and Digital Innovation —

Chemical, Computer, Management, Mechanical Engineering

parts. They are the beating heart of diodes, transistors, and photovoltaic cells.

also treated showing new developments about semiconductor physics.

(E-ALD) of metal chalcogenides is also discussed.

#### Chapter 8 **Semiconducting Electrospun Nanofibers for Energy Conversion 159**

Giulia Massaglia and Marzia Quaglio

## Preface

Chapter 7 **The Electrochemical Performance of Deposited Manganese**

Chapter 8 **Semiconducting Electrospun Nanofibers for Energy**

Giulia Massaglia and Marzia Quaglio

**Capacitor Application 133** Chan Pei Yi and Siti Rohana Majid

**Conversion 159**

**VI** Contents

**Oxide-Based Film as Electrode Material for Electrochemical**

The discovery of semiconducting materials allowed for tremendous and important advance‐ ments in the field of photonics, photovoltaics, electronics, and thermoelectrics. Semiconduc‐ tors play a key role for the miniaturization of computers and manufacturing of computer parts. They are the beating heart of diodes, transistors, and photovoltaic cells.

Over the years, many efforts have been devoted to increase the knowledge on the numerous and fascinating properties of semiconducting materials, not only on silicon and germanium, certainly the most famous and used semiconductors, but also on multielement semiconduc‐ tors such as chalcopyrite and kesterite. Furthermore, semiconducting plastics have attracted great attention in plastic light-emitting diodes (LEDs), which are flexible and can be molded to any desired shape.

This book deals with the synthesis and characterization of different semiconducting materi‐ als with special attention to the possible technological applications. Fundamental issues are also treated showing new developments about semiconductor physics.

Different deposition methods are discussed in this book, such as chemical bath deposition and spray pyrolysis for fabricating thin film and electrospinning process for obtaining semi‐ conducting nanofibers. Particular attention has been devoted to the use of electrochemical methods for synthesizing semiconducting materials of interest for the hi-tech industry. To make semiconductor-based devices more attractive, low-cost and easily scalable processes must be developed. In this context, electrodeposition is a valuable tool because it is cheap, simple, and quick in comparison to other fabrication methods. Good-quality large-area thin films and nanostructures can be obtained by this technique. In particular, template-based electrochemical synthesis is of advantage for fabricating semiconductors with nanostruc‐ tured morphology. A further advantage of the electrochemical route is the possibility to si‐ multaneously control both electrical conduction type and energy gap of the semiconductors by adjusting the deposition parameters. In the case of multielement semiconductors, it is also possible to control the composition of deposit by adjusting the nature and composition of the deposition bath. Apart from conventional electrochemical methods such as direct elec‐ troplating and pulse or pulse-reverse electroplating, electrochemical atomic layer deposition (E-ALD) of metal chalcogenides is also discussed.

#### **Rosalinda Inguanta and Carmelo Sunseri**

Applied Physical Chemistry Laboratory Department of Industrial and Digital Innovation — Chemical, Computer, Management, Mechanical Engineering University of Palermo Palermo, Italy

**Section 1**

**General Methods for Semiconductors Growth**

**General Methods for Semiconductors Growth**

**Chapter 1**

Provisional chapter

**Investigation of the Nanostructured Semiconductor**

DOI: 10.5772/intechopen.72801

The presence of electromagnetic waves on two-dimensional interfaces has been extensively studied over the last several decades. Surface plasmonic polariton (SPP), which normally exists at the interface between a noble metal and a dielectric, is treated as the most widely investigated surface wave. SPPs have promoted new applications in many fields such as microelectronics, photovoltaics, etc. Recently, it has been shown that by nanostructuring the metal surface, it is possible to modify the dispersion of SPPs in a prescribed manner. Herein, we demonstrate the existence of a new kind of surface wave between two anisotropic meta-materials. In contrast to extensively studied surface waves such as SPPs and Dyakonov waves, the surface waves supported by the nanostructured semiconductor metamaterial cross the light line, and a substantial portion at lower frequencies lies above the free-space light line. Consequently, the proposed structure will

Plasmonics and the recent birth of metamaterials (MMs) [1–4] and transformation optics [5, 6] are currently opening a gateway to the development of a family of novel devices with unprecedented functionalities ranging from sub-wavelength plasmonic waveguides and optical nanoresonators [7] to superlenses, hyperlenses [8] and light concentrators [9]. Coupling between photons and surface plasmon polaritons (SPPs) [10, 11] is enabled by the periodically nanostructured metallic films allowing for exceptional and tunable optical properties determined by a combination of design geometry, the surrounding dielectric permittivity and the choice of metal [12, 13]. SPPs have promoted new applications in many fields such as

> © 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 eproduction in any medium, provided the original work is properly cited.

© 2018 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.

Investigation of the Nanostructured Semiconductor

Aleksej Trofimov, Tatjana Gric and Ortwin Hess

Aleksej Trofimov, Tatjana Gric and Ortwin Hess

Additional information is available at the end of the chapter

interact with the material via leaky waves.

Keywords: metamaterials, semiconductor, surface plasmon polaritons

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.72801

**Metamaterials**

Abstract

1. Introduction

Metamaterials

#### **Investigation of the Nanostructured Semiconductor Metamaterials** Investigation of the Nanostructured Semiconductor Metamaterials

DOI: 10.5772/intechopen.72801

Aleksej Trofimov, Tatjana Gric and Ortwin Hess Aleksej Trofimov, Tatjana Gric and Ortwin Hess

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/intechopen.72801

#### Abstract

The presence of electromagnetic waves on two-dimensional interfaces has been extensively studied over the last several decades. Surface plasmonic polariton (SPP), which normally exists at the interface between a noble metal and a dielectric, is treated as the most widely investigated surface wave. SPPs have promoted new applications in many fields such as microelectronics, photovoltaics, etc. Recently, it has been shown that by nanostructuring the metal surface, it is possible to modify the dispersion of SPPs in a prescribed manner. Herein, we demonstrate the existence of a new kind of surface wave between two anisotropic meta-materials. In contrast to extensively studied surface waves such as SPPs and Dyakonov waves, the surface waves supported by the nanostructured semiconductor metamaterial cross the light line, and a substantial portion at lower frequencies lies above the free-space light line. Consequently, the proposed structure will interact with the material via leaky waves.

Keywords: metamaterials, semiconductor, surface plasmon polaritons

#### 1. Introduction

Plasmonics and the recent birth of metamaterials (MMs) [1–4] and transformation optics [5, 6] are currently opening a gateway to the development of a family of novel devices with unprecedented functionalities ranging from sub-wavelength plasmonic waveguides and optical nanoresonators [7] to superlenses, hyperlenses [8] and light concentrators [9]. Coupling between photons and surface plasmon polaritons (SPPs) [10, 11] is enabled by the periodically nanostructured metallic films allowing for exceptional and tunable optical properties determined by a combination of design geometry, the surrounding dielectric permittivity and the choice of metal [12, 13]. SPPs have promoted new applications in many fields such as

Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and eproduction in any medium, provided the original work is properly cited. © 2018 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

microelectronics [14], photovoltaics [15], near-field sensing [16], laser technology [17, 18], photonics [19], meta-materials design [2], high-order harmonics generation [20] or charged particle acceleration [21]. Recently, it has been shown that by nanostructuring the metal surface, it is possible to modify the dispersion of SPPs or excite the SPPs in a prescribed manner [22, 23].

The process of replacement of the uniaxial medium by a biaxial crystal [24], an indefinite medium [25] and a structurally chiral material [26] may enforce the presence of hybrid surface waves with some parallel characteristics. In the latter case, a methodology developed by Tamm [27] was adopted seeking to find a new type of surface wave, called as Dyakonov-Tamm wave, as it combines the features of Dyakonov surface waves (DSWs) and Tamm states. The use of structured materials with extreme anisotropy provided a fertile background aiming to increase the range of directions of DSWs substantially, as it is compared with the rather narrow range observed with natural birefringent materials [28]. Especially, outstanding results take place if the metallic nanoelements are employed to the anisotropic structures, as it occurs, for example, with a simple metal-dielectric multilayer, a case where the angular range may surpass half of a right angle [29]. The propagation length of these DSWs is drastically limited by the penetration depth inside the lossy MM [30] as it is caused by the specific damping capacity of metals.

The examination of two different interfaces, i.e. MM/dielectric and MM/TCO, is of the particular importance. Surface waves of different kinds, including DSWs along with traditional-like SPPs, are examined. Contrarily, the introduction of MM/TCO interface leads to a transformation of the traditional-like SPPs. As a consequence, the new types of surface waves are found.

Consequently, an example to estimate the limitation on the structure period under the effectivemedium theory (EMT) is considered. Proposing model, when the wavelength of radiation is much larger than the thickness of any layer, one can apply the effective-medium approach based

Figure 1. One-dimensional MM composed of alternating TCO and dielectric layers. The MM/dielectric (ε<sup>R</sup> ¼ 1 or

First, the dispersion features of HSPPs are investigated. On the contrary to the approach presented in [36–38], the damping term in the TCO is not ignored in the process of analysing and calculating their dispersion properties. It is worth to mention that this particular MM is equivalent to a uniaxial-anisotropy effective medium, with its anisotropy axis (the optical axis) along the structure periodicity in the long-wavelength limit. Its effective permittivity tensor is

> ε<sup>k</sup> 0 0 0 ε<sup>k</sup> 0 0 0 ε<sup>⊥</sup>

in the principal-axis coordinate system. The principal values of the tensor are expressed with

where f TCO ¼ dTCO=L and f <sup>d</sup> ¼ dd=L, L ¼ dd þ dTCO represent the TCO and dielectric filling ratios, respectively. It should be realized that MMs with a very large permittivity or a nearzero permittivity exhibit interesting properties [40]. While one may consider MMs with a very large permittivity as optical conductors, those with a near-zero permittivity can be used as optical insulators [40]. The zero point and divergence point of the principal values were also of

1

CA, (1)

ε<sup>k</sup> ¼ f TCOεTCO þ f <sup>d</sup>ε<sup>d</sup> (2)

Investigation of the Nanostructured Semiconductor Metamaterials

http://dx.doi.org/10.5772/intechopen.72801

5

<sup>ε</sup><sup>⊥</sup> <sup>¼</sup> <sup>ε</sup>TCOεd<sup>=</sup> <sup>f</sup> TCO<sup>ε</sup> <sup>þ</sup> <sup>f</sup> <sup>d</sup>εTCO � �, (3)

ε ¼ ε<sup>0</sup>

0 B@

on averaging the structure parameters.

ε<sup>R</sup> ¼ 2:25) and MM/TCO interfaces are considered.

written as

[39, 40]

Moreover, hyperbolic metamaterials, being special kind of anisotropic metamaterial with dielectric tenor elements having the mixed signs, have attracted growing attention due to their ability to support very large wave vectors. Their exotic features give rise to many intriguing applications, such as sub-wavelength imaging [31, 32] and hyper-lens [33, 34] that are infeasible with natural materials. In this paper, we demonstrate the existence of a new kind of surface wave between two anisotropic metamaterials. In contrast to extensively studied surface waves such as SPPs and Dyakonov waves, whose in-plane wave vector is greater than that of the bulk modes, the surface waves supported by the nanostructured semiconductor metamaterial cross the light line, and a substantial portion at lower frequencies lies above the free-space light line, which typically separates non-radiative (bound) and radiative (leaky) regions.

## 2. Transparent conducting oxide (TCO)—dielectric composite heterostructure-based multilayer metamaterial

The structure of the metamaterial (MM) is shown in Figure 1, where dTCO and dd represent the thicknesses of TCO and dielectric layers, respectively. All the involved media are nonmagnetic, so the magnetic permeability of every medium is the same as that of vacuum. In our numerical calculation, we use a semi-infinite TCO/PbS MM as an example to explore the dispersive features of hyperbolic surface plasmon polaritons (HSPPs), where ε<sup>d</sup> = 18.8 for PbS layers, and εTCO is calculated using the parameters presented in [35].

microelectronics [14], photovoltaics [15], near-field sensing [16], laser technology [17, 18], photonics [19], meta-materials design [2], high-order harmonics generation [20] or charged particle acceleration [21]. Recently, it has been shown that by nanostructuring the metal surface, it is possible to modify the dispersion of SPPs or excite the SPPs in a prescribed

The process of replacement of the uniaxial medium by a biaxial crystal [24], an indefinite medium [25] and a structurally chiral material [26] may enforce the presence of hybrid surface waves with some parallel characteristics. In the latter case, a methodology developed by Tamm [27] was adopted seeking to find a new type of surface wave, called as Dyakonov-Tamm wave, as it combines the features of Dyakonov surface waves (DSWs) and Tamm states. The use of structured materials with extreme anisotropy provided a fertile background aiming to increase the range of directions of DSWs substantially, as it is compared with the rather narrow range observed with natural birefringent materials [28]. Especially, outstanding results take place if the metallic nanoelements are employed to the anisotropic structures, as it occurs, for example, with a simple metal-dielectric multilayer, a case where the angular range may surpass half of a right angle [29]. The propagation length of these DSWs is drastically limited by the penetration depth inside the lossy MM [30] as it is caused by the specific damping capacity of metals.

The examination of two different interfaces, i.e. MM/dielectric and MM/TCO, is of the particular importance. Surface waves of different kinds, including DSWs along with traditional-like SPPs, are examined. Contrarily, the introduction of MM/TCO interface leads to a transformation of the traditional-like SPPs. As a consequence, the new types of surface waves are found. Moreover, hyperbolic metamaterials, being special kind of anisotropic metamaterial with dielectric tenor elements having the mixed signs, have attracted growing attention due to their ability to support very large wave vectors. Their exotic features give rise to many intriguing applications, such as sub-wavelength imaging [31, 32] and hyper-lens [33, 34] that are infeasible with natural materials. In this paper, we demonstrate the existence of a new kind of surface wave between two anisotropic metamaterials. In contrast to extensively studied surface waves such as SPPs and Dyakonov waves, whose in-plane wave vector is greater than that of the bulk modes, the surface waves supported by the nanostructured semiconductor metamaterial cross the light line, and a substantial portion at lower frequencies lies above the free-space light line,

which typically separates non-radiative (bound) and radiative (leaky) regions.

2. Transparent conducting oxide (TCO)—dielectric composite

The structure of the metamaterial (MM) is shown in Figure 1, where dTCO and dd represent the thicknesses of TCO and dielectric layers, respectively. All the involved media are nonmagnetic, so the magnetic permeability of every medium is the same as that of vacuum. In our numerical calculation, we use a semi-infinite TCO/PbS MM as an example to explore the dispersive features of hyperbolic surface plasmon polaritons (HSPPs), where ε<sup>d</sup> = 18.8 for PbS layers,

heterostructure-based multilayer metamaterial

and εTCO is calculated using the parameters presented in [35].

manner [22, 23].

4 Semiconductors - Growth and Characterization

Figure 1. One-dimensional MM composed of alternating TCO and dielectric layers. The MM/dielectric (ε<sup>R</sup> ¼ 1 or ε<sup>R</sup> ¼ 2:25) and MM/TCO interfaces are considered.

Consequently, an example to estimate the limitation on the structure period under the effectivemedium theory (EMT) is considered. Proposing model, when the wavelength of radiation is much larger than the thickness of any layer, one can apply the effective-medium approach based on averaging the structure parameters.

First, the dispersion features of HSPPs are investigated. On the contrary to the approach presented in [36–38], the damping term in the TCO is not ignored in the process of analysing and calculating their dispersion properties. It is worth to mention that this particular MM is equivalent to a uniaxial-anisotropy effective medium, with its anisotropy axis (the optical axis) along the structure periodicity in the long-wavelength limit. Its effective permittivity tensor is written as

$$
\varepsilon = \varepsilon\_0 \begin{pmatrix}
\varepsilon\_{\parallel} & 0 & 0 \\
0 & \varepsilon\_{\parallel} & 0 \\
0 & 0 & \varepsilon\_{\perp}
\end{pmatrix},
\tag{1}
$$

in the principal-axis coordinate system. The principal values of the tensor are expressed with [39, 40]

$$
\varepsilon\_{\parallel} = f\_{\text{TCO}} \varepsilon\_{\text{TCO}} + f\_d \varepsilon\_d \tag{2}
$$

$$
\varepsilon\_{\perp} = \varepsilon\_{\text{TCO}} \varepsilon\_d / \left( f\_{\text{TCO}} \varepsilon + f\_d \varepsilon\_{\text{TCO}} \right) \tag{3}
$$

where f TCO ¼ dTCO=L and f <sup>d</sup> ¼ dd=L, L ¼ dd þ dTCO represent the TCO and dielectric filling ratios, respectively. It should be realized that MMs with a very large permittivity or a nearzero permittivity exhibit interesting properties [40]. While one may consider MMs with a very large permittivity as optical conductors, those with a near-zero permittivity can be used as optical insulators [40]. The zero point and divergence point of the principal values were also of

Figure 2. The principal values of effective permittivity and the different frequency regions for typical TCO-filling ratios corresponding to the MM/dielectric interface. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d).

the particular interest to discuss the dispersion features of SPPs [41]. The zero point and divergence point of ε<sup>⊥</sup> or ε<sup>k</sup> will be applied to discuss HSPPs.

Electric and magnetic fields' tangential components need to be evaluated at the interface in order to get metamaterial interface confined surface mode unique dispersion [42]:

$$\beta = k \sqrt{\frac{\left(\varepsilon\_{\parallel}^{R} - \varepsilon\_{\parallel}^{L}\right) \varepsilon\_{\perp}^{R} \varepsilon\_{\perp}^{L}}{\varepsilon\_{\perp}^{R} \varepsilon\_{\parallel}^{R} - \varepsilon\_{\perp}^{L} \varepsilon\_{\parallel}^{L}}},\tag{4}$$

β ¼ k

where <sup>a</sup> <sup>¼</sup> ddε<sup>d</sup> � ddεTCOf TCO

right-hand side of the interface.

0

BBBB@

<sup>f</sup> TCO�<sup>1</sup> , <sup>b</sup> <sup>¼</sup> ddεTCO � ddεdf TCO

<sup>ε</sup>TCOεdε<sup>R</sup> dd � ddf TCO

hand side is the same as employed in the MM. This dispersion relation reads

f TCO�1 � � <sup>ε</sup>dR � <sup>a</sup>

Figure 3. The principal values of effective permittivity and the different frequency regions for typical TCO-filling ratios, corresponding to the MM/TCO interface. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d).

> <sup>R</sup> � <sup>ε</sup>TCOεda b � �b

It is of particular interest to obtain the dispersion relation for the interface states in the effective media approach corresponding to the MM interface, having in mind that material at the right-

> <sup>β</sup> <sup>¼</sup> <sup>k</sup> <sup>ε</sup>TCOε<sup>d</sup> εTCO þ ε<sup>d</sup> � �<sup>1</sup>=<sup>2</sup>

ε2

dd�ddf TCO f TCO�1 1

Investigation of the Nanostructured Semiconductor Metamaterials

http://dx.doi.org/10.5772/intechopen.72801

7

1=2

, (5)

(6)

CCCCA

<sup>f</sup> TCO�<sup>1</sup> , <sup>ε</sup><sup>R</sup> is the permittivity of the material at the

!

where k is the absolute value of wave vector in vacuum and β is the component of the wave vector parallel to the interface.

It is interesting to notice that in the case of the MM interface, the obtained result for the dispersion is as follows:

Investigation of the Nanostructured Semiconductor Metamaterials http://dx.doi.org/10.5772/intechopen.72801 7

Figure 3. The principal values of effective permittivity and the different frequency regions for typical TCO-filling ratios, corresponding to the MM/TCO interface. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d).

the particular interest to discuss the dispersion features of SPPs [41]. The zero point and

Figure 2. The principal values of effective permittivity and the different frequency regions for typical TCO-filling ratios corresponding to the MM/dielectric interface. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d).

Electric and magnetic fields' tangential components need to be evaluated at the interface in

εR <sup>k</sup> � <sup>ε</sup><sup>L</sup> k � �

> εR ⊥ε<sup>R</sup> <sup>k</sup> � <sup>ε</sup><sup>L</sup> ⊥ε<sup>L</sup> k

where k is the absolute value of wave vector in vacuum and β is the component of the wave

It is interesting to notice that in the case of the MM interface, the obtained result for the

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

εR ⊥ε<sup>L</sup> ⊥

vuuut , (4)

order to get metamaterial interface confined surface mode unique dispersion [42]:

β ¼ k

divergence point of ε<sup>⊥</sup> or ε<sup>k</sup> will be applied to discuss HSPPs.

vector parallel to the interface.

6 Semiconductors - Growth and Characterization

dispersion is as follows:

$$\beta = k \left( \frac{\varepsilon\_{\rm TCO} \varepsilon\_d \varepsilon\_R \left( d\_d - \frac{d\_d f\_{\rm CO}}{f\_{\rm CO} - 1} \right) \left( \varepsilon\_{\rm dR} - \frac{d}{d\_d - \frac{d\_d f\_{\rm CO}}{f\_{\rm CO} - 1}} \right)}{\left( \varepsilon\_R^2 - \frac{\varepsilon\_{\rm TCO} \varepsilon\_d \varepsilon\_0}{b} \right) b} \right)^{1/2} \,, \tag{5}$$

where <sup>a</sup> <sup>¼</sup> ddε<sup>d</sup> � ddεTCOf TCO <sup>f</sup> TCO�<sup>1</sup> , <sup>b</sup> <sup>¼</sup> ddεTCO � ddεdf TCO <sup>f</sup> TCO�<sup>1</sup> , <sup>ε</sup><sup>R</sup> is the permittivity of the material at the right-hand side of the interface.

It is of particular interest to obtain the dispersion relation for the interface states in the effective media approach corresponding to the MM interface, having in mind that material at the righthand side is the same as employed in the MM. This dispersion relation reads

$$\beta = k \left( \frac{\varepsilon\_{TCO} \varepsilon\_d}{\varepsilon\_{TCO} + \varepsilon\_d} \right)^{1/2} \tag{6}$$

Figure 4. Dispersion curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d), corresponding to the MM/dielectric interface.

As a matter of fact, we obtain a surprising result: the dispersion of a (single) interface mode does not depend on the thicknesses of the layers, and it coincides with the dispersion of a conventional surface plasmon at metal-dielectric interface.

among ε⊥, ε<sup>k</sup> and ε<sup>R</sup> as shown in Figures 2 and 3. ε<sup>⊥</sup> > ε<sup>R</sup> > ε<sup>k</sup> and ε<sup>k</sup> < 0 in the cyan region, ε<sup>⊥</sup> > ε<sup>R</sup> > ε<sup>k</sup> in the green region, ε<sup>k</sup> > ε<sup>R</sup> > ε<sup>⊥</sup> and ε<sup>⊥</sup> < 0 in the grey region, ε<sup>R</sup> > ε<sup>k</sup> > ε<sup>⊥</sup> and ε<sup>⊥</sup> < 0 in the magenta region, and ε<sup>⊥</sup> < ε<sup>k</sup> < 0 in the orange region. The tunability of the effective parameters presented in Figures 2 and 3 has an effect on the dispersion curves of the HSPPs. As shown in Figures 2 and 3, the parameter that principally defines the tunability of the effective optical parameters of our metamaterial is the TCO-filling ratio, fTCO. The disper-

Figure 5. Dispersion curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in

Investigation of the Nanostructured Semiconductor Metamaterials

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The curve colors correspond to different frequency regions. For fTCO = 0.3, there are three HSPPs belonging to three different kinds, respectively, in the case of MM/dielectric interface and two kinds of HSPPs in the case of MM/TCO interface. The upper waves exist in the gray region for MM/TCO case. Three exceptional cases related to fTCO = 0.5, 0.7 and 0.9. In the case of fTCO = 0.5, two kinds of HSPPs exist in two color regions, i.e. cyan and grey for MM/dielectric and MM/TCO interfaces. The upper short curves in the case of fTCO = 0.7 lie in the gray region,

sion curves are illustrated in Figures 4 and 5 for various TCO-filling ratios.

(c) and f TCO ¼ 0:9 in (d), corresponding to the MM/dielectric interface.

#### 2.1. The mode structure

In the case of a spatially infinite anisotropic material, invariant in two directions, the electromagnetic wave dispersion can be plotted for both MM/dielectric and MM/TCO cases. Thus, herein, we present analysis performed after the homogenization of the MM corresponding to the MM/dielectric and MM/TCO interfaces. Doing so, in the numerical calculations, the semiinfinite AZO/PbS MM is taken as an example. We will first review the optical properties by depicting the curves of the principal values (ε<sup>⊥</sup> and εk) and the dielectric constant (εR) and then by distinguishing different frequency regions according to the properties of and relation

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Figure 5. Dispersion curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d), corresponding to the MM/dielectric interface.

among ε⊥, ε<sup>k</sup> and ε<sup>R</sup> as shown in Figures 2 and 3. ε<sup>⊥</sup> > ε<sup>R</sup> > ε<sup>k</sup> and ε<sup>k</sup> < 0 in the cyan region, ε<sup>⊥</sup> > ε<sup>R</sup> > ε<sup>k</sup> in the green region, ε<sup>k</sup> > ε<sup>R</sup> > ε<sup>⊥</sup> and ε<sup>⊥</sup> < 0 in the grey region, ε<sup>R</sup> > ε<sup>k</sup> > ε<sup>⊥</sup> and ε<sup>⊥</sup> < 0 in the magenta region, and ε<sup>⊥</sup> < ε<sup>k</sup> < 0 in the orange region. The tunability of the effective parameters presented in Figures 2 and 3 has an effect on the dispersion curves of the HSPPs. As shown in Figures 2 and 3, the parameter that principally defines the tunability of the effective optical parameters of our metamaterial is the TCO-filling ratio, fTCO. The dispersion curves are illustrated in Figures 4 and 5 for various TCO-filling ratios.

As a matter of fact, we obtain a surprising result: the dispersion of a (single) interface mode does not depend on the thicknesses of the layers, and it coincides with the dispersion of a

Figure 4. Dispersion curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in

In the case of a spatially infinite anisotropic material, invariant in two directions, the electromagnetic wave dispersion can be plotted for both MM/dielectric and MM/TCO cases. Thus, herein, we present analysis performed after the homogenization of the MM corresponding to the MM/dielectric and MM/TCO interfaces. Doing so, in the numerical calculations, the semiinfinite AZO/PbS MM is taken as an example. We will first review the optical properties by depicting the curves of the principal values (ε<sup>⊥</sup> and εk) and the dielectric constant (εR) and then by distinguishing different frequency regions according to the properties of and relation

conventional surface plasmon at metal-dielectric interface.

(c) and f TCO ¼ 0:9 in (d), corresponding to the MM/dielectric interface.

2.1. The mode structure

8 Semiconductors - Growth and Characterization

The curve colors correspond to different frequency regions. For fTCO = 0.3, there are three HSPPs belonging to three different kinds, respectively, in the case of MM/dielectric interface and two kinds of HSPPs in the case of MM/TCO interface. The upper waves exist in the gray region for MM/TCO case. Three exceptional cases related to fTCO = 0.5, 0.7 and 0.9. In the case of fTCO = 0.5, two kinds of HSPPs exist in two color regions, i.e. cyan and grey for MM/dielectric and MM/TCO interfaces. The upper short curves in the case of fTCO = 0.7 lie in the gray region,

Figure 6. Absorption curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d), corresponding to the MM/dielectric interface.

(Figures 2 and 4). The former extension is demonstrated by the dark-green color in Figure 2. It is worthwhile mentioning that extension of the gray region in Figure 3 is possible due to the employment of ITO instead of GZO at the right-hand side of the interface (the extension of the gray range is depicted by the light-gray color). Due to the negative principal values of the effective permittivity in the orange region, the HSPP in this region is similar to the traditional SPP, so we should name it the traditional-like SPP. The others are new types of HSPPs. Thus, all the HSPPs are divided into five kinds, situated in the five color regions in Figures 4 and 5, respectively. In the case of AZO/PbS MM and air/dielectric interface, five types of the HSPPs

Figure 7. Absorption curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in

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As expected, the traditional-like SPP has a typical dispersion in the case of the MM/dielectric interface, lying to the right of the light line. At the same time, the degradation of the dispersion in the orange region takes place in the case of the MM/TCO interface. The dispersion properties can be tuned with the TCO-filling ratio of the MM realization. While for low TCO-filling ratio in the

exist, and only three are seen in the case of AZO/PbS MM and TCO interface.

(c) and f TCO ¼ 0:9 in (d), corresponding to the MM/TCO interface.

as illustrated in Figures 4 and 5. The other HSPPs lie in the orange region. The additional HSSPs lying in the magenta region correspond to the case of MM/dielectric interface. The extension of the magenta range for the case ε<sup>R</sup> ¼ 2:25 is displayed by the dark-magenta color. In the case of fTCO = 0.9, two kinds of HSPPs are found lying in the orange and cyan regions for MM/TCO case and three kinds of HSPPS lying in the cyan, orange and gray regions for MM/ dielectric case. Figure 4 also demonstrates that there always is one HSPP in the cyan region for various TCO-filling ratios in MM/dielectric case. It is worthwhile mentioning that the case ε<sup>R</sup> ¼ εITO also allows for the rich phenomenon as the HSPP always exists in the cyan region for various filling ratios (Figure 5).

Based on the necessary condition for the existence of the DSW [39], the HSPP in the green region (Figure 4) is similar to the DSW so that it should be called Dyakonov-like SPP [2] or the Dyakonov defined in [39]. Moreover, the frequency range of the DSW existence can be extended by replacing the material at the right-hand side of the interface with ε<sup>R</sup> ¼ 2:25

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Figure 7. Absorption curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in (c) and f TCO ¼ 0:9 in (d), corresponding to the MM/TCO interface.

as illustrated in Figures 4 and 5. The other HSPPs lie in the orange region. The additional HSSPs lying in the magenta region correspond to the case of MM/dielectric interface. The extension of the magenta range for the case ε<sup>R</sup> ¼ 2:25 is displayed by the dark-magenta color. In the case of fTCO = 0.9, two kinds of HSPPs are found lying in the orange and cyan regions for MM/TCO case and three kinds of HSPPS lying in the cyan, orange and gray regions for MM/ dielectric case. Figure 4 also demonstrates that there always is one HSPP in the cyan region for various TCO-filling ratios in MM/dielectric case. It is worthwhile mentioning that the case ε<sup>R</sup> ¼ εITO also allows for the rich phenomenon as the HSPP always exists in the cyan region

Figure 6. Absorption curves of HSPPs for various TCO-filling ratios, i.e. f TCO ¼ 0:3 in (a), f TCO ¼ 0:5 in (b), f TCO ¼ 0:7 in

Based on the necessary condition for the existence of the DSW [39], the HSPP in the green region (Figure 4) is similar to the DSW so that it should be called Dyakonov-like SPP [2] or the Dyakonov defined in [39]. Moreover, the frequency range of the DSW existence can be extended by replacing the material at the right-hand side of the interface with ε<sup>R</sup> ¼ 2:25

for various filling ratios (Figure 5).

10 Semiconductors - Growth and Characterization

(c) and f TCO ¼ 0:9 in (d), corresponding to the MM/dielectric interface.

(Figures 2 and 4). The former extension is demonstrated by the dark-green color in Figure 2. It is worthwhile mentioning that extension of the gray region in Figure 3 is possible due to the employment of ITO instead of GZO at the right-hand side of the interface (the extension of the gray range is depicted by the light-gray color). Due to the negative principal values of the effective permittivity in the orange region, the HSPP in this region is similar to the traditional SPP, so we should name it the traditional-like SPP. The others are new types of HSPPs. Thus, all the HSPPs are divided into five kinds, situated in the five color regions in Figures 4 and 5, respectively. In the case of AZO/PbS MM and air/dielectric interface, five types of the HSPPs exist, and only three are seen in the case of AZO/PbS MM and TCO interface.

As expected, the traditional-like SPP has a typical dispersion in the case of the MM/dielectric interface, lying to the right of the light line. At the same time, the degradation of the dispersion in the orange region takes place in the case of the MM/TCO interface. The dispersion properties can be tuned with the TCO-filling ratio of the MM realization. While for low TCO-filling ratio in the case of the MM/TCO interface, two types of the modes are always present, for higher TCO-filling ratio, the disappearance of modes in the gray region takes place (Figure 5(d)).

The complex mode structure (Figures 4 and 5) corresponding to either MM/dielectric or MM/ TCO interface emerges as a consequence of the confinement of plasmon polaritons in the direction perpendicular to the wave propagation. These electromagnetic surface waves arise via the coupling of the electromagnetic fields to oscillations of the conductor's electron plasma.

The imaginary parts of the wave vector (i.e. absorption) are plotted in Figures 6 and 7. It should be mentioned that negative values of the absorption in Figures 6 and 7 result from non-physical solutions of the dispersion equation and have been omitted in line with [43]. Taking advantage of the absorption resonances, one can show that the simple multilayer structures without possessing any periodic corrugation have the prospective to act as directive and monochromatic thermal sources [44].

#### 3. Nanostructured semiconductor metamaterial

Another interesting MM structure depicted in Figure 8 is periodic stack of semiconductordielectric layers called hyperbolic metamaterial heterostructure.

The effective permittivity of the semiconductor (Si) can be calculated as follows:

$$
\varepsilon\_{1,3}(\omega) = \varepsilon\_{\circ} - \frac{\omega\_p^2}{\omega^2 + i\delta\omega'} \tag{7}
$$

than the thickness of any layer is applied aiming to describe the optical response of such a system. The dispersion relation for the surface modes localized at the boundary separating two anisotropic media [42] is found by applying the appropriate boundary conditions, i.e. matching the tangential components of the electrical and magnetic fields at the interface. It is

Figure 9. The influence of (a) doping level N and (b) thickness of dielectric dd, on the real part of ε⊥. dd = 10 nm in (a) and

metallic properties at terahertz frequencies [46, 47] and has the potential to replace metals in such applications [48]. The case of a heavy-doped Si is considered, assuming that the doping

It should be mentioned that dramatic control of the frequency range of the surface wave existence is mostly concerned with the modifications of the permittivities and thicknesses of the layers [50] employed in the HMMs. To further study the surface waves, the tangential components of the electric and magnetic fields at the interface should be evaluated, and a single surface mode with the propagation constant should be obtained aiming to get the unique dispersion

relation for the surface modes confined at the interface between two metamaterials [42].

) has been shown to exhibit

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interesting to note that heavily doped silicon (<sup>n</sup> > 2.2 <sup>10</sup><sup>19</sup> cm<sup>3</sup>

level is <sup>N</sup> = 5 1019 cm<sup>3</sup> [49].

N=5 <sup>10</sup><sup>25</sup> <sup>m</sup><sup>3</sup> in (b).

where ε<sup>∞</sup> is the background permittivity and ω<sup>p</sup> is the plasma frequency. The effective-medium approach [45] which is justified if the wavelength of the radiation considered is much larger

Figure 8. Geometry of the HMM. An interface separating two different semiconductor-dielectric-layered structures. Herein, indexes "1 and 2" correspond to the semiconductor and dielectric layers correspondingly.

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case of the MM/TCO interface, two types of the modes are always present, for higher TCO-filling

The complex mode structure (Figures 4 and 5) corresponding to either MM/dielectric or MM/ TCO interface emerges as a consequence of the confinement of plasmon polaritons in the direction perpendicular to the wave propagation. These electromagnetic surface waves arise via the coupling of the electromagnetic fields to oscillations of the conductor's electron plasma. The imaginary parts of the wave vector (i.e. absorption) are plotted in Figures 6 and 7. It should be mentioned that negative values of the absorption in Figures 6 and 7 result from non-physical solutions of the dispersion equation and have been omitted in line with [43]. Taking advantage of the absorption resonances, one can show that the simple multilayer structures without possessing any periodic corrugation have the prospective to act as directive

Another interesting MM structure depicted in Figure 8 is periodic stack of semiconductor-

<sup>ε</sup>1,3ð Þ¼ <sup>ω</sup> <sup>ε</sup><sup>∞</sup> � <sup>ω</sup><sup>2</sup>

where ε<sup>∞</sup> is the background permittivity and ω<sup>p</sup> is the plasma frequency. The effective-medium approach [45] which is justified if the wavelength of the radiation considered is much larger

Figure 8. Geometry of the HMM. An interface separating two different semiconductor-dielectric-layered structures.

Herein, indexes "1 and 2" correspond to the semiconductor and dielectric layers correspondingly.

p

<sup>ω</sup><sup>2</sup> <sup>þ</sup> <sup>i</sup>δω , (7)

The effective permittivity of the semiconductor (Si) can be calculated as follows:

ratio, the disappearance of modes in the gray region takes place (Figure 5(d)).

and monochromatic thermal sources [44].

12 Semiconductors - Growth and Characterization

3. Nanostructured semiconductor metamaterial

dielectric layers called hyperbolic metamaterial heterostructure.

Figure 9. The influence of (a) doping level N and (b) thickness of dielectric dd, on the real part of ε⊥. dd = 10 nm in (a) and N=5 <sup>10</sup><sup>25</sup> <sup>m</sup><sup>3</sup> in (b).

than the thickness of any layer is applied aiming to describe the optical response of such a system. The dispersion relation for the surface modes localized at the boundary separating two anisotropic media [42] is found by applying the appropriate boundary conditions, i.e. matching the tangential components of the electrical and magnetic fields at the interface. It is interesting to note that heavily doped silicon (<sup>n</sup> > 2.2 <sup>10</sup><sup>19</sup> cm<sup>3</sup> ) has been shown to exhibit metallic properties at terahertz frequencies [46, 47] and has the potential to replace metals in such applications [48]. The case of a heavy-doped Si is considered, assuming that the doping level is <sup>N</sup> = 5 1019 cm<sup>3</sup> [49].

It should be mentioned that dramatic control of the frequency range of the surface wave existence is mostly concerned with the modifications of the permittivities and thicknesses of the layers [50] employed in the HMMs. To further study the surface waves, the tangential components of the electric and magnetic fields at the interface should be evaluated, and a single surface mode with the propagation constant should be obtained aiming to get the unique dispersion relation for the surface modes confined at the interface between two metamaterials [42].

Using the (4) formula, we can describe the case ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> ¼ ε4, d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> reveals the dispersion as follows:

$$
\beta = k \sqrt{\frac{\varepsilon\_1 \varepsilon\_2}{\varepsilon\_1 + \varepsilon\_2}}.\tag{8}
$$

β ¼ k

various HM-based optical devices.

N=5 � <sup>10</sup><sup>25</sup> <sup>m</sup>�<sup>3</sup> in (b).

d1ε<sup>2</sup>

s

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ε1ε2ε4ð Þ d<sup>1</sup> þ d<sup>2</sup>

The permittivity spectra for the perpendicular components of the considered multilayer heterostructure are demonstrated in Figures 2(c), 3(a)–(c) and 9(a). Tuning the doping level of the semiconductor may open a gateway to the frequency control of the hyperbolic dispersion curve as shown in Figures 3(a), 4(a), 5(a) and 9(a). It is assumed that ds = 0.35 nm. PbS with relative permittivity ε<sup>d</sup> = 18.8 and slab thickness td = 10 nm is chosen as the dielectric layer. It is clear that one has the potential to achieve the resonant behaviour of ε<sup>⊥</sup> by varying the doping level; moreover, the increase in the doping level causes a tuning of the resonant frequencies over the higher frequency range. Because of these attractive properties, our semiconductorbased layered structure has the great potential in the application as the building block for

Figure 11. The influence of (a) doping level N and (b) thickness of dielectric dd, on the real part of εk. dd = 10 nm in (a) and

<sup>1</sup> þ d2ε1ε<sup>2</sup> þ d1ε2ε<sup>4</sup> þ d2ε1ε<sup>4</sup>

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(10)

15

The dispersion for the case of ε<sup>1</sup> ¼ ε<sup>3</sup> and ε<sup>2</sup> 6¼ ε4, d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> is as follows:

$$\beta = k \sqrt{\frac{\varepsilon\_1^2 \varepsilon\_2 \varepsilon\_4 (d\_1 + d\_2)(d\_3 + d\_4) \left(\frac{D}{d\_1 + d\_2} - \frac{B}{d\_3 + d\_4}\right)}{\left(\frac{\varepsilon\_1 \varepsilon\_2 D}{C} - \frac{\varepsilon\_1 \varepsilon\_4 B}{A}\right) \Box A}}\tag{9}$$

where A ¼ d4ε<sup>1</sup> þ d3ε4, B ¼ d3ε<sup>1</sup> þ d4ε4, C ¼ d1ε<sup>2</sup> þ d2ε<sup>1</sup> and D ¼ d1ε<sup>1</sup> þ d2ε2. In the case of ε<sup>1</sup> ¼ ε<sup>3</sup> and ε<sup>2</sup> 6¼ ε4, d<sup>1</sup> ¼ d<sup>3</sup> and d<sup>4</sup> ¼ d2:

Figure 10. The influence of (a) doping level N and (b) thickness of dielectric dd, on the imaginary part of ε⊥. dd = 10 nm in (a) and N = 5 � <sup>10</sup><sup>25</sup> <sup>m</sup>�<sup>3</sup> in (b).

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$$\beta = k \sqrt{\frac{\varepsilon\_1 \varepsilon\_2 \varepsilon\_4 (d\_1 + d\_2)}{d\_1 \varepsilon\_1^2 + d\_2 \varepsilon\_1 \varepsilon\_2 + d\_1 \varepsilon\_2 \varepsilon\_4 + d\_2 \varepsilon\_1 \varepsilon\_4}} \tag{10}$$

The permittivity spectra for the perpendicular components of the considered multilayer heterostructure are demonstrated in Figures 2(c), 3(a)–(c) and 9(a). Tuning the doping level of the semiconductor may open a gateway to the frequency control of the hyperbolic dispersion curve as shown in Figures 3(a), 4(a), 5(a) and 9(a). It is assumed that ds = 0.35 nm. PbS with relative permittivity ε<sup>d</sup> = 18.8 and slab thickness td = 10 nm is chosen as the dielectric layer. It is clear that one has the potential to achieve the resonant behaviour of ε<sup>⊥</sup> by varying the doping level; moreover, the increase in the doping level causes a tuning of the resonant frequencies over the higher frequency range. Because of these attractive properties, our semiconductorbased layered structure has the great potential in the application as the building block for various HM-based optical devices.

Using the (4) formula, we can describe the case ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> ¼ ε4, d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> reveals the

r

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ε1ε<sup>2</sup> ε<sup>1</sup> þ ε<sup>2</sup>

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

D <sup>d</sup>1þd<sup>2</sup> � <sup>B</sup> d3þd<sup>4</sup> � �

vuut , (9)

: (8)

β ¼ k

The dispersion for the case of ε<sup>1</sup> ¼ ε<sup>3</sup> and ε<sup>2</sup> 6¼ ε4, d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> is as follows:

where A ¼ d4ε<sup>1</sup> þ d3ε4, B ¼ d3ε<sup>1</sup> þ d4ε4, C ¼ d1ε<sup>2</sup> þ d2ε<sup>1</sup> and D ¼ d1ε<sup>1</sup> þ d2ε2.

<sup>1</sup>ε2ε4ð Þ d<sup>1</sup> þ d<sup>2</sup> ð Þ d<sup>3</sup> þ d<sup>4</sup>

ε1ε2D <sup>C</sup> � <sup>ε</sup>1ε4<sup>B</sup> A � �CA

Figure 10. The influence of (a) doping level N and (b) thickness of dielectric dd, on the imaginary part of ε⊥. dd = 10 nm in

β ¼ k

In the case of ε<sup>1</sup> ¼ ε<sup>3</sup> and ε<sup>2</sup> 6¼ ε4, d<sup>1</sup> ¼ d<sup>3</sup> and d<sup>4</sup> ¼ d2:

ε2

dispersion as follows:

14 Semiconductors - Growth and Characterization

(a) and N = 5 � <sup>10</sup><sup>25</sup> <sup>m</sup>�<sup>3</sup> in (b).

Figure 11. The influence of (a) doping level N and (b) thickness of dielectric dd, on the real part of εk. dd = 10 nm in (a) and N=5 � <sup>10</sup><sup>25</sup> <sup>m</sup>�<sup>3</sup> in (b).

Other than the doping level, the resonant behaviour of ε<sup>⊥</sup> was found to depend on the fill factions of the dielectric and semiconductor sheet, as shown in Figures 9(b) and 3(b). From Figure 9(b), the shift of the resonant frequency of ε<sup>⊥</sup> to the higher frequencies as the thickness dd is increased can be clearly distinguished.

#### 3.1. The mode structure

Guided by the homogenization of two HMs, the computed dispersion curves are demonstrated. Thus, in Figure 13, the dispersion curves for the case ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> ¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> ¼6 d<sup>3</sup> ¼6 d<sup>4</sup> are shown. In this case, we deal with the boundary of two metamaterials with d1 = 0.35 nm, d2 = 10 nm, d3 = 0.25 nm and d4 = 10.1 nm. Furthermore, due to the great interest in this case, Figure 13 refers to the dispersion of surface waves with the calculated effective parameters shown in Figures 9(a), 10(a), 11(a) and 12(a); the blue line is the free-space light line.

The frequency ranges of surface wave can be tuned by changing the doping level of silicon. As it is shown in Figure 9(a), doping level is correlated to the permittivity orthogonal component ε⊥. Moreover, dispersion curves are shifted to the lower and higher frequencies with the

decreases and increases in the doping level N accordingly. These tunability properties can be observed in Figure 13. Furthermore, this correlation can be used to engineer the metamaterial

Figure 13. The dispersion of surface waves (a); propagation lengths (b) and absorption (c) at different doping levels of silicon, where dd = 10 nm, ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> ¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> and the blue line in (a) is the free-space light line.

As the silicon is not modeled as lossless, β is complex, leading to a finite propagation length

radiative (bound) SP modes throughout the certain frequency range. All the considered cases are of particular interest due to the fact that their dispersion relations cross the light line and a significant portion at lower frequencies lies above the free-space light line, which usually splits up non-radiative (bound) and radiative (leaky) regions. For the bound modes, longer propagation lengths take place at lower frequencies owning the dispersion that is close to linear. Mode corresponding to the case N = 4 � 1025 <sup>m</sup>�<sup>3</sup> possesses the longer propagation length than the

.

The existence of the boundary modes associated with the second case under consideration, i.e. ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> 6¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d4, is also of the particular importance. Thus, Figure 14

) always lie to the right side of the light line and remain non-

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,N=3 � 1025 <sup>m</sup>�<sup>3</sup>

,

(Eq. (2)), drawn in Figure 13(b). In Figure 13 the four modes (N = 2 � 1025 <sup>m</sup>�<sup>3</sup>

surface wave just by controlling silicon sheet doping level.

,N=5 � 1025 <sup>m</sup>�<sup>3</sup>

mode corresponding to the case N = 5 � 1025 <sup>m</sup>�<sup>3</sup>

N=4 � 1025 <sup>m</sup>�<sup>3</sup>

Figure 12. The influence of (a) doping level N and (b) thickness of dielectric dd, on the imaginary part of εk. dd = 10 nm in (a) and N = 5 � 1025 <sup>m</sup>�<sup>3</sup> in (b).

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Other than the doping level, the resonant behaviour of ε<sup>⊥</sup> was found to depend on the fill factions of the dielectric and semiconductor sheet, as shown in Figures 9(b) and 3(b). From Figure 9(b), the shift of the resonant frequency of ε<sup>⊥</sup> to the higher frequencies as the thickness

Guided by the homogenization of two HMs, the computed dispersion curves are demonstrated. Thus, in Figure 13, the dispersion curves for the case ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> ¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> ¼6 d<sup>3</sup> ¼6 d<sup>4</sup> are shown. In this case, we deal with the boundary of two metamaterials with d1 = 0.35 nm, d2 = 10 nm, d3 = 0.25 nm and d4 = 10.1 nm. Furthermore, due to the great interest in this case, Figure 13 refers to the dispersion of surface waves with the calculated effective parameters

The frequency ranges of surface wave can be tuned by changing the doping level of silicon. As it is shown in Figure 9(a), doping level is correlated to the permittivity orthogonal component ε⊥. Moreover, dispersion curves are shifted to the lower and higher frequencies with the

Figure 12. The influence of (a) doping level N and (b) thickness of dielectric dd, on the imaginary part of εk. dd = 10 nm in

shown in Figures 9(a), 10(a), 11(a) and 12(a); the blue line is the free-space light line.

dd is increased can be clearly distinguished.

16 Semiconductors - Growth and Characterization

3.1. The mode structure

(a) and N = 5 � 1025 <sup>m</sup>�<sup>3</sup> in (b).

Figure 13. The dispersion of surface waves (a); propagation lengths (b) and absorption (c) at different doping levels of silicon, where dd = 10 nm, ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> ¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> and the blue line in (a) is the free-space light line.

decreases and increases in the doping level N accordingly. These tunability properties can be observed in Figure 13. Furthermore, this correlation can be used to engineer the metamaterial surface wave just by controlling silicon sheet doping level.

As the silicon is not modeled as lossless, β is complex, leading to a finite propagation length (Eq. (2)), drawn in Figure 13(b). In Figure 13 the four modes (N = 2 � 1025 <sup>m</sup>�<sup>3</sup> ,N=3 � 1025 <sup>m</sup>�<sup>3</sup> , N=4 � 1025 <sup>m</sup>�<sup>3</sup> ,N=5 � 1025 <sup>m</sup>�<sup>3</sup> ) always lie to the right side of the light line and remain nonradiative (bound) SP modes throughout the certain frequency range. All the considered cases are of particular interest due to the fact that their dispersion relations cross the light line and a significant portion at lower frequencies lies above the free-space light line, which usually splits up non-radiative (bound) and radiative (leaky) regions. For the bound modes, longer propagation lengths take place at lower frequencies owning the dispersion that is close to linear. Mode corresponding to the case N = 4 � 1025 <sup>m</sup>�<sup>3</sup> possesses the longer propagation length than the mode corresponding to the case N = 5 � 1025 <sup>m</sup>�<sup>3</sup> .

The existence of the boundary modes associated with the second case under consideration, i.e. ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> 6¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d4, is also of the particular importance. Thus, Figure 14

Figure 14. The dispersion of surface waves (a), propagation lengths (b) and absorption (c) at different doping levels, where dd = 10 nm, ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> 6¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> and the blue line in (a) is the free-space light line.

tackles this problem by displaying four modes at the boundary of two different metamaterials with ε<sup>4</sup> = 2.25. As seen from Figure 14, the smallest asymptotic frequency corresponds to the case N = 2 � 1025 <sup>m</sup>�<sup>3</sup> .

4. Conclusion

thickness of dielectric dd. Herein N = 5 <sup>10</sup><sup>25</sup> <sup>m</sup><sup>3</sup>

During a study of the HSPPs in a one-dimensional TCO-dielectric MM, we can see that similar to graphene-dielectric MM [45], TCO-dielectric MM supports traditional-like SPPs having different patterns corresponding to two different interfaces. The dispersion equations of HSPPs are obtained based on the theoretical approach [42, 45]. Five kinds of HSPP, among which three kinds are new types of HSPPs and one is the Dyakonov-like SPP and another is the traditionallike SPP have been predicted. The existence of these HSPPs is dramatically influenced by the properties of and the relation among the principal values of the effective permittivity and the dielectric constant of the covering medium. It is worthwhile mentioning that the new types of the HSPPs arise because the principal values of the effective permittivity used in this chapter are functions of frequency and can be negative or positive. Moreover, it was demonstrated that used approach allows to predict surface mode with the dispersion that coincides with the dispersion of a surface plasmon at the boundary of two isotropic media corresponding to the MM interface

Figure 15. The dependences of dispersion of surface waves, propagation lengths and absorption on (a), (b) and (c) the

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19

.

if the material at the right-hand side is the same as employed in the MM.

In contrast to the previous case, we now discuss the instance denoted as ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> 6¼ ε<sup>4</sup> and d<sup>1</sup> ¼ d3, d<sup>2</sup> ¼ d<sup>4</sup> and shown by means of the dispersion diagrams of the TM modes. Thus, Figure 15 shows the dispersion curves of four different modes. The assessment and control of variation of the dielectric and semiconductor sheets' fill factors is of critical importance (Figure 15). First, the impact of the thickness of the dielectric dd on the dispersion curve (see Figure 15(a)) is considered. It is found that the upper limit moves to the higher frequencies as dd is increased. The former is consistent with the effect of dd on the frequency range of ε⊥. The dependence of the frequency range of the surface waves existence on the thickness of dielectric stands for as the most critical feature of the HMs providing an unprecedented degree of freedom to control the surface wave at the near-infrared frequencies. In Figures 14(a) and 15(a), it is interesting to observe the Ferrell-Berreman modes which exist at energies near the ENZ of the hyperbolic metamaterial to the left of the light line [51–53].

Investigation of the Nanostructured Semiconductor Metamaterials http://dx.doi.org/10.5772/intechopen.72801 19

Figure 15. The dependences of dispersion of surface waves, propagation lengths and absorption on (a), (b) and (c) the thickness of dielectric dd. Herein N = 5 <sup>10</sup><sup>25</sup> <sup>m</sup><sup>3</sup> .

#### 4. Conclusion

tackles this problem by displaying four modes at the boundary of two different metamaterials with ε<sup>4</sup> = 2.25. As seen from Figure 14, the smallest asymptotic frequency corresponds to the

Figure 14. The dispersion of surface waves (a), propagation lengths (b) and absorption (c) at different doping levels,

where dd = 10 nm, ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> 6¼ ε<sup>4</sup> and d<sup>1</sup> 6¼ d<sup>2</sup> 6¼ d<sup>3</sup> 6¼ d<sup>4</sup> and the blue line in (a) is the free-space light line.

In contrast to the previous case, we now discuss the instance denoted as ε<sup>1</sup> ¼ ε3, ε<sup>2</sup> 6¼ ε<sup>4</sup> and d<sup>1</sup> ¼ d3, d<sup>2</sup> ¼ d<sup>4</sup> and shown by means of the dispersion diagrams of the TM modes. Thus, Figure 15 shows the dispersion curves of four different modes. The assessment and control of variation of the dielectric and semiconductor sheets' fill factors is of critical importance (Figure 15). First, the impact of the thickness of the dielectric dd on the dispersion curve (see Figure 15(a)) is considered. It is found that the upper limit moves to the higher frequencies as dd is increased. The former is consistent with the effect of dd on the frequency range of ε⊥. The dependence of the frequency range of the surface waves existence on the thickness of dielectric stands for as the most critical feature of the HMs providing an unprecedented degree of freedom to control the surface wave at the near-infrared frequencies. In Figures 14(a) and 15(a), it is interesting to observe the Ferrell-Berreman modes which exist at energies near the ENZ

case N = 2 � 1025 <sup>m</sup>�<sup>3</sup>

.

18 Semiconductors - Growth and Characterization

of the hyperbolic metamaterial to the left of the light line [51–53].

During a study of the HSPPs in a one-dimensional TCO-dielectric MM, we can see that similar to graphene-dielectric MM [45], TCO-dielectric MM supports traditional-like SPPs having different patterns corresponding to two different interfaces. The dispersion equations of HSPPs are obtained based on the theoretical approach [42, 45]. Five kinds of HSPP, among which three kinds are new types of HSPPs and one is the Dyakonov-like SPP and another is the traditionallike SPP have been predicted. The existence of these HSPPs is dramatically influenced by the properties of and the relation among the principal values of the effective permittivity and the dielectric constant of the covering medium. It is worthwhile mentioning that the new types of the HSPPs arise because the principal values of the effective permittivity used in this chapter are functions of frequency and can be negative or positive. Moreover, it was demonstrated that used approach allows to predict surface mode with the dispersion that coincides with the dispersion of a surface plasmon at the boundary of two isotropic media corresponding to the MM interface if the material at the right-hand side is the same as employed in the MM.

Moreover, a new kind of surface wave between two nanostructured semiconductor metamaterials was demonstrated. It is shown that the dispersion diagrams are sensitively dependent on the semiconductor parameters. These findings open the gateway towards potential applications in both classical and quantum optical signal communication and processing.

[11] Ozbay E. Plasmonics: Merging photonics and electronics at nanoscale dimensions. Sci-

Investigation of the Nanostructured Semiconductor Metamaterials

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21

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#### Author details

Aleksej Trofimov<sup>1</sup> \*, Tatjana Gric<sup>1</sup> and Ortwin Hess<sup>2</sup>


#### References


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1 Vilnius Gediminas Technical University, Vilnius, Lithuania

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**Chapter 2**

**Provisional chapter**

**Understanding the Mechanisms that Affect the Quality**

**of Electrochemically Grown Semiconducting Nanowires**

Template-assisted synthesis of nanowires is a simple electrochemical technique commonly used in the fabrication of semiconducting nanowires. It is an easy and cost-effective approach compared to conventional lithography, which requires expensive equipment. The focus of this chapter is on the various mechanisms involving mass transport of ions during successive stages of the template-assisted electrochemical growth of indium antimonide (InSb) nanowires. The nanowires were grown in two different templates such as commercially available anodic aluminum oxide (AAO) templates and polycarbonate membranes. The chapter also presents the results of characterizing the InSb nanowires connected in a field effect transistor (FET) configuration. The Sb-rich InSb nanowires that were fabricated by DC electrodeposition in nanoporous AAO exhibited hole-dominated transport (p-type conduction). Temperature-dependent transport measurement shows

**Keywords:** semiconducting nanowire, InSb, electrodeposition, anodic aluminum oxide

In recent years, anodic aluminum oxide (AAO) templates have been used extensively for fabricating myriads of nanostructures. AAO is a self-organized nanostructured material containing a high density of uniform cylindrical pores that are aligned perpendicular to the surface and extend over the entire thickness of the template. AAO is also optically transparent, electrically insulating, thermally and mechanically robust and chemically inert. Nanowire arrays that are grown in AAO template pores are used in many applications such as energy conversion [1], energy-storage devices [2, 3], electronics [4], metamaterials [5],

**Understanding the Mechanisms that Affect the Quality** 

DOI: 10.5772/intechopen.71631

© 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,

© 2018 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.

and reproduction in any medium, provided the original work is properly cited.

**of Electrochemically Grown Semiconducting**

**Nanowires**

**Abstract**

**1. Introduction**

Abhay Singh and Usha Philipose

Abhay Singh and Usha Philipose

http://dx.doi.org/10.5772/intechopen.71631

Additional information is available at the end of the chapter

the semiconducting nature of these nanowires.

(AAO), nanowire growth mechanism

Additional information is available at the end of the chapter

**Provisional chapter**

## **Understanding the Mechanisms that Affect the Quality of Electrochemically Grown Semiconducting Nanowires Understanding the Mechanisms that Affect the Quality of Electrochemically Grown Semiconducting Nanowires**

DOI: 10.5772/intechopen.71631

Abhay Singh and Usha Philipose Additional information is available at the end of the chapter

Abhay Singh and Usha Philipose

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.71631

#### **Abstract**

Template-assisted synthesis of nanowires is a simple electrochemical technique commonly used in the fabrication of semiconducting nanowires. It is an easy and cost-effective approach compared to conventional lithography, which requires expensive equipment. The focus of this chapter is on the various mechanisms involving mass transport of ions during successive stages of the template-assisted electrochemical growth of indium antimonide (InSb) nanowires. The nanowires were grown in two different templates such as commercially available anodic aluminum oxide (AAO) templates and polycarbonate membranes. The chapter also presents the results of characterizing the InSb nanowires connected in a field effect transistor (FET) configuration. The Sb-rich InSb nanowires that were fabricated by DC electrodeposition in nanoporous AAO exhibited hole-dominated transport (p-type conduction). Temperature-dependent transport measurement shows the semiconducting nature of these nanowires.

**Keywords:** semiconducting nanowire, InSb, electrodeposition, anodic aluminum oxide (AAO), nanowire growth mechanism

#### **1. Introduction**

In recent years, anodic aluminum oxide (AAO) templates have been used extensively for fabricating myriads of nanostructures. AAO is a self-organized nanostructured material containing a high density of uniform cylindrical pores that are aligned perpendicular to the surface and extend over the entire thickness of the template. AAO is also optically transparent, electrically insulating, thermally and mechanically robust and chemically inert. Nanowire arrays that are grown in AAO template pores are used in many applications such as energy conversion [1], energy-storage devices [2, 3], electronics [4], metamaterials [5],

© 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. © 2018 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.

optoelectronics [6], photonics [7], and piezoelectrics [8]. Template-assisted electrochemical growth of large-scale nanowire arrays is attractive because they are readily scalable to mass production. Moreover, it is possible to obtain an array of uniform diameter and length and a surface that can be engineered by nonuniformity of the template tubes.

with time; a fact that contradicts the previous study where current remains constant. Here, we discuss the more realistic model that describes a situation where the electrolyte concentration

Understanding the Mechanisms that Affect the Quality of Electrochemically Grown…

http://dx.doi.org/10.5772/intechopen.71631

27

The electrodeposition growth process of nanowires in porous templates has been modeled by considering different mechanisms of mass transport at different stages of the growth process [24]. There are mainly three different mechanisms that could explain the growth process: (i) diffusion process (arising from concentration gradients existing between electrolyte in the template pores and in the solution), (ii) convection process (arising from movement of deposition ions in the electrolyte), and (iii) migration of ions into the template pores. Contribution of each process can be controlled by modifying the cell design, reducing depositing ion concentrations compared to mixing electrolyte, and stirring the electrolyte during growth pro-

Mass transport of ions during nanowire growth is mainly controlled by a diffusion process. The three main stages of this process are shown in **Figure 1**. **Figure 1(a)** shows the earlier stage of diffusion where diffusion front propagates in only one dimension inside the pores.

**Figure 1(b)** shows the second stage where diffusion front reached the mouth of the pore and opens in three-dimensional hemispherical front. As soon as diffusion front reaches at the mouth, then there is a concentration gradient builds up between the mouth and the bulk. **Figure 1(c)** shows the third stage where linear diffusion happens inside and outside of the pores. This diffusion-limited electrodeposition results in a morphological instability driven by individual nanowires [25]. In template-assisted electrodeposition, nanowire length must follow the length of AAO pore, but some nanowire grows fast compared to others and ends up mushroom-like structures at the top surface of AAO template [26–31]. Recent study [25] also showed that this growth instability can be avoided by introducing temperature gradient

**Figure 1.** Schematic of growth of nanowire in a single nanopore: (a) linear diffusion, (b) linear and semi-infinite hemispherical

diffusion, and (c) semi-infinite planar diffusion outside the pore (Adapted from Ref. [24]).

at the mouth of the pores is different than the bulk during electrodeposition.

cess to avoid the concentration gradient.

between bottom and top of the AAO pores.

**2.1. Diffusion mechanism in nanowire growth process**

This chapter reviews the various processes that govern the filling of the nanochannel alumina template pores resulting in a systematic growth of nanowires within the pores. We also discuss the challenges faced in achieving uniform growth including effects of pore-wetting and aeration. Any nonuniformity in the alumina template tubes also result in nanowires with rough surfaces, an exploitable result for use of nanowires in thermoelectric applications.

Over the last decade, there has been increasing interest in nanostructures of III–V semiconducting materials due to their potential applications in electronic and optoelectronic devices [9]. After identifying the various mechanisms of electrochemical growth of semiconducting nanowires, we present the specific growth process conditions for growth of indium antimonide (InSb) nanowires. Indium antimonide (InSb) bulk is a promising III–V direct bandgap semiconductor material with zinc blende (FCC) structure [10–14]. InSb has a high room temperature mobility (electron and hole mobility [15] of 77,000 and 850 cm<sup>2</sup> V−1 s−1, respectively), low electron effective mass [16] of 0.014, and low direct bandgap (E<sup>g</sup> = 0.17 eV, at 300 K), and large Lande g-factor of 51, [17] making it suitable for use in applications such as high speed, low-power transistors, tunneling field effect transistors (FETs), infrared optoelectronics [18], and magnetoresistive sensors [19].

Nanowires can be grown via electrodeposition in different nanoporous templates like anodic aluminum oxide (AAO) and ion track-etched polycarbonate. Every template type has its own advantage, disadvantage, and application.

### **2. Growth process: convection, migration, and diffusion**

The benchmark method for synthesis of nanowire was suggested in 1986 by Martin et al. [20], where nanowires were grown inside polycarbonate membranes. This method has since been widely used to synthesize nanowires in both polycarbonate and anodic aluminum oxide (AAO) membranes. These membranes contain a high density of cylindrical pores that are perpendicular to the membrane surface and in most cases, these pores penetrate the entire membrane thickness. There have been several studies on electrochemical growth of nanowires but very small number of investigations on understanding the growth mechanism of nanowires in AAO pores.

Template-assisted nanowire growth process consists of two major steps: electrochemical reduction of the cation inside the pore, followed by removal of the template. Several studies [21, 22] have shown that during the pore-filling stage (nanowire growth), the total constant current density is independent of flux variation inside the pore and concentration of electrolyte at the mouth of pore during electrodeposition. Recently, another study [23] showed that the electrolyte concentration at the mouth of the pore remains constant and current decays with time; a fact that contradicts the previous study where current remains constant. Here, we discuss the more realistic model that describes a situation where the electrolyte concentration at the mouth of the pores is different than the bulk during electrodeposition.

The electrodeposition growth process of nanowires in porous templates has been modeled by considering different mechanisms of mass transport at different stages of the growth process [24]. There are mainly three different mechanisms that could explain the growth process: (i) diffusion process (arising from concentration gradients existing between electrolyte in the template pores and in the solution), (ii) convection process (arising from movement of deposition ions in the electrolyte), and (iii) migration of ions into the template pores. Contribution of each process can be controlled by modifying the cell design, reducing depositing ion concentrations compared to mixing electrolyte, and stirring the electrolyte during growth process to avoid the concentration gradient.

#### **2.1. Diffusion mechanism in nanowire growth process**

optoelectronics [6], photonics [7], and piezoelectrics [8]. Template-assisted electrochemical growth of large-scale nanowire arrays is attractive because they are readily scalable to mass production. Moreover, it is possible to obtain an array of uniform diameter and length and

This chapter reviews the various processes that govern the filling of the nanochannel alumina template pores resulting in a systematic growth of nanowires within the pores. We also discuss the challenges faced in achieving uniform growth including effects of pore-wetting and aeration. Any nonuniformity in the alumina template tubes also result in nanowires with rough surfaces, an exploitable result for use of nanowires in thermoelectric applications.

Over the last decade, there has been increasing interest in nanostructures of III–V semiconducting materials due to their potential applications in electronic and optoelectronic devices [9]. After identifying the various mechanisms of electrochemical growth of semiconducting nanowires, we present the specific growth process conditions for growth of indium antimonide (InSb) nanowires. Indium antimonide (InSb) bulk is a promising III–V direct bandgap semiconductor material with zinc blende (FCC) structure [10–14]. InSb has a high room temperature mobility (electron and hole mobility [15] of 77,000 and 850 cm<sup>2</sup> V−1 s−1, respectively), low electron effective mass [16] of 0.014, and low direct bandgap (E<sup>g</sup> = 0.17 eV, at 300 K), and large Lande g-factor of 51, [17] making it suitable for use in applications such as high speed, low-power transistors, tunneling field effect transistors (FETs), infrared optoelectronics [18],

Nanowires can be grown via electrodeposition in different nanoporous templates like anodic aluminum oxide (AAO) and ion track-etched polycarbonate. Every template type has its own

The benchmark method for synthesis of nanowire was suggested in 1986 by Martin et al. [20], where nanowires were grown inside polycarbonate membranes. This method has since been widely used to synthesize nanowires in both polycarbonate and anodic aluminum oxide (AAO) membranes. These membranes contain a high density of cylindrical pores that are perpendicular to the membrane surface and in most cases, these pores penetrate the entire membrane thickness. There have been several studies on electrochemical growth of nanowires but very small number of investigations on understanding the growth mechanism of

Template-assisted nanowire growth process consists of two major steps: electrochemical reduction of the cation inside the pore, followed by removal of the template. Several studies [21, 22] have shown that during the pore-filling stage (nanowire growth), the total constant current density is independent of flux variation inside the pore and concentration of electrolyte at the mouth of pore during electrodeposition. Recently, another study [23] showed that the electrolyte concentration at the mouth of the pore remains constant and current decays

**2. Growth process: convection, migration, and diffusion**

a surface that can be engineered by nonuniformity of the template tubes.

and magnetoresistive sensors [19].

26 Semiconductors - Growth and Characterization

nanowires in AAO pores.

advantage, disadvantage, and application.

Mass transport of ions during nanowire growth is mainly controlled by a diffusion process. The three main stages of this process are shown in **Figure 1**. **Figure 1(a)** shows the earlier stage of diffusion where diffusion front propagates in only one dimension inside the pores.

**Figure 1(b)** shows the second stage where diffusion front reached the mouth of the pore and opens in three-dimensional hemispherical front. As soon as diffusion front reaches at the mouth, then there is a concentration gradient builds up between the mouth and the bulk. **Figure 1(c)** shows the third stage where linear diffusion happens inside and outside of the pores. This diffusion-limited electrodeposition results in a morphological instability driven by individual nanowires [25]. In template-assisted electrodeposition, nanowire length must follow the length of AAO pore, but some nanowire grows fast compared to others and ends up mushroom-like structures at the top surface of AAO template [26–31]. Recent study [25] also showed that this growth instability can be avoided by introducing temperature gradient between bottom and top of the AAO pores.

**Figure 1.** Schematic of growth of nanowire in a single nanopore: (a) linear diffusion, (b) linear and semi-infinite hemispherical diffusion, and (c) semi-infinite planar diffusion outside the pore (Adapted from Ref. [24]).

The diffusion mechanism during various stages of nanowire growth is described by the S-curve of current-time plot, as shown in **Figure 2(a)**.

region (III) of the S-curve plot, current does not increase significantly and tends to saturate, as compared to region (II) of the plot. **Figure 2(b)** and **(c)** shows cross-sectional SEM images of template in the region (I) and (III), respectively. **Figure 2(b)** shows partially filled pores, which correspond to a linear diffusion in region (I) of the S-curve. **Figure 4(c)** shows the completely filled pores and hemispherical caps that merge to form a dense mushroom-like

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The effect of convection on nanowire growth process is dependent on cell design. The three electrodes in an electrochemical cell are: (i) working electrode (substrate); (ii) counter electrode (Pt wire/mesh); and (iii) reference electrode. To minimize the mechanism of natural convection, electrodeposition cell configuration is an integral part of nanowire growth in nanoporous templates. In **Figure 3(a)** and **(b)**, the cathode and anode are placed vertically above each other, whereas in **Figure 3(c)**, the cathode and anode are placed parallel to each other. The effects of convection are considered to be minimal for the design in **Figure 3(c)**. Konishi et al. [33] reported that the electrodeposition current increases in the case of anode over cathode **Figure 3(b)** configuration, and current decreases in the case of cathode over anode configuration **Figure 3(a)** during early stages of the electrodeposition. They have found the considerable difference in the transient behavior of the current in 100- and 200-nm-sized pores in certain stage of electrodeposition, which is caused by the

The effects of natural convection are proven to be significant and reported in Ref [32] for the schematic shown in **Figure 4**, where an increase in convection current because of this mechanism is shown for pores size as small as 100 nm. This effect was observed during Cu electrodeposition.

To explain the InSb nanowire growth process in the following section, the effect of migration of In and Sb ions has not been considered in this discussion. Since concentration of these ions in supporting electrolyte is low, the contribution of ion migration is considered to be lower than diffusion.

**Figure 3.** Schematic of cell designs: (a) cathode over anode; (b) anode over cathode; and (c) cathode and electrode placed laterally and parallel to each other. Letter A represents working electrode (Cathode), C represents secondary electrode

overgrowth on the surface of the pores.

electrolytic cell configuration.

(Anode), and B represents reference electrode.

**2.2. Convection mechanism in growth of nanowires**

**2.3. Migration of ions during growth of nanowires**

The entire mechanism involves three consecutive growth stages: (i) one dimensional diffusion front inside AAO pores; (ii) linear diffusion inside the pores and hemispherical diffusion at the mouth of the pores; and (iii) merging of hemispherical diffusion front and resulting in planned front over the surface of the mouth. The electrodeposition time and average length of wire have been controlled by monitoring the deposition current versus time at given deposition voltage −1.5 V which was maintained with respect to reference electrode. Inset of **Figure 4(a)** shows the schematic of the various stages of electrodeposition.

In first stage of growth process, concentration of the electrolyte at the mouth of pores matches the concentration of the bulk of the electrolyte (c<sup>b</sup> ). The linear diffusion front passing through pore length results in transportation of In and Sb ions inside the pores of the AAO template. The diffusion fronts between individual pores have not been overlapped at this stage. The concentration within the pore changes as nanowire growth starts and it has been expressed by Fick's law as: *c*(*x*, *t*) = *c <sup>b</sup>* erf[ \_\_\_\_ *<sup>x</sup>* 2 √ \_\_\_ *Dt*], where 2√ \_\_\_ *Dt* is the diffusion length (L) that corresponds how far concentration varies along the pore length at time t. The initial pore length is L<sup>0</sup> at t = 0. There is a sharp decrease in electrodeposition current during this time, as is expected for diffusion-controlled process. The current was observed to decrease from 35 to 20 mA in region (I) of the plot in **Figure 2(a)** which corresponds to region (i) of growth process. In region (i), In and Sb ions diffuse to the bottom of the pore in the tubular column of AAO template to initiate nanowire growth.

During region (II) of the growth process, initial linear diffusion front reaches the mouth of the pores and three dimensional hemispherical diffusion front develops at the mouth. During this stage in region (II) of plot **Figure 4**, the steady increase in current has been observed and is believed to increase the diffusion of ions to the nanowire growth front. During the final stage in region (III), the nanowires have filled the pore completely and the hemispherical cap merges with each other and formed a continuous rough film on the surface of template. In the

**Figure 2.** Evolution of nanowire growth in the AAO template pores: (a) current-time plot of InSb nanowire growth showing different growth regions; (b) cross-sectional SEM image of the template in region I. The thin bottom layer is a gold film. The pores are partially filled with InSb. This region corresponds to a planar diffusion inside the pore; and (c) cross-sectional SEM image of region III where NWs have filled the pores and the dome-shaped tips collapse to form mushroom-shaped clusters, corresponding to semi-infinite planar diffusion at the mouth of the pore [32].

region (III) of the S-curve plot, current does not increase significantly and tends to saturate, as compared to region (II) of the plot. **Figure 2(b)** and **(c)** shows cross-sectional SEM images of template in the region (I) and (III), respectively. **Figure 2(b)** shows partially filled pores, which correspond to a linear diffusion in region (I) of the S-curve. **Figure 4(c)** shows the completely filled pores and hemispherical caps that merge to form a dense mushroom-like overgrowth on the surface of the pores.

#### **2.2. Convection mechanism in growth of nanowires**

The diffusion mechanism during various stages of nanowire growth is described by the

The entire mechanism involves three consecutive growth stages: (i) one dimensional diffusion front inside AAO pores; (ii) linear diffusion inside the pores and hemispherical diffusion at the mouth of the pores; and (iii) merging of hemispherical diffusion front and resulting in planned front over the surface of the mouth. The electrodeposition time and average length of wire have been controlled by monitoring the deposition current versus time at given deposition voltage −1.5 V which was maintained with respect to reference electrode. Inset of **Figure 4(a)**

In first stage of growth process, concentration of the electrolyte at the mouth of pores matches

pore length results in transportation of In and Sb ions inside the pores of the AAO template. The diffusion fronts between individual pores have not been overlapped at this stage. The concentration within the pore changes as nanowire growth starts and it has been expressed

\_\_\_

t = 0. There is a sharp decrease in electrodeposition current during this time, as is expected for diffusion-controlled process. The current was observed to decrease from 35 to 20 mA in region (I) of the plot in **Figure 2(a)** which corresponds to region (i) of growth process. In region (i), In and Sb ions diffuse to the bottom of the pore in the tubular column of AAO tem-

During region (II) of the growth process, initial linear diffusion front reaches the mouth of the pores and three dimensional hemispherical diffusion front develops at the mouth. During this stage in region (II) of plot **Figure 4**, the steady increase in current has been observed and is believed to increase the diffusion of ions to the nanowire growth front. During the final stage in region (III), the nanowires have filled the pore completely and the hemispherical cap merges with each other and formed a continuous rough film on the surface of template. In the

**Figure 2.** Evolution of nanowire growth in the AAO template pores: (a) current-time plot of InSb nanowire growth showing different growth regions; (b) cross-sectional SEM image of the template in region I. The thin bottom layer is a gold film. The pores are partially filled with InSb. This region corresponds to a planar diffusion inside the pore; and (c) cross-sectional SEM image of region III where NWs have filled the pores and the dome-shaped tips collapse to form

mushroom-shaped clusters, corresponding to semi-infinite planar diffusion at the mouth of the pore [32].

how far concentration varies along the pore length at time t. The initial pore length is L<sup>0</sup>

*Dt*], where 2√

). The linear diffusion front passing through

*Dt* is the diffusion length (L) that corresponds

at

S-curve of current-time plot, as shown in **Figure 2(a)**.

28 Semiconductors - Growth and Characterization

shows the schematic of the various stages of electrodeposition.

the concentration of the bulk of the electrolyte (c<sup>b</sup>

erf[ \_\_\_\_ *<sup>x</sup>* 2 √ \_\_\_

by Fick's law as: *c*(*x*, *t*) = *c <sup>b</sup>*

plate to initiate nanowire growth.

The effect of convection on nanowire growth process is dependent on cell design. The three electrodes in an electrochemical cell are: (i) working electrode (substrate); (ii) counter electrode (Pt wire/mesh); and (iii) reference electrode. To minimize the mechanism of natural convection, electrodeposition cell configuration is an integral part of nanowire growth in nanoporous templates. In **Figure 3(a)** and **(b)**, the cathode and anode are placed vertically above each other, whereas in **Figure 3(c)**, the cathode and anode are placed parallel to each other. The effects of convection are considered to be minimal for the design in **Figure 3(c)**. Konishi et al. [33] reported that the electrodeposition current increases in the case of anode over cathode **Figure 3(b)** configuration, and current decreases in the case of cathode over anode configuration **Figure 3(a)** during early stages of the electrodeposition. They have found the considerable difference in the transient behavior of the current in 100- and 200-nm-sized pores in certain stage of electrodeposition, which is caused by the electrolytic cell configuration.

The effects of natural convection are proven to be significant and reported in Ref [32] for the schematic shown in **Figure 4**, where an increase in convection current because of this mechanism is shown for pores size as small as 100 nm. This effect was observed during Cu electrodeposition.

#### **2.3. Migration of ions during growth of nanowires**

To explain the InSb nanowire growth process in the following section, the effect of migration of In and Sb ions has not been considered in this discussion. Since concentration of these ions in supporting electrolyte is low, the contribution of ion migration is considered to be lower than diffusion.

**Figure 3.** Schematic of cell designs: (a) cathode over anode; (b) anode over cathode; and (c) cathode and electrode placed laterally and parallel to each other. Letter A represents working electrode (Cathode), C represents secondary electrode (Anode), and B represents reference electrode.

SEM images of the (a) top and (b) bottom surfaces of commercial AAO templates. Inset shows higher magnification of the same surfaces. It is obvious from these images that the top and bottom pores do not have the same diameter, and the top surface pores are larger in diameter compared to the bottom pores. It is this difference in pore diameters that leads to most researchers depositing gold on the lower diameter side of the AAO template, thus preventing overfilling

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**3. Fabrication and characterization of InSb nanowires in commercial** 

, 0.1 M SbCl<sup>3</sup>

, 0.36 M C<sup>6</sup>

value of the electrolyte was adjusted to 1.7 and electrodeposition was performed at room temperature (200°C). The citrate ions have been used as complexing agents that bring the deposition potential of In and Sb closer to maintain a binary growth during the deposition of the InSb nanowires. The complete electrodeposition process was performed for 50 min at a potential of −1.5 V with respect to the reference electrode (Ag/AgCl). Also, magnetic stirring was used to maintain the uniform concentration of electrolyte during growth of nanowires. On completion of the deposition process, the AAO template was carefully rinsed several times with DI water, and then AAO template was removed from copper tape for further processing.

Following steps are involved in extraction of nanowires from AAO template: (a) AAO template was placed in a clean 80 ml beaker with the gold side facing upward and gold film was removed by 1 M potassium iodide (KI). To dissolve the AAM template, a few drops of 1 M KOH solution was used. This was done with vigorous shaking for about 20 min (b) following dissolution of the template; 15 ml of DI water was added to the solution, which was then sonicated for 10 min. The sonication step helps to dissolve the AAO completely, while breaking the InSb nanowires off the InSb crust (overgrown areas). (c) DI water as added to wash out the precipitates in the bottom of the beaker into the vial; the process was repeated until there was no visible precipitate in the beaker. This step was extremely important, since most of the nanowires were obtained from these precipitates. (d) To further dilute the KOH and AAO in the solution, both vials were placed in a centrifuge (Fisher Scientific Centrifuge, model: 228

H8 O7 H2

O, and C<sup>6</sup>

H5 Na3O<sup>7</sup>

. The pH

Using the commercial AAO templates discussed in the previous section, one surface of the AAO template was coated with a thin film of Au (∼150 nm) using thermal evaporation. Prior to Au evaporation, the pores were widened nominally by soaking the template in 5% phosphoric acid (H3PO3) for 2 min at 30°C. The electrodeposition process was conducted in a three-electrode cell with the AAO template Au side as the working electrode, platinum electrode as the counter electrode and Ag/AgCl as the reference electrode. Crucial part of deposition process was to cover the edges of the AAO template with an insulating material and then attaching it to a conducting copper tape. The side of the AAO pore, which was covered with thin film of Au, was attached to conducting glue side of copper tape. After that, insulating polymer was applied to avoid any conducting path except through the open AAO pores which terminates at the Au film at the bottom of the pores. The entire electrodeposition process was controlled and carried out by a potentiostat (Princeton Applied Research, model: Potentiostat/Galvanostat 263A). The deposition parameters used for electrolyte in this experi-

of the pores.

**AAO template**

ment are as follows: 0.15 M InCl3

**Figure 4.** Schematic diagram convection in and around nanoporous pores in the configuration of anode over cathode.

To conclude the discussion on mechanisms affecting nanowire growth in template pores, we can say that there are various parameters that affect the nanowire growth such as cell design, pH of electrolyte, applied bias as well as various growth mechanisms like convection, migration and diffusion. However, of these, diffusion-controlled mass transfer flux is more important. Influence of migration can be ignored by lowering the concentration of depositing metal and increasing the concentration of supporting electrolyte. The contribution due to natural convection can be avoided by electrodeposition cell design in which working electrode can be placed in up or vertical positions.

Commercial self-organized porous AAO templates purchased from Synkera Technologies, Inc., have AAO pore length and diameter of ~50 μm and ~100 nm, respectively. **Figure 5** shows the

**Figure 5.** SEM image of commercial anodic aluminum oxide (AAO) template purchased from Synkera Technologies Inc. (a) SEM image of top surface of AAO template at 500-nm scale bar and inset shows 1-μm scale bar. These pores are nonuniform pores. (b) SEM image of bottom surface of AAO template at 500-nm scale bar.

SEM images of the (a) top and (b) bottom surfaces of commercial AAO templates. Inset shows higher magnification of the same surfaces. It is obvious from these images that the top and bottom pores do not have the same diameter, and the top surface pores are larger in diameter compared to the bottom pores. It is this difference in pore diameters that leads to most researchers depositing gold on the lower diameter side of the AAO template, thus preventing overfilling of the pores.

## **3. Fabrication and characterization of InSb nanowires in commercial AAO template**

Using the commercial AAO templates discussed in the previous section, one surface of the AAO template was coated with a thin film of Au (∼150 nm) using thermal evaporation. Prior to Au evaporation, the pores were widened nominally by soaking the template in 5% phosphoric acid (H3PO3) for 2 min at 30°C. The electrodeposition process was conducted in a three-electrode cell with the AAO template Au side as the working electrode, platinum electrode as the counter electrode and Ag/AgCl as the reference electrode. Crucial part of deposition process was to cover the edges of the AAO template with an insulating material and then attaching it to a conducting copper tape. The side of the AAO pore, which was covered with thin film of Au, was attached to conducting glue side of copper tape. After that, insulating polymer was applied to avoid any conducting path except through the open AAO pores which terminates at the Au film at the bottom of the pores. The entire electrodeposition process was controlled and carried out by a potentiostat (Princeton Applied Research, model: Potentiostat/Galvanostat 263A). The deposition parameters used for electrolyte in this experiment are as follows: 0.15 M InCl3 , 0.1 M SbCl<sup>3</sup> , 0.36 M C<sup>6</sup> H8 O7 H2 O, and C<sup>6</sup> H5 Na3O<sup>7</sup> . The pH value of the electrolyte was adjusted to 1.7 and electrodeposition was performed at room temperature (200°C). The citrate ions have been used as complexing agents that bring the deposition potential of In and Sb closer to maintain a binary growth during the deposition of the InSb nanowires. The complete electrodeposition process was performed for 50 min at a potential of −1.5 V with respect to the reference electrode (Ag/AgCl). Also, magnetic stirring was used to maintain the uniform concentration of electrolyte during growth of nanowires. On completion of the deposition process, the AAO template was carefully rinsed several times with DI water, and then AAO template was removed from copper tape for further processing.

To conclude the discussion on mechanisms affecting nanowire growth in template pores, we can say that there are various parameters that affect the nanowire growth such as cell design, pH of electrolyte, applied bias as well as various growth mechanisms like convection, migration and diffusion. However, of these, diffusion-controlled mass transfer flux is more important. Influence of migration can be ignored by lowering the concentration of depositing metal and increasing the concentration of supporting electrolyte. The contribution due to natural convection can be avoided by electrodeposition cell design in which working electrode can be

**Figure 4.** Schematic diagram convection in and around nanoporous pores in the configuration of anode over cathode.

Commercial self-organized porous AAO templates purchased from Synkera Technologies, Inc., have AAO pore length and diameter of ~50 μm and ~100 nm, respectively. **Figure 5** shows the

**Figure 5.** SEM image of commercial anodic aluminum oxide (AAO) template purchased from Synkera Technologies Inc. (a) SEM image of top surface of AAO template at 500-nm scale bar and inset shows 1-μm scale bar. These pores are

nonuniform pores. (b) SEM image of bottom surface of AAO template at 500-nm scale bar.

placed in up or vertical positions.

30 Semiconductors - Growth and Characterization

Following steps are involved in extraction of nanowires from AAO template: (a) AAO template was placed in a clean 80 ml beaker with the gold side facing upward and gold film was removed by 1 M potassium iodide (KI). To dissolve the AAM template, a few drops of 1 M KOH solution was used. This was done with vigorous shaking for about 20 min (b) following dissolution of the template; 15 ml of DI water was added to the solution, which was then sonicated for 10 min. The sonication step helps to dissolve the AAO completely, while breaking the InSb nanowires off the InSb crust (overgrown areas). (c) DI water as added to wash out the precipitates in the bottom of the beaker into the vial; the process was repeated until there was no visible precipitate in the beaker. This step was extremely important, since most of the nanowires were obtained from these precipitates. (d) To further dilute the KOH and AAO in the solution, both vials were placed in a centrifuge (Fisher Scientific Centrifuge, model: 228 Benchtop Centrifuge) at 3000 rpm for 10 min. Eighty percent of the solution was then carefully removed from the top of the centrifuge vial. Dilution step was repeated until there was no residue of KOH and AAO left, a fact that was verified by SEM. The final dark solution was found to contain a high density of InSb nanowires that could then be drop-cast for making devices. After growth, overgrown part is removed by a gentle mechanical polishing. Average length and diameter of InSb nanowire after processing have been found to be 20–30 μm and 100 nm, respectively.

**Figure 6(a)** is an SEM image of dispersed InSb nanowires, which show that they have a rough surface. This surface roughness is attributed to nonuniform pores and roughness inside the pores of the AAO template. This roughness can be avoided by using homemade AAO template or track-etched polycarbonate membranes where the tubular pores are uniform resulting in smooth and uniform nanowires as compared to those grown in commercially available AAO templates. The nanowire composition was verified by energy dispersive X-ray spectroscopy (EDX) in **Figure 6(b)**, which shows Sb-rich composition, with an average In:Sb weight ratio of 40:60.

To verify the composition and crystallinity, Raman spectroscopy and X-ray diffraction studies were made on the as-grown InSb nanowires. **Figure 7** shows the Raman spectrum obtained from two crossed InSb nanowires that were dispersed on a Si substrate. Optical image of the region where the spectrum was collected from is shown in the inset of **Figure 7**. Room temperature Raman spectrum shows two distinct peaks at 178 and 188 cm−1, which correspond to TO and LO phonon modes, respectively. These peaks are matched to prior studies on InSb nanowire [9]. Additional peaks at 150 (TO-TA) and 110 cm−1 have been reported earlier in InSb nanorods by Wada et al. [34].

Above InSb NWs was characterized with high-resolution X-ray diffraction system from Rigaku model Ultima III. **Figure 8** shows the X-ray diffraction spectrum of as-grown InSb nanowires in AAO template. All strong peaks were indexed to 2θ = 23.0, 39.0, and 46 identify

the dominant phase as zinc blende InSb. No other crystalline impurities peaks, such as In2O3, were detected in the XRD pattern. Similar XRD pattern has been reported earlier for electro-

**Figure 7.** Raman spectrum obtained from two crossed InSb nanowires; inset shows the region from where the Raman spectrum was collected. The two peaks at 150 and 110 cm−1 are most likely related to surface roughness and defects in the

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nanowire. The characteristic TO and LO peaks attest to crystallinity of the InSb nanowires [32].

**Figure 8.** XRD spectrum of InSb NWs grown in AAO template in the scan range of 2θ = 20–60°.

To obtain additional information on the electronic quality of the as-grown InSb nanowires, electron transport measurement was performed on individual nanowires by connecting them in a back-gated FET type structure as shown in **Figure 9**. To determine type, concentration and mobility of carriers in the nanowire, 2-terminal and 3-terminal current–voltage measurements were performed on the fabricated device. **Figure 9(a)** shows the device schematic where back-gate has been used and **Figure 9(b)** shows the SEM image of the InSb nanowire

chemically grown InSb nanowires [35–37].

contacted by Cr/Au as a source and drain.

**Figure 6.** SEM images of (a) InSb nanowires that were removed from the AAO template and dispersed on a cleaned Si substrate. Inset shows a single nanowire with roughened surface and (b) EDX spectrum of Sb-rich InSb nanowires [32].

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Benchtop Centrifuge) at 3000 rpm for 10 min. Eighty percent of the solution was then carefully removed from the top of the centrifuge vial. Dilution step was repeated until there was no residue of KOH and AAO left, a fact that was verified by SEM. The final dark solution was found to contain a high density of InSb nanowires that could then be drop-cast for making devices. After growth, overgrown part is removed by a gentle mechanical polishing. Average length and diameter of InSb nanowire after processing have been found to be 20–30 μm and

**Figure 6(a)** is an SEM image of dispersed InSb nanowires, which show that they have a rough surface. This surface roughness is attributed to nonuniform pores and roughness inside the pores of the AAO template. This roughness can be avoided by using homemade AAO template or track-etched polycarbonate membranes where the tubular pores are uniform resulting in smooth and uniform nanowires as compared to those grown in commercially available AAO templates. The nanowire composition was verified by energy dispersive X-ray spectroscopy (EDX) in **Figure 6(b)**, which shows Sb-rich composition, with an average In:Sb weight

To verify the composition and crystallinity, Raman spectroscopy and X-ray diffraction studies were made on the as-grown InSb nanowires. **Figure 7** shows the Raman spectrum obtained from two crossed InSb nanowires that were dispersed on a Si substrate. Optical image of the region where the spectrum was collected from is shown in the inset of **Figure 7**. Room temperature Raman spectrum shows two distinct peaks at 178 and 188 cm−1, which correspond to TO and LO phonon modes, respectively. These peaks are matched to prior studies on InSb nanowire [9]. Additional peaks at 150 (TO-TA) and 110 cm−1 have been reported earlier in InSb nanorods by

Above InSb NWs was characterized with high-resolution X-ray diffraction system from Rigaku model Ultima III. **Figure 8** shows the X-ray diffraction spectrum of as-grown InSb nanowires in AAO template. All strong peaks were indexed to 2θ = 23.0, 39.0, and 46 identify

**Figure 6.** SEM images of (a) InSb nanowires that were removed from the AAO template and dispersed on a cleaned Si substrate. Inset shows a single nanowire with roughened surface and (b) EDX spectrum of Sb-rich InSb nanowires [32].

100 nm, respectively.

32 Semiconductors - Growth and Characterization

ratio of 40:60.

Wada et al. [34].

**Figure 7.** Raman spectrum obtained from two crossed InSb nanowires; inset shows the region from where the Raman spectrum was collected. The two peaks at 150 and 110 cm−1 are most likely related to surface roughness and defects in the nanowire. The characteristic TO and LO peaks attest to crystallinity of the InSb nanowires [32].

**Figure 8.** XRD spectrum of InSb NWs grown in AAO template in the scan range of 2θ = 20–60°.

the dominant phase as zinc blende InSb. No other crystalline impurities peaks, such as In2O3, were detected in the XRD pattern. Similar XRD pattern has been reported earlier for electrochemically grown InSb nanowires [35–37].

To obtain additional information on the electronic quality of the as-grown InSb nanowires, electron transport measurement was performed on individual nanowires by connecting them in a back-gated FET type structure as shown in **Figure 9**. To determine type, concentration and mobility of carriers in the nanowire, 2-terminal and 3-terminal current–voltage measurements were performed on the fabricated device. **Figure 9(a)** shows the device schematic where back-gate has been used and **Figure 9(b)** shows the SEM image of the InSb nanowire contacted by Cr/Au as a source and drain.

**Figure 9.** (a) Schematic of back-gated InSb nanowire field effect transistor and (b) SEM image of a single InSb nanowire contacted by two Cr/Au contacts [32].

**Figure 10** shows the results of electron transport measurements made on the back-gated nanowire at different Vds (from 0.2 to 1.0 V, in steps of 0.2 V). The transfer characteristic at Vds = 1.0 V was found to have a subthreshold region between −0.5 and +0.5 V in both linear and log scale (Inset of **Figure 10**.). These values most likely correspond to complete depletion of the nanowire. The conductance of InSb nanowire has been shown to decrease with the increasing V<sup>g</sup> , that is, confirmation of a p-channel FET behavior. The p-type behavior is attributed to the Sb-rich nature of the InSb nanowire and attributed to the two common defects in Sb-rich InSb and Sb antisite and In interstitial defects [38].

The equilibrium hole concentration p0 in the nanowire was determined by using the Eq. (1) [39, 40]:

$$p\_0 = \frac{CV\_T}{q\pi R^2 L} \tag{1}$$

where transconductance *gm* = (

ratio of 103

crystalline nanowire.

*<sup>d</sup> <sup>I</sup>* \_\_\_\_*ds*

effect mobility was determined to be 507 cm<sup>2</sup> V−1 s−1 at Vds = 1.0 V. The lowest mobility has been found to be 277 cm2 V−1 s−1 at Vds = 0.2 V, which is higher than the previously reported [11] value of 57 cm<sup>2</sup> V−1 s−1 obtained at Vds = 0.1 V on unintentionally doped 5-μm long p-type InSb nanowire grown by electrochemical method. Similarly, carbon-doped p-type InSb nanowire of length 1.8 μm has been reported to be has the mobility of 127 ± 21 cm<sup>2</sup> V−1 s−1 at Vds = 0.05 V [41]. Inset of **Figure 10** shows p-type nanowire FET has been shown a relatively high ON/OFF

**Figure 10.** Transfer characteristics of InSb nanowire back-gated FET at increasing *V*ds from 0.2 to 1.0 V. Inset shows a logarithm plot of the *Ids vs. Vgs* curve at *Vds* = 1.0 V. The ON–OFF current ratio is estimated to be of the order of 10<sup>3</sup>

of the electrolyte. This most likely causes an increase in adsorption of Sb anions on the growing

**Figure 11.** (a) Temperature-dependent I-V shows current increasing with temperature, which is characteristic semiconducting behavior. (b) Temperature-dependent conductivity measurements (Arrhenius plot) for the extraction of activation energy from a single InSb nanowire. Inset shows normalized resistance versus temperature, which shows the

exponential decrease of resistance with increasing temperature [32].

. One possible cause for the Sb-rich nature of these nanowires is the lower pH (1.7)

*<sup>d</sup> Vgs*) has been deduced at different constant Vds. The field

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[32].

35

where the gate capacitance (C) has been determined from the Eq. (2) [40]:

$$\mathbf{C} = \frac{2\pi\varepsilon\_{o}\varepsilon\_{g'}L}{\cosh^{-1}\left(\frac{t\_{ai}+R}{R}\right)}\tag{2}$$

where ϵ*eff* ~ 2.2 is the effective dielectric constant of the SiO2 dielectric, ϵ<sup>0</sup> is the permittivity of free space, tox = 250 nm is the thickness of dielectric, R = 50 nm is the nanowire radius, and L = 10 μm is the length of nanowire channel, which is used for this measurement. Using Eqs. (1) and (2), the hole concentration has been calculated to be 1.96 × 1016 cm−3. From the linear part of the Ids versus Vgs as shown in **Figure 10**, the field effect mobility (μFE) was determined at different constant Vds using Eq. (3):

$$
\mu\_{\rm PE} = \left(\frac{L^2}{C V\_{\rm de}}\right) \times \left(\frac{dI\_{\rm de}}{dV\_{\rm p}}\right) \tag{3}
$$

Understanding the Mechanisms that Affect the Quality of Electrochemically Grown… http://dx.doi.org/10.5772/intechopen.71631 35

**Figure 10.** Transfer characteristics of InSb nanowire back-gated FET at increasing *V*ds from 0.2 to 1.0 V. Inset shows a logarithm plot of the *Ids vs. Vgs* curve at *Vds* = 1.0 V. The ON–OFF current ratio is estimated to be of the order of 10<sup>3</sup> [32].

**Figure 10** shows the results of electron transport measurements made on the back-gated nanowire at different Vds (from 0.2 to 1.0 V, in steps of 0.2 V). The transfer characteristic at Vds = 1.0 V was found to have a subthreshold region between −0.5 and +0.5 V in both linear and log scale (Inset of **Figure 10**.). These values most likely correspond to complete depletion of the nanowire. The conductance of InSb nanowire has been shown to decrease with the

**Figure 9.** (a) Schematic of back-gated InSb nanowire field effect transistor and (b) SEM image of a single InSb nanowire

uted to the Sb-rich nature of the InSb nanowire and attributed to the two common defects in

*cosh*<sup>−</sup><sup>1</sup> ( *t ox* <sup>+</sup> *<sup>R</sup>* \_\_\_\_ *<sup>R</sup>* )

free space, tox = 250 nm is the thickness of dielectric, R = 50 nm is the nanowire radius, and L = 10 μm is the length of nanowire channel, which is used for this measurement. Using Eqs. (1) and (2), the hole concentration has been calculated to be 1.96 × 1016 cm−3. From the linear part of the Ids versus Vgs as shown in **Figure 10**, the field effect mobility (μFE) was determined

*<sup>C</sup> Vds*) <sup>×</sup> (

*d I* \_\_\_\_*ds*

Sb-rich InSb and Sb antisite and In interstitial defects [38].

*<sup>p</sup>*<sup>0</sup> <sup>=</sup> *<sup>C</sup> <sup>V</sup>* \_\_\_\_\_\_*<sup>T</sup>*

*<sup>C</sup>* <sup>=</sup> <sup>2</sup>*<sup>π</sup> <sup>ε</sup>*<sup>0</sup> <sup>ϵ</sup>*eff <sup>L</sup>* \_\_\_\_\_\_\_\_\_\_\_

where the gate capacitance (C) has been determined from the Eq. (2) [40]:

where ϵ*eff* ~ 2.2 is the effective dielectric constant of the SiO2 dielectric, ϵ<sup>0</sup>

The equilibrium hole concentration p0

contacted by two Cr/Au contacts [32].

34 Semiconductors - Growth and Characterization

at different constant Vds using Eq. (3):

*μFE* <sup>=</sup> ( *<sup>L</sup>*<sup>2</sup> \_\_\_\_

, that is, confirmation of a p-channel FET behavior. The p-type behavior is attrib-

in the nanowire was determined by using the Eq. (1)

*<sup>q</sup> <sup>R</sup>*<sup>2</sup> *<sup>L</sup>* (1)

*<sup>d</sup> Vgs*) (3)

(2)

is the permittivity of

increasing V<sup>g</sup>

[39, 40]:

where transconductance *gm* = ( *<sup>d</sup> <sup>I</sup>* \_\_\_\_*ds <sup>d</sup> Vgs*) has been deduced at different constant Vds. The field effect mobility was determined to be 507 cm<sup>2</sup> V−1 s−1 at Vds = 1.0 V. The lowest mobility has been found to be 277 cm2 V−1 s−1 at Vds = 0.2 V, which is higher than the previously reported [11] value of 57 cm<sup>2</sup> V−1 s−1 obtained at Vds = 0.1 V on unintentionally doped 5-μm long p-type InSb nanowire grown by electrochemical method. Similarly, carbon-doped p-type InSb nanowire of length 1.8 μm has been reported to be has the mobility of 127 ± 21 cm<sup>2</sup> V−1 s−1 at Vds = 0.05 V [41]. Inset of **Figure 10** shows p-type nanowire FET has been shown a relatively high ON/OFF ratio of 103 . One possible cause for the Sb-rich nature of these nanowires is the lower pH (1.7) of the electrolyte. This most likely causes an increase in adsorption of Sb anions on the growing crystalline nanowire.

**Figure 11.** (a) Temperature-dependent I-V shows current increasing with temperature, which is characteristic semiconducting behavior. (b) Temperature-dependent conductivity measurements (Arrhenius plot) for the extraction of activation energy from a single InSb nanowire. Inset shows normalized resistance versus temperature, which shows the exponential decrease of resistance with increasing temperature [32].

The results of temperature-dependent I-V have been shown in **Figure 11(a)**. Linear I-V shows that the Cr/Au contacts (source-drain) to the nanowire are ohmic contacts. The InSb nanowire resistivity was determined to be 248 Ω cm−1, which is higher than the value reported for bulk InSb. The reasons for higher resistivity of the nanowire compared to bulk are as follows: (i) Sb antisite and In interstitials and (ii) surface roughness of nanowire (**Figure 6(a)**). The increased resistivity has been also reported in the references [15, 42], attributed to a significantly reduced hole mobility caused by the scattering at nanowire surface. The nanowire resistance was found to decrease with increasing temperature (inset of **Figure 11(b)**). This large temperature dependence of resistance of InSb nanowire has been reported previously in Ref.s [15, 43]; this is the characteristic behavior of a semiconductor where carrier concentration varies exponentially with temperature. The nanowire conductivity has been measured using device geometry. **Figure 11(b)** shows that below T = 200 K, the nanowire conductivity (σ) exhibits the characteristics of thermal activation which can be analyzed using Eq. (4):

$$
\sigma = \sigma\_0 \exp\left(\frac{-E\_s}{k\_y T}\right) \tag{4}
$$

After second step anodization, aluminum is dissolved in 1 M mercury chloride (H<sup>g</sup>

cal, magnetic, and thermoelectric properties of nanowires.

which shows that diameter of pore is approximately 50 nm.

rated solution. We have achieved ~50 μm pore length in about 4 h 16 min with pore diameter of ~50 nm. One advantage for the home-grown route of synthesizing templates is the possibility of obtaining smaller pore diameters. In this chapter, results of pore widths of ~50–70 nm are presented, in contrast to the larger diameter ~100 nm obtained from the commercial templates. Nanowires of different diameters will enable a study of its dependence on the electri-

**Figure 12.** SEM images of AAO template top surfaces (a) at scale bar of 500 nm and (b) with scale bar in pore diameter

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A noteworthy challenge in homemade AAO templates is the presence of a thin barrier layer comprising aluminum oxide. Following fabrication of the alumina layer and removal of the underlying metallic aluminum, the bottom layer of the template was found to exhibit bulges or protrusions. **Figure 13(a)** shows an SEM image showing the capped protrusions (hemispherical caps) of the barrier later at the bottom of the pore. Inset shows an enlarged image of the barrier caps. **Figure 16(b)** shows the completely opened barrier later after etching in 5% phosphoric acid at 30°C for 30min. It has also been observed that if etching time in 5% phosphoric acid is increased from 30min to 35 min, the thin membrane tears in sections from the top surface and the pores merge together as their walls collapse. Such membranes cannot be used for nanowire growth via electrodeposition, since the collapsed walls will result in direct contact with the underlying metal layer. **Figure 14** shows SEM images of bottom and top

**Figure 13.** SEM images of bottom side of AAO template: (a) hemispherical cap of barrier layer formed during anodization

and (b) complete barrier layer etching in 5% phosphoric acid at 30°C for 30min.

Cl2 ) satu37

where σ<sup>0</sup> is the pre-exponent factor, E<sup>a</sup> is the activation energy, kB is Boltzmann's constant (8.617 × 10−5 eVK−1), and T is the absolute temperature. E<sup>a</sup> has been estimated to be about 0.1 eV corresponding to carrier generation across the bandgap with activation energy equal to half of the bandgap. The roughness of the nanowire surface can be exploited to control phonon transport. Future work can be directed toward reducing the intrinsic point defects within the nanowire so that electrical conductivity can be enhanced, while the surface roughness can impede transport of phonons to reduce the lattice contribution to thermal conductivity, making it a promising material for thermoelectric applications.

#### **4. Homemade AAO template fabrication for semiconducting nanowire growth of desired dimensions**

Anodic films on aluminum have received considerable attention due to their extensive application as templates for synthesizing various nanostructures in the forms of nanowires and nanotubes [44–46]. To obtain a higher control of the nanowire dimensions and surface quality, an electrochemical self-assembly technique was used to fabricate a hexagonally ordered array of cylindrical nanopores on an aluminum substrate. Starting with high purity, unpolished and annealed aluminum (99.997%, AlfaAsar) foil with thickness of ~250 μm, the unpolished samples were first chemically polished [47] using 15 parts of 68% nitric acid and 85 parts of 85% phosphoric acid for 5 min at 850°C. The samples were then neutralized in 1 M sodium hydroxide (NaOH) for 20 min. This was followed by a multistep anodization process using 3% oxalic acid and 40 V DC at room temperature. Final step of anodization is carried out for 5 min, which produces pore length of ~1 μm [48]. **Figure 12** shows SEM images of top surface of AAO template after second anodization, **Figure 12(a)** is the SEM images of top surface at a scale bar of 500 nm and **Figure 12(b)** is the top surface with a scale bar which shows the pore diameter of AAO template is approximately 50 nm.

Understanding the Mechanisms that Affect the Quality of Electrochemically Grown… http://dx.doi.org/10.5772/intechopen.71631 37

The results of temperature-dependent I-V have been shown in **Figure 11(a)**. Linear I-V shows that the Cr/Au contacts (source-drain) to the nanowire are ohmic contacts. The InSb nanowire resistivity was determined to be 248 Ω cm−1, which is higher than the value reported for bulk InSb. The reasons for higher resistivity of the nanowire compared to bulk are as follows: (i) Sb antisite and In interstitials and (ii) surface roughness of nanowire (**Figure 6(a)**). The increased resistivity has been also reported in the references [15, 42], attributed to a significantly reduced hole mobility caused by the scattering at nanowire surface. The nanowire resistance was found to decrease with increasing temperature (inset of **Figure 11(b)**). This large temperature dependence of resistance of InSb nanowire has been reported previously in Ref.s [15, 43]; this is the characteristic behavior of a semiconductor where carrier concentration varies exponentially with temperature. The nanowire conductivity has been measured using device geometry. **Figure 11(b)** shows that below T = 200 K, the nanowire conductivity (σ) exhibits the characteristics of thermal activation which can be analyzed using Eq. (4):

<sup>−</sup>*E*\_\_\_*<sup>a</sup>*

0.1 eV corresponding to carrier generation across the bandgap with activation energy equal to half of the bandgap. The roughness of the nanowire surface can be exploited to control phonon transport. Future work can be directed toward reducing the intrinsic point defects within the nanowire so that electrical conductivity can be enhanced, while the surface roughness can impede transport of phonons to reduce the lattice contribution to thermal conductivity, mak-

**4. Homemade AAO template fabrication for semiconducting nanowire** 

Anodic films on aluminum have received considerable attention due to their extensive application as templates for synthesizing various nanostructures in the forms of nanowires and nanotubes [44–46]. To obtain a higher control of the nanowire dimensions and surface quality, an electrochemical self-assembly technique was used to fabricate a hexagonally ordered array of cylindrical nanopores on an aluminum substrate. Starting with high purity, unpolished and annealed aluminum (99.997%, AlfaAsar) foil with thickness of ~250 μm, the unpolished samples were first chemically polished [47] using 15 parts of 68% nitric acid and 85 parts of 85% phosphoric acid for 5 min at 850°C. The samples were then neutralized in 1 M sodium hydroxide (NaOH) for 20 min. This was followed by a multistep anodization process using 3% oxalic acid and 40 V DC at room temperature. Final step of anodization is carried out for 5 min, which produces pore length of ~1 μm [48]. **Figure 12** shows SEM images of top surface of AAO template after second anodization, **Figure 12(a)** is the SEM images of top surface at a scale bar of 500 nm and **Figure 12(b)** is the top surface with a scale bar which shows the pore

*kB <sup>T</sup>*) (4)

has been estimated to be about

is the activation energy, kB is Boltzmann's constant

*σ* = *σ*<sup>0</sup> *exp*(

36 Semiconductors - Growth and Characterization

is the pre-exponent factor, E<sup>a</sup>

**growth of desired dimensions**

(8.617 × 10−5 eVK−1), and T is the absolute temperature. E<sup>a</sup>

ing it a promising material for thermoelectric applications.

diameter of AAO template is approximately 50 nm.

where σ<sup>0</sup>

**Figure 12.** SEM images of AAO template top surfaces (a) at scale bar of 500 nm and (b) with scale bar in pore diameter which shows that diameter of pore is approximately 50 nm.

After second step anodization, aluminum is dissolved in 1 M mercury chloride (H<sup>g</sup> Cl2 ) saturated solution. We have achieved ~50 μm pore length in about 4 h 16 min with pore diameter of ~50 nm. One advantage for the home-grown route of synthesizing templates is the possibility of obtaining smaller pore diameters. In this chapter, results of pore widths of ~50–70 nm are presented, in contrast to the larger diameter ~100 nm obtained from the commercial templates. Nanowires of different diameters will enable a study of its dependence on the electrical, magnetic, and thermoelectric properties of nanowires.

A noteworthy challenge in homemade AAO templates is the presence of a thin barrier layer comprising aluminum oxide. Following fabrication of the alumina layer and removal of the underlying metallic aluminum, the bottom layer of the template was found to exhibit bulges or protrusions. **Figure 13(a)** shows an SEM image showing the capped protrusions (hemispherical caps) of the barrier later at the bottom of the pore. Inset shows an enlarged image of the barrier caps. **Figure 16(b)** shows the completely opened barrier later after etching in 5% phosphoric acid at 30°C for 30min. It has also been observed that if etching time in 5% phosphoric acid is increased from 30min to 35 min, the thin membrane tears in sections from the top surface and the pores merge together as their walls collapse. Such membranes cannot be used for nanowire growth via electrodeposition, since the collapsed walls will result in direct contact with the underlying metal layer. **Figure 14** shows SEM images of bottom and top

**Figure 13.** SEM images of bottom side of AAO template: (a) hemispherical cap of barrier layer formed during anodization and (b) complete barrier layer etching in 5% phosphoric acid at 30°C for 30min.

**Figure 14.** SEM images of bottom and top surfaces after etching in 5% phosphoric acid (H<sup>3</sup> PO<sup>4</sup> ) for 35 min at 30°C. (a) SEM image of bottom surface, which shows all the barrier layers are etched and pores are uniformly open and (b) SEM image of the same sample as in (a), which shows over etching of top surface which looks like pores are bundled together*.*

comparison. **Figure 16** shows the SEM images of AAO top surface (a) as-grown and without etching, (b) after 17 min etching in 1 M phosphoric acid at 30°C and (c) after 20 min etching. It is clear from SEM image in **Figure 16(b)** that after 17 min of etching in 1 M phosphoric acid, the AAO template top pores are widened to its maximum capacity, and any increase in etching time will result in merging of the pores. **Figure 16(c)** shows that 20 min of etching is over etching and the pores are merged. After appropriate time etching, barrier layer is thinned and pore diameter is widened. Now, pores can be filled with cobalt by AC electro-

**Figure 16.** SEM images of top surface of AAO template: (a) shows nonporous hexagonal AAO pores without any etching, (b) after 17 min etching in 1 M phosphoric acid at 30°C and, (c) after 20 min etching in 1 M phosphoric acid at 30°C.

Barrier layer thinning is an alternative technique to remove the barrier layer formed at the end of second anodization process. After thinning of barrier layer, individual pores terminate in the metallic aluminum layer at the bottom of the template. Such membranes grown directly on aluminum foil are suitable for electrodeposition since the bottom aluminum layer works

As discussed in Section 3, nanowire growth can also be carried out in the pores of a track-

Following the pioneering work of pore creation in track-etched mica by Possin [51], metallic (Ag) nanowires were grown in 8 nm wide pores by Williams and Giordano [52]. Penner and Martin [53–55] subsequently created pores in polycarbonate membranes by track-etch method. In this work, polycarbonate membranes were purchased from Whatman and had pore lengths of 20 μm and diameter of the order of 200 nm. Following metal deposition on one surface of this template, InSb nanowires were grown in its pores. Unlike the pores in the AAO template, the polycarbonate pores were of uniform diameter and had smooth surfaces. This results in a very

As-grown InSb nanowires in polycarbonate were first dissolved in dichloromethane and then in chloroform. After InSb nanowires were also grown in track-etched polycarbonate membranes and an SEM image of the as-grown nanowires is shown in **Figure 17**. The biggest

**5. Semiconducting InSb nanowires grown in polycarbonate template**

O solution stabilized with 2% H<sup>3</sup>

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39

BO<sup>3</sup>

at 20 V AC

.7H<sup>2</sup>

chemical deposition using a 5% CoSO<sup>4</sup>

etched polycarbonate membrane.

smooth surface for the as-grown InSb nanowires.

and 250 Hz [50].

as one electrode.

surfaces after etching in 5% phosphoric acid (H<sup>3</sup> PO<sup>4</sup> ) for 35 min at 30°C. **Figure 15(a)** shows SEM image of bottom surface, which shows all the barrier layers are etched and pores are uniformly open and **Figure 15(b)** shows SEM image of the same sample as shown in **Figure 13**, which shows over etching of top surface which looks like pores are bundled together.

To check the anodization rate, cross-sectional image of AAO template has been taken at two different times of second anodization, 5 and 250 min, respectively. **Figure 15** shows SEM images of cross-section of homemade templates. **Figure 15(a)** shows ~50-μm long pore which was obtained after 250 min of second anodization, and inset shows smooth and straight pore and; **Figure 15(b)** shows ~ 5-μm long pore which was obtained after 15 min of second anodization. Based on **Figure 15**, anodization rate in the abovementioned condition for the growth of homemade AAO template is 1-μm long in 5 min, which is a faster rate than any other anodization conditions reported before [49] .

As mentioned earlier, removal of the barrier layer from the bottom of the AAO membrane is very challenging; if the etching time is not well controlled in most cases, it results in collapse of the tubular walls and destruction of the pores from the top surface of the template. Results of barrier removal using 1 M H<sup>3</sup> PO<sup>4</sup> at 30°C at different times are presented for

**Figure 15.** SEM images of cross-section of homemade template: (a) shows ~50-μm long pore which was obtained after 250 min of second anodization, and inset shows smooth and straight pore and (b) shows ~5-μm long pore which was obtained after 15 min of second anodization.

Understanding the Mechanisms that Affect the Quality of Electrochemically Grown… http://dx.doi.org/10.5772/intechopen.71631 39

**Figure 16.** SEM images of top surface of AAO template: (a) shows nonporous hexagonal AAO pores without any etching, (b) after 17 min etching in 1 M phosphoric acid at 30°C and, (c) after 20 min etching in 1 M phosphoric acid at 30°C.

comparison. **Figure 16** shows the SEM images of AAO top surface (a) as-grown and without etching, (b) after 17 min etching in 1 M phosphoric acid at 30°C and (c) after 20 min etching. It is clear from SEM image in **Figure 16(b)** that after 17 min of etching in 1 M phosphoric acid, the AAO template top pores are widened to its maximum capacity, and any increase in etching time will result in merging of the pores. **Figure 16(c)** shows that 20 min of etching is over etching and the pores are merged. After appropriate time etching, barrier layer is thinned and pore diameter is widened. Now, pores can be filled with cobalt by AC electrochemical deposition using a 5% CoSO<sup>4</sup> .7H<sup>2</sup> O solution stabilized with 2% H<sup>3</sup> BO<sup>3</sup> at 20 V AC and 250 Hz [50].

surfaces after etching in 5% phosphoric acid (H<sup>3</sup>

38 Semiconductors - Growth and Characterization

together*.*

ization conditions reported before [49] .

Results of barrier removal using 1 M H<sup>3</sup>

obtained after 15 min of second anodization.

PO<sup>4</sup>

SEM image of bottom surface, which shows all the barrier layers are etched and pores are uniformly open and **Figure 15(b)** shows SEM image of the same sample as shown in **Figure 13**,

(a) SEM image of bottom surface, which shows all the barrier layers are etched and pores are uniformly open and (b) SEM image of the same sample as in (a), which shows over etching of top surface which looks like pores are bundled

To check the anodization rate, cross-sectional image of AAO template has been taken at two different times of second anodization, 5 and 250 min, respectively. **Figure 15** shows SEM images of cross-section of homemade templates. **Figure 15(a)** shows ~50-μm long pore which was obtained after 250 min of second anodization, and inset shows smooth and straight pore and; **Figure 15(b)** shows ~ 5-μm long pore which was obtained after 15 min of second anodization. Based on **Figure 15**, anodization rate in the abovementioned condition for the growth of homemade AAO template is 1-μm long in 5 min, which is a faster rate than any other anod-

As mentioned earlier, removal of the barrier layer from the bottom of the AAO membrane is very challenging; if the etching time is not well controlled in most cases, it results in collapse of the tubular walls and destruction of the pores from the top surface of the template.

**Figure 15.** SEM images of cross-section of homemade template: (a) shows ~50-μm long pore which was obtained after 250 min of second anodization, and inset shows smooth and straight pore and (b) shows ~5-μm long pore which was

PO<sup>4</sup>

which shows over etching of top surface which looks like pores are bundled together.

**Figure 14.** SEM images of bottom and top surfaces after etching in 5% phosphoric acid (H<sup>3</sup>

) for 35 min at 30°C. **Figure 15(a)** shows

PO<sup>4</sup>

) for 35 min at 30°C.

at 30°C at different times are presented for

Barrier layer thinning is an alternative technique to remove the barrier layer formed at the end of second anodization process. After thinning of barrier layer, individual pores terminate in the metallic aluminum layer at the bottom of the template. Such membranes grown directly on aluminum foil are suitable for electrodeposition since the bottom aluminum layer works as one electrode.

As discussed in Section 3, nanowire growth can also be carried out in the pores of a tracketched polycarbonate membrane.

### **5. Semiconducting InSb nanowires grown in polycarbonate template**

Following the pioneering work of pore creation in track-etched mica by Possin [51], metallic (Ag) nanowires were grown in 8 nm wide pores by Williams and Giordano [52]. Penner and Martin [53–55] subsequently created pores in polycarbonate membranes by track-etch method. In this work, polycarbonate membranes were purchased from Whatman and had pore lengths of 20 μm and diameter of the order of 200 nm. Following metal deposition on one surface of this template, InSb nanowires were grown in its pores. Unlike the pores in the AAO template, the polycarbonate pores were of uniform diameter and had smooth surfaces. This results in a very smooth surface for the as-grown InSb nanowires.

As-grown InSb nanowires in polycarbonate were first dissolved in dichloromethane and then in chloroform. After InSb nanowires were also grown in track-etched polycarbonate membranes and an SEM image of the as-grown nanowires is shown in **Figure 17**. The biggest

Major challenge is the pore wettability, if this step is not optimized prior to nanowire growth, then even if all other conditions are met, this will result in either no growth or in nonuniform growth in some pores and formation of a mushroom-like crust. To avoid this, it is important that a few basic steps to be performed prior to nanowire growth. Pore-wetting can be achieved by: (i) sonication of nanoporous template in water to remove any air bubbles at the bottom of the pores and (ii) aeration of nanoporous template in vacuum chamber to remove any air bubbles at the bottom of the template. Another challenge is the nonuniform growth which can be avoided by

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41

maintaining a temperature gradient between working electrode and mouth of the pore.

The growth process of semiconducting nanowires in template pores is a relatively simple process that can be used for synthesizing a high density of well-ordered nanowires in an array. However, the growth process involves several complex basic mechanisms that are affected by cell design, pore wettability, and electrolyte condition. This chapter presents a brief review of work done in this area. It has been experimentally found that commercial AAO template pores have rough tubular surfaces and this results in nanowires with jagged edges. Most commercial AAO templates with narrow pores (of the order of 100 nm) are sometimes tapered and therefore, the metal evaporation and pore-wetting of the pores in these templates are very critical. To circumvent this problem, homemade templates with controlled thickness and pore diameters can be used. However, in this case, the challenge is with the barrier layer formed on the backside of the template. A controlled etching process typically achieves barrier layer removal, where the etching time and etchant concentration as well as temperature are controlled. If the process is not optimized, this will cause the template to tear along the top surface. The tear does not extend over the entire template thickness. The advantage of homemade templates is the pores are relatively uniform and therefore, the nanowires grown in these pores have smooth surfaces. Such smooth nanowires can also be grown in track-etched polycarbonate membranes.

**6. Conclusions**

**Author details**

**References**

Aug;**8**(8):648-653

Abhay Singh\* and Usha Philipose

\*Address all correspondence to: abhaysingh@my.unt.edu

Department of Physics, University of North Texas, Denton, TX, USA

[1] Fan Z, Razavi H, Do J, Moriwaki A, Ergen O, Chueh Y-L, et al. Three-dimensional nanopillar-array photovoltaics on low-cost and flexible substrates. Nature Materials. 2009

**Figure 17.** SEM image of a piece of InSb nanowire bundle which was drop-casted on the silicon wafer after dissolving as-grown InSb in polycarbonate template.

challenge in the polycarbonate membrane growth process is the clumping of the nanowires into bundles. This is most likely caused by the residues of the polymer membrane which tends to hold the nanowires together. Dissolution of the membrane is typically done using dichloromethane and chloroform, followed by cleaning in alcohol and DI water, and the nanowires grown by this technique were found to have very smooth surfaces and a length of approximately 20 μm. Electrolyte concentration and cell design were same as used to grow the InSb nanowire in AAO template as discussed in Section 3 except for the pH of solution was maintained at 1.9 instead of 1.7.

**Figure 18(a)** and **(b)** shows the SEM images of dispersed on silicon substrate at different places and **Figure 18(c)** shows the EDX spectrum of the InSb nanowires and they are In-rich.

**Figure 18.** SEM image of InSb nanowire after drop casting on silicon wafer after ultrasonication for 1 min and 1 min oxygen plasma etching: (a) showing broken nanowire due to ultrasonication which resulted in reduced lengths of nanowire from ~20 micron to average 5 micron and (b) showing reduced bundle with approximately 20 μm long InSb nanowires. Average diameter of the nanowires are approximately 150 nm and (c) energy dispersive X-ray (EDX) spectrum of in-rich (54.4 wt% In and 45.6 wt% Sb) electrochemically InSb nanowires grown in polycarbonate template with −1.5 V potential at room temperature and pH = 1.9.

Major challenge is the pore wettability, if this step is not optimized prior to nanowire growth, then even if all other conditions are met, this will result in either no growth or in nonuniform growth in some pores and formation of a mushroom-like crust. To avoid this, it is important that a few basic steps to be performed prior to nanowire growth. Pore-wetting can be achieved by: (i) sonication of nanoporous template in water to remove any air bubbles at the bottom of the pores and (ii) aeration of nanoporous template in vacuum chamber to remove any air bubbles at the bottom of the template. Another challenge is the nonuniform growth which can be avoided by maintaining a temperature gradient between working electrode and mouth of the pore.

#### **6. Conclusions**

challenge in the polycarbonate membrane growth process is the clumping of the nanowires into bundles. This is most likely caused by the residues of the polymer membrane which tends to hold the nanowires together. Dissolution of the membrane is typically done using dichloromethane and chloroform, followed by cleaning in alcohol and DI water, and the nanowires grown by this technique were found to have very smooth surfaces and a length of approximately 20 μm. Electrolyte concentration and cell design were same as used to grow the InSb nanowire in AAO template as discussed in Section 3 except for the pH of solution

**Figure 17.** SEM image of a piece of InSb nanowire bundle which was drop-casted on the silicon wafer after dissolving

**Figure 18(a)** and **(b)** shows the SEM images of dispersed on silicon substrate at different places and **Figure 18(c)** shows the EDX spectrum of the InSb nanowires and they are In-rich.

**Figure 18.** SEM image of InSb nanowire after drop casting on silicon wafer after ultrasonication for 1 min and 1 min oxygen plasma etching: (a) showing broken nanowire due to ultrasonication which resulted in reduced lengths of nanowire from ~20 micron to average 5 micron and (b) showing reduced bundle with approximately 20 μm long InSb nanowires. Average diameter of the nanowires are approximately 150 nm and (c) energy dispersive X-ray (EDX) spectrum of in-rich (54.4 wt% In and 45.6 wt% Sb) electrochemically InSb nanowires grown in polycarbonate template

was maintained at 1.9 instead of 1.7.

as-grown InSb in polycarbonate template.

40 Semiconductors - Growth and Characterization

with −1.5 V potential at room temperature and pH = 1.9.

The growth process of semiconducting nanowires in template pores is a relatively simple process that can be used for synthesizing a high density of well-ordered nanowires in an array. However, the growth process involves several complex basic mechanisms that are affected by cell design, pore wettability, and electrolyte condition. This chapter presents a brief review of work done in this area. It has been experimentally found that commercial AAO template pores have rough tubular surfaces and this results in nanowires with jagged edges. Most commercial AAO templates with narrow pores (of the order of 100 nm) are sometimes tapered and therefore, the metal evaporation and pore-wetting of the pores in these templates are very critical. To circumvent this problem, homemade templates with controlled thickness and pore diameters can be used. However, in this case, the challenge is with the barrier layer formed on the backside of the template. A controlled etching process typically achieves barrier layer removal, where the etching time and etchant concentration as well as temperature are controlled. If the process is not optimized, this will cause the template to tear along the top surface. The tear does not extend over the entire template thickness. The advantage of homemade templates is the pores are relatively uniform and therefore, the nanowires grown in these pores have smooth surfaces. Such smooth nanowires can also be grown in track-etched polycarbonate membranes.

#### **Author details**

Abhay Singh\* and Usha Philipose

\*Address all correspondence to: abhaysingh@my.unt.edu

Department of Physics, University of North Texas, Denton, TX, USA

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**Chapter 3**

Provisional chapter

**Semiconductor Quantum Wells with BenDaniel-Duke**

DOI: 10.5772/intechopen.73837

The energy levels of bound states of an electron in a quantum well with BenDaniel-Duke boundary condition are studied. Analytic, explicit, simple, and accurate formulae have been obtained for the ground state and the first excited state. In our approach, the exact, transcendental eigenvalues equations were replaced with approximate, tractable, algebraic equations, using algebraic approximations for some trigonometric functions. Our method can be applied to both type I and type II semiconductors and easily extended to quantum dots. The same approach was used for the semi-quantitative analyze of two toy

Keywords: type I and type II semiconductors, BenDaniel-Duke boundary conditions,

In the last three decades, nanophysics became a domain of increasing interest and intense research, due to the huge number of new effects produced at nanoscale level, in quantum wells (QWs), quantum dots (QDs), Janus nanoparticles, etc. These new effects are fascinating from the perspective of both applied and theoretical physics. The semiconductors provide the largest area of challenging subjects, due to their applications in nanoelectronic devices, multifunctional catalysis, (bio-)chemical sensors, data storage, solar energy con-

An attractive aspect of nanophysics is the fact that a quite large number of interesting problems can be approached using quite simple theoretical tools, sometimes at the level of oneparticle quantum mechanics. In some cases, the properties of nanostructures like quantum

> © 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 eproduction in any medium, provided the original work is properly cited.

> © 2018 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.

Semiconductor Quantum Wells with BenDaniel-Duke

**Boundary Conditions and Janus Nanorods**

Boundary Conditions and Janus Nanorods

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/intechopen.73837

models of Janus nanorods.

Janus nanorods, toy models

Victor Barsan

Victor Barsan

Abstract

1. Introduction

version, etc.

#### **Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods** Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

DOI: 10.5772/intechopen.73837

#### Victor Barsan Victor Barsan

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/intechopen.73837

#### Abstract

The energy levels of bound states of an electron in a quantum well with BenDaniel-Duke boundary condition are studied. Analytic, explicit, simple, and accurate formulae have been obtained for the ground state and the first excited state. In our approach, the exact, transcendental eigenvalues equations were replaced with approximate, tractable, algebraic equations, using algebraic approximations for some trigonometric functions. Our method can be applied to both type I and type II semiconductors and easily extended to quantum dots. The same approach was used for the semi-quantitative analyze of two toy models of Janus nanorods.

Keywords: type I and type II semiconductors, BenDaniel-Duke boundary conditions, Janus nanorods, toy models

#### 1. Introduction

In the last three decades, nanophysics became a domain of increasing interest and intense research, due to the huge number of new effects produced at nanoscale level, in quantum wells (QWs), quantum dots (QDs), Janus nanoparticles, etc. These new effects are fascinating from the perspective of both applied and theoretical physics. The semiconductors provide the largest area of challenging subjects, due to their applications in nanoelectronic devices, multifunctional catalysis, (bio-)chemical sensors, data storage, solar energy conversion, etc.

An attractive aspect of nanophysics is the fact that a quite large number of interesting problems can be approached using quite simple theoretical tools, sometimes at the level of oneparticle quantum mechanics. In some cases, the properties of nanostructures like quantum

wells, quantum dots, or quantum rods can be explained by just solving the Schrodinger equation with simple potentials. For instance, the basic physical properties of a heterostructure consisting of a thin layer of a semiconductor A sandwiched between two somewhat larger semiconductors of identical composition, B can be obtained from the study of the movement of a particle with position-dependent mass (PDM) in a finite square well. This particle is, of course, a charge carrier in the semiconductor, and—in our case—will be an electron. As the effective mass of a charge carrier in a semiconductor depends on the charge carrier-lattice interaction, it changes if the lattice composition or the symmetry changes. So, excepting the case of a charge moving in a perfect crystal, the effective mass of an electron or hole is, rigorously speaking, position dependent.

possible to approximate the trigonometric functions with algebraic expression, and to transform, in this way, the exact transcendental equation into an approximate algebraic one. In its simplest form, for instance, in approximations like sin x ≃x, for x ≪ 1, this "algebraization" is largely used. But what is really interesting is to use algebraic approximations of the trigonometric functions valid on their whole domain of definition, as de Alcantara Bonfin and Griffiths proposed in a recent paper [8]; such analytical approximations have been studied

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

http://dx.doi.org/10.5772/intechopen.73837

49

In this chapter, we shall obtain approximate analytical results for the energy of electronic bound states in quantum wells and in simple models of Janus semiconductor nanorods. As the concept of Janus nanoparticle is less popular than the concept of QWs or QDs, we shall give

Their name derives from the Roman god Janus: his head had two opposite faces. A Janus nanodot can be a sphere composed of two semispheres of different materials. A Janus nanorod can be a nanorod having the left half and the right half made of different materials. Due to their intrinsic duality, the opposite parts of Janus particles can be functionalized differently [10]. Janus particles with an electron-donor and -acceptor side may be used in photovoltaics. As the Janus nanoparticles have lower symmetry than their homogenous counterparts, their theoretical description is more difficult. In this chapter, we shall propose toy models for semiconduc-

The structure of this chapter is the following. We shall firstly formulate the basic theory for the quantum mechanical problem of a quantum well, composed of a thin semiconductor sandwiched between two massive ones. This heterojunction can de-modeled by a quantum well (QW), essentially a finite square well, with BenDaniel-Duke boundary conditions. Such a problem was recently discussed by several authors, like Singh et al. [11, 12], who replaced the trigonometric functions entering in the transcendental equations for the bound states energy by the first few terms of their series expansion; in this way, the equations become simple, tractable algebraic ones. Our approach is different, being based on a more sophisticate "algebraization" of trigonometric functions, as proposed by de Alcantara Bonfim and Griffiths [8]. We shall obtain explicit formulas (series expansions) for the ground state energy and for the first excited state, very accurate if the well is not too shallow. Our results can be

In the last part of our chapter, we shall study two toy models for semiconductor Janus nanorods; for the simplest one, we shall obtain analytical expressions for some energy eigen-

We shall solve the Schrodinger equation for an electron moving in a square well, described by

and extended by other authors [9].

here some short explanations.

tor Janus nanorods.

applied to both type I and type II semiconductors.

values of electronic bound states.

2. Basic theory

the potential:

But, simple as the theoretical tools needed for its investigation are, this problem of quantum mechanics involves two important issues: the position-dependent mass (PDM) quantum physics and semiconductor heterostructures. Let us shortly comment on these points.

The roots of the position-dependent effective mass concept are to be found in the pioneering works of Wannier (1937) and Slater (1949) (see Ref. 1 in [1]). Recent papers give explicit methods to obtain explicit solutions of the Schrödinger equation with PDM, for various forms of this dependence and for several classes of potentials [2–4].

However, in practical situations usually encountered in the physics of semiconductor junctions of two materials, A and B, the simplest and more popular form of position dependence of the effective mass is a step function: the effective mass has a constant value in the material A and another, constant value, in the material B: In such a case, the most convenient approach for obtaining the wave functions or the envelope functions in a heterostructure—for instance, a quantum well (QW) or quantum dot (QD)—is to solve the Schrödinger equation with BenDaniel-Duke boundary conditions for the wave function [5, 6].

The transition from the complex problem of a real semiconductor (for instance, Kane theory) to the simple problem of a particle moving in a square well with BenDaniel-Duke boundary conditions is indicated, for instance, in Chapter III of Bastard's book [5]. This simple problem provides, however, a realistic description of states near the high-symmetry points in the Brillouin zone of a large class of semiconductors. "It [i.e., 'the simple problem'] often leads to analytical results and leaves the user with the feeling that he can trace back, in a relatively transparent way, the physical origin of the numerical results." ([5], p. 63).

The boundary conditions for the wave functions or envelope functions at interfaces generate the eigenvalue equations for energy; of course, different boundary conditions generate different eigenvalue equations. They are transcendental equations, involving algebraic, trigonometric, hyperbolic, or even more complicated functions. With few exceptions (for instance, the Lambert equation [7]), their solutions, which cannot be expressed as a finite combinations of elementary functions, are not systematically studied.

However, in some situations, quite accurate analytical approximate solutions can be obtained. When a transcendental equation mixes algebraic and trigonometric functions, it might be possible to approximate the trigonometric functions with algebraic expression, and to transform, in this way, the exact transcendental equation into an approximate algebraic one. In its simplest form, for instance, in approximations like sin x ≃x, for x ≪ 1, this "algebraization" is largely used. But what is really interesting is to use algebraic approximations of the trigonometric functions valid on their whole domain of definition, as de Alcantara Bonfin and Griffiths proposed in a recent paper [8]; such analytical approximations have been studied and extended by other authors [9].

In this chapter, we shall obtain approximate analytical results for the energy of electronic bound states in quantum wells and in simple models of Janus semiconductor nanorods. As the concept of Janus nanoparticle is less popular than the concept of QWs or QDs, we shall give here some short explanations.

Their name derives from the Roman god Janus: his head had two opposite faces. A Janus nanodot can be a sphere composed of two semispheres of different materials. A Janus nanorod can be a nanorod having the left half and the right half made of different materials. Due to their intrinsic duality, the opposite parts of Janus particles can be functionalized differently [10]. Janus particles with an electron-donor and -acceptor side may be used in photovoltaics. As the Janus nanoparticles have lower symmetry than their homogenous counterparts, their theoretical description is more difficult. In this chapter, we shall propose toy models for semiconductor Janus nanorods.

The structure of this chapter is the following. We shall firstly formulate the basic theory for the quantum mechanical problem of a quantum well, composed of a thin semiconductor sandwiched between two massive ones. This heterojunction can de-modeled by a quantum well (QW), essentially a finite square well, with BenDaniel-Duke boundary conditions. Such a problem was recently discussed by several authors, like Singh et al. [11, 12], who replaced the trigonometric functions entering in the transcendental equations for the bound states energy by the first few terms of their series expansion; in this way, the equations become simple, tractable algebraic ones. Our approach is different, being based on a more sophisticate "algebraization" of trigonometric functions, as proposed by de Alcantara Bonfim and Griffiths [8]. We shall obtain explicit formulas (series expansions) for the ground state energy and for the first excited state, very accurate if the well is not too shallow. Our results can be applied to both type I and type II semiconductors.

In the last part of our chapter, we shall study two toy models for semiconductor Janus nanorods; for the simplest one, we shall obtain analytical expressions for some energy eigenvalues of electronic bound states.

### 2. Basic theory

wells, quantum dots, or quantum rods can be explained by just solving the Schrodinger equation with simple potentials. For instance, the basic physical properties of a heterostructure consisting of a thin layer of a semiconductor A sandwiched between two somewhat larger semiconductors of identical composition, B can be obtained from the study of the movement of a particle with position-dependent mass (PDM) in a finite square well. This particle is, of course, a charge carrier in the semiconductor, and—in our case—will be an electron. As the effective mass of a charge carrier in a semiconductor depends on the charge carrier-lattice interaction, it changes if the lattice composition or the symmetry changes. So, excepting the case of a charge moving in a perfect crystal, the effective mass of an electron or hole is,

But, simple as the theoretical tools needed for its investigation are, this problem of quantum mechanics involves two important issues: the position-dependent mass (PDM) quantum physics and semiconductor heterostructures. Let us shortly comment on these points.

The roots of the position-dependent effective mass concept are to be found in the pioneering works of Wannier (1937) and Slater (1949) (see Ref. 1 in [1]). Recent papers give explicit methods to obtain explicit solutions of the Schrödinger equation with PDM, for various forms

However, in practical situations usually encountered in the physics of semiconductor junctions of two materials, A and B, the simplest and more popular form of position dependence of the effective mass is a step function: the effective mass has a constant value in the material A and another, constant value, in the material B: In such a case, the most convenient approach for obtaining the wave functions or the envelope functions in a heterostructure—for instance, a quantum well (QW) or quantum dot (QD)—is to solve the Schrödinger equation with

The transition from the complex problem of a real semiconductor (for instance, Kane theory) to the simple problem of a particle moving in a square well with BenDaniel-Duke boundary conditions is indicated, for instance, in Chapter III of Bastard's book [5]. This simple problem provides, however, a realistic description of states near the high-symmetry points in the Brillouin zone of a large class of semiconductors. "It [i.e., 'the simple problem'] often leads to analytical results and leaves the user with the feeling that he can trace back, in a relatively transparent way, the physical origin of the numerical results." ([5],

The boundary conditions for the wave functions or envelope functions at interfaces generate the eigenvalue equations for energy; of course, different boundary conditions generate different eigenvalue equations. They are transcendental equations, involving algebraic, trigonometric, hyperbolic, or even more complicated functions. With few exceptions (for instance, the Lambert equation [7]), their solutions, which cannot be expressed as a finite combinations of

However, in some situations, quite accurate analytical approximate solutions can be obtained. When a transcendental equation mixes algebraic and trigonometric functions, it might be

rigorously speaking, position dependent.

48 Semiconductors - Growth and Characterization

p. 63).

of this dependence and for several classes of potentials [2–4].

BenDaniel-Duke boundary conditions for the wave function [5, 6].

elementary functions, are not systematically studied.

We shall solve the Schrodinger equation for an electron moving in a square well, described by the potential:

$$V(\mathbf{x}) = \begin{cases} \mathbf{0}, & |\mathbf{x}| \lessapprox \mathbf{L}/2\\ V\_0 > \mathbf{0}, & |\mathbf{x}| > \mathbf{L}/2 \end{cases} \tag{1}$$

<sup>ψ</sup>2nþ<sup>1</sup>ðx; <sup>0</sup> <sup>&</sup>lt; <sup>x</sup><sup>⩽</sup> <sup>L</sup>=2Þ ¼ <sup>A</sup>2nþ<sup>1</sup> sin kin,2nþ<sup>1</sup><sup>x</sup> <sup>ψ</sup>2nþ<sup>1</sup>ð Þ¼ <sup>x</sup>; <sup>x</sup> <sup>&</sup>gt; <sup>L</sup>=<sup>2</sup> <sup>B</sup>2nþ1,sexp ð Þ �kout, <sup>2</sup>nþ<sup>1</sup><sup>x</sup>

> kin, <sup>2</sup>nL <sup>2</sup> exp

kin, <sup>2</sup>nL

sin kin, <sup>2</sup>nL kin, <sup>2</sup>nL <sup>þ</sup>

It is convenient to use the potential strength P (introduced by Pitkanen [14], who actually used

<sup>ε</sup><sup>n</sup> <sup>¼</sup> En

<sup>β</sup> <sup>¼</sup> mi m<sup>0</sup>

Φ<sup>n</sup> ¼ kin,n

L

s

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi L2 <sup>2</sup>ℏ<sup>2</sup> miV

<sup>1</sup> � sin kin,2nþ<sup>1</sup><sup>L</sup> kin,2nþ<sup>1</sup><sup>L</sup> <sup>þ</sup>

P ¼

These results generalize the formula (24) in [11] and the Eqs. (25.3e, o) in [13].

<sup>2</sup> exp

The continuity of these functions in x ¼ L=2 gives:

So, the wave function outside the well is:

The wave functions are normalized if:

and to define also εn, β, and X as:

α, instead of P)

B2<sup>n</sup> ¼ A2<sup>n</sup> cos

ψ2nð Þ¼ x > L=2 A2<sup>n</sup> cos

1 A2 2n ¼ L <sup>2</sup> <sup>1</sup> <sup>þ</sup>

> ¼ L 2

1 A2 2nþ1

<sup>ψ</sup>2nþ<sup>1</sup>ð Þ¼ <sup>x</sup> <sup>&</sup>gt; <sup>L</sup>=<sup>2</sup> <sup>A</sup>2nþ<sup>1</sup> sin kin, <sup>2</sup>nþ<sup>1</sup><sup>L</sup>

<sup>B</sup>2nþ<sup>1</sup> <sup>¼</sup> <sup>A</sup>2nþ<sup>1</sup> sin kin,2nþ<sup>1</sup><sup>L</sup>

<sup>ψ</sup>2nþ<sup>1</sup>ð Þ¼� <sup>x</sup> <sup>&</sup>lt; <sup>0</sup> <sup>ψ</sup>2nþ<sup>1</sup>ð Þ �<sup>x</sup> (9)

� � (10)

http://dx.doi.org/10.5772/intechopen.73837

� � (11)

� � � � (12)

2 � � � � (13)

<sup>V</sup> (17)

<sup>2</sup> (19)

2

kout,2nL 2

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

<sup>2</sup> exp �kout,2<sup>n</sup> <sup>x</sup> � <sup>L</sup>

<sup>2</sup> exp �kout, <sup>2</sup>nþ<sup>1</sup> <sup>x</sup> � <sup>L</sup>

1 þ cos kin, <sup>2</sup>nL kout, <sup>2</sup>nL

1 � cos kin,2nþ<sup>1</sup>L kout,2nþ<sup>1</sup>L � � (15)

� � (14)

kout,2nþ<sup>1</sup>L 2

(8)

51

(16)

(18)

considering that its mass is position dependent. More exactly, the mass inside the well, mi, and the mass outside the well, mo, are different:

$$m(\mathbf{x}) = \begin{cases} m\_{i\nu} & |\mathbf{x}| \lessapprox L/2\\ m\_{0\nu} & |\mathbf{x}| > L/2 \end{cases} \tag{2}$$

So, the Schrodinger equation for bound states is:

$$H\psi(\mathbf{x}) = \left[ -\frac{\hbar^2}{2}\frac{d}{d\mathbf{x}}\left(\frac{1}{m(\mathbf{x})}\frac{d}{d\mathbf{x}}\right) + V(\mathbf{x})\right]\psi\_n(\mathbf{x}) = E\_n\psi\_n(\mathbf{x}) - \tag{3}$$

Its physically acceptable solutions, that is, the wave functions, have to satisfy two conditions: (1) the continuity of the wave function and (2) the continuity of the probability currents density at the interface. The first one is encountered in all quantum mechanical problems, but the second one is specific to the case of the position-dependent mass [6], defined by the Eq. (2), and takes the form:

$$\frac{1}{m\_i} \frac{d\psi\_{in}(\mathbf{x} < L/2)}{d\mathbf{x}} \Big|\_{\mathbf{x} \to L/2} = \frac{1}{m\_o} \frac{d\psi\_{out}(\mathbf{x} > L/2)}{d\mathbf{x}} \Big|\_{\mathbf{x} \to L/2} \tag{4}$$

Eq. (4) is known as the BenDaniel-Duke boundary condition. The notations ψin, ψout were used here to make more visible the physical content of this special boundary condition, and will not be maintained in the rest of the chapter.

The nth bound state has a unique energy, En, but two wave vectors, one inside the well, kin,n, and another one outside, kout,n:

$$E\_n = \frac{\hbar^2 k\_{in,n}^2}{2m}, \quad V\_0 - E\_n = \frac{\hbar^2 k\_{out,n}^2}{2m} \tag{5}$$

Due to the parity of the potential, V xð Þ¼ Vð Þ �x , the wave functions can be chosen to be symmetric or antisymmetric.

The symmetric wave functions, describing the even states, are:

$$
\psi\_{2n}(\mathbf{x}, \ 0 < \mathbf{x} \lessapprox L/2) = A\_{2n} \cos k\_{\mathrm{in}, 2n} \mathbf{x}; \ \psi\_{2n}(\mathbf{x}, \ \mathbf{x} > L/2) = B\_{2n} \exp \left( -k\_{\mathrm{out}, 2n} \mathbf{x} \right) \tag{6}
$$

$$
\psi\_{2n}(\mathbf{x} < \mathbf{0}) = \psi\_{2n}(-\mathbf{x})\tag{7}
$$

The ground state wave function is, of course, ψ0ð Þx . The antisymmetric wave functions, describing the odd states, are:

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods http://dx.doi.org/10.5772/intechopen.73837 51

$$\begin{aligned} \psi\_{2n+1}(\mathbf{x}, \ 0 < \mathbf{x} \lessapprox \mathbf{L}/2) &= A\_{2n+1} \sin k\_{in, 2n+1} \mathbf{x} \\ \psi\_{2n+1}(\mathbf{x}, \ \mathbf{x} > \mathbf{L}/2) &= B\_{2n+1, s} \exp \left( -k\_{out, 2n+1} \mathbf{x} \right) \end{aligned} \tag{8}$$

$$
\psi\_{2n+1}(\mathbf{x} < \mathbf{0}) = -\psi\_{2n+1}(-\mathbf{x})\tag{9}
$$

The continuity of these functions in x ¼ L=2 gives:

$$B\_{2n} = A\_{2n} \cos \frac{k\_{in, 2n}L}{2} \exp\left(\frac{k\_{out, 2n}L}{2}\right) \tag{10}$$

$$B\_{2n+1} = A\_{2n+1} \sin \frac{k\_{in, 2n+1}L}{2} \exp\left(\frac{k\_{out, 2n+1}L}{2}\right) \tag{11}$$

So, the wave function outside the well is:

V xð Þ¼ <sup>0</sup>, xj j⩽L=<sup>2</sup>

considering that its mass is position dependent. More exactly, the mass inside the well, mi, and

m xð Þ¼ mi, xj j<sup>⩽</sup> <sup>L</sup>=<sup>2</sup>

1 m xð Þ

d dx 

Its physically acceptable solutions, that is, the wave functions, have to satisfy two conditions: (1) the continuity of the wave function and (2) the continuity of the probability currents density at the interface. The first one is encountered in all quantum mechanical problems, but the second one is specific to the case of the position-dependent mass [6], defined by the Eq. (2),

> ¼ 1 mo

Eq. (4) is known as the BenDaniel-Duke boundary condition. The notations ψin, ψout were used here to make more visible the physical content of this special boundary condition, and

The nth bound state has a unique energy, En, but two wave vectors, one inside the well, kin,n,

Due to the parity of the potential, V xð Þ¼ Vð Þ �x , the wave functions can be chosen to be

The ground state wave function is, of course, ψ0ð Þx . The antisymmetric wave functions,

, V<sup>0</sup> � En <sup>¼</sup> <sup>ℏ</sup><sup>2</sup>

ψ2nðx; 0 < x⩽L=2Þ ¼ A2<sup>n</sup> cos kin, <sup>2</sup>nx; ψ2nð Þ¼ x; x > L=2 B2nexp ð Þ �kout,2nx (6)

dψoutð Þ x > L=2 dx

> k2 out,n

ψ2nð Þ¼ x < 0 ψ2nð Þ �x (7)

 x!L=2

the mass outside the well, mo, are different:

50 Semiconductors - Growth and Characterization

So, the Schrodinger equation for bound states is:

Hψð Þ¼ � x

1 mi

will not be maintained in the rest of the chapter.

and another one outside, kout,n:

symmetric or antisymmetric.

describing the odd states, are:

and takes the form:

ℏ2 2 d dx

dψinð Þ x < L=2 dx

En <sup>¼</sup> <sup>ℏ</sup><sup>2</sup>

The symmetric wave functions, describing the even states, are:

k2 in,n 2m

 x!L=2

V<sup>0</sup> > 0, xj j > L=2

mo, xj j > L=2

þ V xð Þ

(1)

(2)

(4)

ψnð Þ¼ x Enψnð Þ�x (3)

<sup>2</sup><sup>m</sup> (5)

$$\psi\_{2n}(\mathbf{x} > L/2) = A\_{2n} \cos \frac{k\_{in, 2n}L}{2} \exp \left( -k\_{out, 2n} \left( \mathbf{x} - \frac{L}{2} \right) \right) \tag{12}$$

$$\psi\_{2n+1}(\mathbf{x} > L/2) = A\_{2n+1} \sin \frac{k\_{\mathrm{in},2n+1}L}{2} \exp \left( -k\_{\mathrm{out},2n+1} \left( \mathbf{x} - \frac{L}{2} \right) \right) \tag{13}$$

The wave functions are normalized if:

$$\frac{1}{A\_{2n}^2} = \frac{L}{2} \left( 1 + \frac{\sin k\_{in, 2n} L}{k\_{in, 2n} L} + \frac{1 + \cos k\_{in, 2n} L}{k\_{out, 2n} L} \right) \tag{14}$$

$$\frac{1}{A\_{2n+1}^2} = \frac{L}{2} \left( 1 - \frac{\sin k\_{in, 2n+1} L}{k\_{in, 2n+1} L} + \frac{1 - \cos k\_{in, 2n+1} L}{k\_{out, 2n+1} L} \right) \tag{15}$$

These results generalize the formula (24) in [11] and the Eqs. (25.3e, o) in [13].

It is convenient to use the potential strength P (introduced by Pitkanen [14], who actually used α, instead of P)

$$P = \sqrt{\frac{L^2}{2\hbar^2} m\_i V} \tag{16}$$

and to define also εn, β, and X as:

$$
\varepsilon\_n = \frac{E\_n}{V} \tag{17}
$$

$$
\beta = \frac{m\_i}{m\_0} \tag{18}
$$

$$\Phi\_n = k\_{\text{in},n} \frac{L}{2} \tag{19}$$

P, εn, β, and Φ<sup>n</sup> are dimensionless quantities; Φ<sup>n</sup> will be sometimes called dimensionless wave vector.

It is easy to see that:

$$k\_{in,n} \frac{L}{2} = P\sqrt{\epsilon\_n} \tag{20}$$

sin Φ2nþ<sup>1</sup> Φ2nþ<sup>1</sup>

wells ð Þ p ≪ 1 [15] and in the general case [8, 9, 16, 17].

3.1. The first even state (the ground state)

3.1.1. The case β > 1

ground state, is the smallest positive root of the equation:

We shall discuss separately the cases β > 1 and β < 1:

It is useful to introduce the new parameters <sup>γ</sup>>, g>, A<sup>2</sup>

<sup>γ</sup><sup>&</sup>gt; <sup>¼</sup> <sup>β</sup> � <sup>1</sup>, g<sup>&</sup>gt; <sup>¼</sup> <sup>1</sup>

<sup>¼</sup> <sup>1</sup> ffiffiffiffiffiffi γ> p ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi A2 <sup>&</sup>gt; � <sup>Φ</sup><sup>2</sup> 0 <sup>q</sup> , <sup>0</sup> <sup>&</sup>lt; <sup>Φ</sup><sup>0</sup> <sup>&</sup>lt;

because the eigenvalue equation can be written in a simpler form:

cos Φ<sup>0</sup> Φ<sup>0</sup>

cos Φ<sup>0</sup> Φ<sup>0</sup>

¼ �ð Þ<sup>1</sup> <sup>n</sup> <sup>p</sup> ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

3. Approximate analytical solutions for eigenvalue equations

<sup>β</sup> <sup>þ</sup> <sup>1</sup> � <sup>β</sup> � �p<sup>2</sup>Φ<sup>2</sup>

2nþ1

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

For β ¼ 1, they take the form of the well-known equations for the energy eigenvalues of the finite square well. Approximate analytical solutions of these equations were obtained for deep

If 0 < β < 1 β > 1 � �, the rhs of Eqs. (29) and (30) is a monotonically decreasing (increasing) function of Φ; in both cases, the roots of these equations can be obtained using the same approach. In this chapter, we shall obtain precise analytical approximations for the energy of the first two states, that is, for the ground state and for the first excited state, considering the cases β < 1 and β > 1 separately. For moderate and deep wells, the formulae are both simple and accurate. In the limit β ! 1, we shall obtain the result of de Alcantara Bonfim and Griffiths, Eq. (17) of [8].

According to Eq. (29), the dimensionless momentum of the first even state, which is also the

0 <sup>q</sup> , <sup>0</sup> <sup>&</sup>lt; <sup>x</sup> <sup>&</sup>lt;

>:

<sup>&</sup>gt; <sup>¼</sup> <sup>P</sup><sup>2</sup> β <sup>β</sup> � <sup>1</sup> <sup>¼</sup> <sup>P</sup><sup>2</sup> π

π

<sup>2</sup> (31)

βg<sup>&</sup>gt; (32)

<sup>2</sup> (33)

<sup>¼</sup> <sup>p</sup> ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi <sup>β</sup> <sup>þ</sup> <sup>1</sup> � <sup>β</sup> � �p<sup>2</sup>Φ<sup>2</sup>

> γ> , A<sup>2</sup>

In the most physically interesting cases, P is quite large (the wells are quite deep), and according to (32), A<sup>&</sup>gt; is even larger, so it is more convenient to use A instead of P as "large parameter".

<sup>q</sup> , n <sup>¼</sup> <sup>0</sup>, <sup>1</sup>, … odd states (30)

http://dx.doi.org/10.5772/intechopen.73837

53

$$k\_{out,n} \frac{L}{2} = P \sqrt{\frac{1 - \varepsilon\_n}{\beta}} \tag{21}$$

$$k\_{in,n}^2 + \beta k\_{out,n}^2 = \frac{1}{\left(pL/2\right)^2} \tag{22}$$

Let us mention that, if the mass is position-independent, that is, if mi ¼ mo, the eigenvalue equations are (see for instance [5], p. 3, Eqs. (15) and (16)):

$$\tan\frac{k\_{in,2n}L}{2} = \frac{k\_{out,2n}}{k\_{in,2n}},\quad \text{even}\quad \text{states}\tag{23}$$

$$\tan\frac{k\_{in,2n+1}L}{2} = -\frac{k\_{in,2n+1}}{k\_{out,2n+1}},\ \text{odd states}\tag{24}$$

If the mass is position dependent, according to (2), the eigenvalue equations obtained from the Schrodinger equations, using BenDaniel-Duke boundary conditions have the form:

$$\tan\frac{k\_{in,2n}L}{2} = \frac{m\_i}{m\_o}\frac{k\_{out,2n}}{k\_{in,2n}} = = \beta \frac{k\_{out,2n}}{k\_{in,2n}}, \quad \text{even} \quad \text{states} \tag{25}$$

$$\tan\frac{k\_{in,2n+1}L}{2} = -\frac{k\_{in,2n+1}}{k\_{out,2n+1}},\ \text{odd states}\tag{26}$$

We shall consider that both mi, mo are positive; this corresponds to type I semiconductors. So, with β > 0, with kinL=2 replaced by Φ2<sup>n</sup> for even states and by Φ2nþ<sup>1</sup> for odd states, we can put the Eqs. (25) and (26) in a more convenient form:

$$\Phi\_{2n}\tan\Phi\_{2n} = \frac{\sqrt{\beta}}{p}\sqrt{1 - p^2\Phi\_{2n'}^2} \quad n = 0, \ 1, \ldots \\ \text{even states} \tag{27}$$

$$\mathbb{1}\Phi\_{2n+1}\cot\Phi\_{2n+1} = -\frac{\sqrt{\beta}}{p}\sqrt{1 - p^2\Phi\_{2n+1}^2} \quad n = 0, \ 1, \ldots \\ odd \tag{28}$$

or, equivalently:

$$\frac{\cos \Phi\_{2n}}{\Phi\_{2n}} = (-1)^n \frac{p}{\sqrt{\beta + (1 - \beta)p^2 \Phi\_{2n}^2}}, \quad n = 0, \ 1, \ \ldots \quad \text{even states} \tag{29}$$

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods http://dx.doi.org/10.5772/intechopen.73837 53

$$\frac{\sin\Phi\_{2n+1}}{\Phi\_{2n+1}} = (-1)^n \frac{p}{\sqrt{\beta + (1-\beta)p^2\Phi\_{2n+1}^2}}, \quad n = 0, \ 1, \ \dots \text{ odd states} \tag{30}$$

For β ¼ 1, they take the form of the well-known equations for the energy eigenvalues of the finite square well. Approximate analytical solutions of these equations were obtained for deep wells ð Þ p ≪ 1 [15] and in the general case [8, 9, 16, 17].

If 0 < β < 1 β > 1 � �, the rhs of Eqs. (29) and (30) is a monotonically decreasing (increasing) function of Φ; in both cases, the roots of these equations can be obtained using the same approach.

In this chapter, we shall obtain precise analytical approximations for the energy of the first two states, that is, for the ground state and for the first excited state, considering the cases β < 1 and β > 1 separately. For moderate and deep wells, the formulae are both simple and accurate. In the limit β ! 1, we shall obtain the result of de Alcantara Bonfim and Griffiths, Eq. (17) of [8].

#### 3. Approximate analytical solutions for eigenvalue equations

#### 3.1. The first even state (the ground state)

According to Eq. (29), the dimensionless momentum of the first even state, which is also the ground state, is the smallest positive root of the equation:

$$\frac{\cos \Phi\_0}{\Phi\_0} = \frac{p}{\sqrt{\beta + (1 - \beta)p^2 \Phi\_0^2}}, \quad 0 < x < \frac{\pi}{2} \tag{31}$$

We shall discuss separately the cases β > 1 and β < 1:

#### 3.1.1. The case β > 1

P, εn, β, and Φ<sup>n</sup> are dimensionless quantities; Φ<sup>n</sup> will be sometimes called dimensionless

ffiffiffiffiffiffiffiffiffiffiffiffiffi 1 � ε<sup>n</sup> β

s

out,n <sup>¼</sup> <sup>1</sup>

Let us mention that, if the mass is position-independent, that is, if mi ¼ mo, the eigenvalue

<sup>2</sup> <sup>¼</sup> kout, <sup>2</sup><sup>n</sup> kin,2<sup>n</sup>

<sup>2</sup> ¼ � kin,2nþ<sup>1</sup>

Schrodinger equations, using BenDaniel-Duke boundary conditions have the form:

kout,2<sup>n</sup> kin,2<sup>n</sup>

<sup>2</sup> ¼ � kin,2nþ<sup>1</sup>

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi <sup>1</sup> � <sup>p</sup><sup>2</sup>Φ<sup>2</sup> 2n

> ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi <sup>1</sup> � <sup>p</sup><sup>2</sup>Φ<sup>2</sup>

> > 2n

2nþ1

<sup>q</sup> , n <sup>¼</sup> <sup>0</sup>, <sup>1</sup>, … even states (29)

We shall consider that both mi, mo are positive; this corresponds to type I semiconductors. So, with β > 0, with kinL=2 replaced by Φ2<sup>n</sup> for even states and by Φ2nþ<sup>1</sup> for odd states, we can put

If the mass is position dependent, according to (2), the eigenvalue equations obtained from the

¼¼ β

kout, <sup>2</sup>nþ<sup>1</sup>

kout,2<sup>n</sup> kin,2<sup>n</sup>

kout, <sup>2</sup>nþ<sup>1</sup>

p (20)

ð Þ pL=<sup>2</sup> <sup>2</sup> (22)

, even states (23)

, odd states (24)

, even states (25)

, odd states (26)

, n ¼ 0, 1, …even states (27)

, n ¼ 0, 1, …odd states (28)

(21)

kin,n L <sup>2</sup> <sup>¼</sup> <sup>P</sup> ffiffiffiffiffi εn

kout,n L <sup>2</sup> <sup>¼</sup> <sup>P</sup>

k 2 in,n þ βk 2

kin,2nL

kin,2nþ<sup>1</sup>L

kin,2nþ<sup>1</sup>L

ffiffiffi β p p

¼ �ð Þ<sup>1</sup> <sup>n</sup> <sup>p</sup> ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

q

ffiffiffi β p p

q

<sup>β</sup> <sup>þ</sup> <sup>1</sup> � <sup>β</sup> � �p<sup>2</sup>Φ<sup>2</sup>

equations are (see for instance [5], p. 3, Eqs. (15) and (16)):

tan

tan

kin,2nL <sup>2</sup> <sup>¼</sup> mi mo

tan

tan

the Eqs. (25) and (26) in a more convenient form:

Φ2<sup>n</sup> tan Φ2<sup>n</sup> ¼

Φ2nþ<sup>1</sup> cot Φ2nþ<sup>1</sup> ¼ �

cos Φ2<sup>n</sup> Φ2<sup>n</sup>

or, equivalently:

wave vector.

It is easy to see that:

52 Semiconductors - Growth and Characterization

It is useful to introduce the new parameters <sup>γ</sup>>, g>, A<sup>2</sup> >:

$$
\gamma\_{>} = \beta - 1, \text{ g}\_{>} = \frac{1}{\gamma\_{>}}, \quad A\_{>}^{2} = \frac{P^{2}\beta}{\beta - 1} = P^{2}\beta \text{g}\_{>} \tag{32}
$$

because the eigenvalue equation can be written in a simpler form:

$$\frac{\cos \Phi\_0}{\Phi\_0} = \frac{1}{\sqrt{\mathcal{V}\_>} \sqrt{A\_>^2 - \Phi\_0^2}}, \quad 0 < \Phi\_0 < \frac{\pi}{2} \tag{33}$$

In the most physically interesting cases, P is quite large (the wells are quite deep), and according to (32), A<sup>&</sup>gt; is even larger, so it is more convenient to use A instead of P as "large parameter".

We shall replace the exact, transcendental Eq. (31) with an approximate, algebraic equation, using one of the formulae proposed in [8] for cos x, namely:

$$\cos x \simeq f(\mathbf{x}, \mathbf{c}) = \frac{1 - \left(\frac{2\mathbf{x}}{\pi}\right)^2}{\sqrt{1 + c\mathbf{x}^2}} \tag{34}$$

It is useful to write (39) in terms of more physical parameters, p and β: In order to do this, let us

<sup>β</sup> , <sup>α</sup><sup>2</sup>

<sup>β</sup><sup>1</sup>=<sup>2</sup> <sup>þ</sup>

π6

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 þ C 2 s� � <sup>p</sup>

<sup>β</sup><sup>3</sup>=<sup>2</sup> <sup>þ</sup>

<sup>32</sup> ð Þ <sup>1</sup> <sup>þ</sup> <sup>C</sup> <sup>p</sup><sup>2</sup> <sup>þ</sup>

It is a simple exercise to check that the first three terms of the previous formula coincides with

If the parameter A<sup>&</sup>gt; cannot be considered "large," the exact expression of the root can be obtained using the standard approach [19]; they are elementary, but cumbersome, and will be

C 2 � � p<sup>3</sup>

π4

<sup>&</sup>gt; <sup>¼</sup> <sup>β</sup> � <sup>1</sup>

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

π4

<sup>128</sup> <sup>β</sup> � <sup>1</sup> <sup>þ</sup>

π5 32

<sup>32</sup> ð Þ <sup>1</sup> <sup>þ</sup> <sup>C</sup> <sup>p</sup><sup>2</sup>

� �

ffiffiffiffiffiffiffiffiffiffiffiffi 1 þ C 2 r C

β

<sup>2</sup> <sup>p</sup><sup>3</sup> <sup>þ</sup>

π

<sup>0</sup> ¼ z (45)

ð Þ <sup>1</sup> <sup>þ</sup> <sup>C</sup> <sup>p</sup><sup>4</sup>

π6 128 C

C 2

<sup>β</sup> <sup>p</sup><sup>2</sup> (40)

http://dx.doi.org/10.5772/intechopen.73837

<sup>β</sup><sup>2</sup> <sup>þ</sup> …

<sup>2</sup> ð Þ <sup>1</sup> <sup>þ</sup> <sup>C</sup> <sup>p</sup><sup>4</sup> <sup>þ</sup> … (42)

<sup>2</sup> (43)

βγ<sup>&</sup>lt; (44)

(41)

55

<sup>g</sup>>α<sup>2</sup> <sup>&</sup>gt; <sup>¼</sup> <sup>p</sup><sup>2</sup>

<sup>4</sup> � <sup>π</sup><sup>3</sup> 8

β � 1 þ

notice that:

and:

<sup>z</sup> <sup>β</sup> <sup>¼</sup> <sup>1</sup> � � <sup>¼</sup> <sup>π</sup><sup>2</sup>

3.1.2. The case β < 1

a cubic equation:

so Eq. (39) takes the form:

þ π5 32

<sup>4</sup> � <sup>π</sup><sup>3</sup> 8

<sup>z</sup> <sup>β</sup> <sup>&</sup>gt; <sup>1</sup> � � <sup>¼</sup> <sup>π</sup><sup>2</sup>

r

ffiffiffiffiffiffiffiffiffiffiffiffi 1 þ C 2

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 þ C 2 s� �<sup>p</sup> <sup>þ</sup>

the first three terms of the power series given by Eq. (17) of [8].

not given here; the interested reader can find them in [18] .

cos Φ<sup>0</sup> Φ<sup>0</sup>

<sup>γ</sup><sup>&</sup>lt; <sup>¼</sup> <sup>1</sup> � <sup>β</sup>, <sup>1</sup>

with the following definitions for the parameters:

equation, which becomes, with the same substitution

If β < 1, the eigenvalue equation for the dimensionless wave vector is:

<sup>¼</sup> <sup>1</sup> ffiffiffiffiffiffi γ< p ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi A2 <sup>&</sup>lt; <sup>þ</sup> <sup>Φ</sup><sup>2</sup> 0 <sup>q</sup> , <sup>0</sup> <sup>&</sup>lt; <sup>Φ</sup><sup>0</sup> <sup>&</sup>lt;

γ<

<sup>¼</sup> <sup>g</sup><, A<sup>2</sup>

Using the de Alcantara Bonfim-Griffiths algebraization for cos x (34) [8], it gives an algebraic

Φ2

<sup>&</sup>lt; <sup>¼</sup> <sup>P</sup><sup>2</sup> β <sup>1</sup> � <sup>β</sup> <sup>¼</sup> <sup>P</sup><sup>2</sup>

The precision of this approximation on various subintervals of 0ð Þ ; 1 depends on the exact value of c, with 0:18 ≲c≲0:23; in our numerical evaluation, we shall use the value c ¼ 0:22: For a detailed discussion on this issue, see [18].

The algebraic approximation of the eigenvalue equation, we get with (34) is:

$$\frac{1}{\Phi\_0} \frac{1 - \left(\frac{2\Phi\_0}{\pi}\right)^2}{\sqrt{1 + c\Phi\_0^2}} = \frac{1}{\sqrt{\mathcal{V}\_>} \sqrt{A\_>^2 - \Phi\_0^2}}\tag{35}$$

with

$$
\Phi\_0^2 = z \tag{36}
$$

(35) can be written as:

$$z^3 + \left(\frac{1}{16}\pi^4 c \mathfrak{g}\_{>} - A\_{>}^2 - \frac{1}{2}\pi^2\right)z^2 + \frac{\pi^2}{2}\left(\frac{\pi^2}{8}\mathfrak{g}\_{>} + A\_{>}^2 + \frac{1}{8}\pi^2\right)z - \frac{1}{16}\pi^4 A\_{>}^2 = 0\tag{37}$$

Following the approach outlined in [19] and applied to this problem in [18], introducing the notation:

$$\mathbf{C} = \frac{\pi^2}{2}\mathbf{c} \tag{38}$$

and considering that the well is not too shallow:

A2 <sup>&</sup>gt; ≫ 1

we obtain for the physically interesting root the expression:

$$\begin{split} z(\beta > 1) &= \frac{\pi^2}{4} - \frac{\pi^3}{8} \sqrt{\left(1 + \frac{\mathsf{C}}{2}\right)} \left(\sqrt{\mathsf{g}\_>} a\_>\right) + \frac{\pi^4}{32} (1 + \mathsf{C}) \left(\sqrt{\mathsf{g}\_>} a\_>\right)^2 \\ &+ \frac{\pi^5}{32} \left(1 + \frac{\mathsf{C} \mathsf{g}\_>}{2}\right) \sqrt{\left(1 + \frac{\mathsf{C}}{2}\right)} \sqrt{\mathsf{g}\_>} a\_>^3 + \frac{\pi^6}{128} \mathsf{g}\_> (1 + \mathsf{C}) \left(1 + \frac{\mathsf{C} \mathsf{g}\_>}{2}\right) a\_>^4 + \dots \end{split} \tag{39}$$

If the depth of the well increases indefinitely, <sup>α</sup><sup>&</sup>gt; ! 0 and <sup>z</sup><sup>1</sup> ! <sup>π</sup><sup>2</sup>=4, <sup>Φ</sup><sup>0</sup> ! <sup>π</sup>=2, as requested. Indeed, in a finite well, the energy of a bound state is smaller than the corresponding energy in an infinite one, so the first term in ffiffiffiffiffi g> p α<sup>&</sup>gt; in the previous formula is negative.

It is useful to write (39) in terms of more physical parameters, p and β: In order to do this, let us notice that:

$$g\_{>}a\_{>}^{2} = \frac{p^{2}}{\beta}, \quad a\_{>}^{2} = \frac{\beta - 1}{\beta}p^{2} \tag{40}$$

so Eq. (39) takes the form:

$$\begin{split} z(\beta > 1) &= \frac{\pi^2}{4} - \frac{\pi^3}{8} \sqrt{\left(1 + \frac{\mathsf{C}}{2}\right)} \frac{p}{\beta^{1/2}} + \frac{\pi^4}{32} (1 + \mathsf{C}) \frac{p^2}{\beta} \\ &+ \frac{\pi^5}{32} \sqrt{1 + \frac{\mathsf{C}}{2}} \left(\beta - 1 + \frac{\mathsf{C}}{2}\right) \frac{p^3}{\beta^{3/2}} + \frac{\pi^6}{128} \left(\beta - 1 + \frac{\mathsf{C}}{2}\right) (1 + \mathsf{C}) \frac{p^4}{\beta^2} + \dots \end{split} \tag{41}$$

and:

We shall replace the exact, transcendental Eq. (31) with an approximate, algebraic equation,

The precision of this approximation on various subintervals of 0ð Þ ; 1 depends on the exact value of c, with 0:18 ≲c≲0:23; in our numerical evaluation, we shall use the value c ¼ 0:22:

> ffiffiffiffiffiffi γ> p ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi A2 <sup>&</sup>gt; � <sup>Φ</sup><sup>2</sup> 0

π2

Following the approach outlined in [19] and applied to this problem in [18], introducing the

<sup>C</sup> <sup>¼</sup> <sup>π</sup><sup>2</sup>

A2 <sup>&</sup>gt; ≫ 1

g> p α<sup>&</sup>gt; � � <sup>þ</sup>

If the depth of the well increases indefinitely, <sup>α</sup><sup>&</sup>gt; ! 0 and <sup>z</sup><sup>1</sup> ! <sup>π</sup><sup>2</sup>=4, <sup>Φ</sup><sup>0</sup> ! <sup>π</sup>=2, as requested. Indeed, in a finite well, the energy of a bound state is smaller than the corresponding energy in

g> p α<sup>3</sup> <sup>&</sup>gt; þ π6 π4

p α<sup>&</sup>gt; in the previous formula is negative.

<sup>32</sup> ð Þ <sup>1</sup> <sup>þ</sup> <sup>C</sup> ffiffiffiffiffi

<sup>128</sup> <sup>g</sup>>ð Þ <sup>1</sup> <sup>þ</sup> <sup>C</sup> <sup>1</sup> <sup>þ</sup> Cg<sup>&</sup>gt;

g> p α<sup>&</sup>gt; � �<sup>2</sup>

2 � �

α4 <sup>&</sup>gt; þ …

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 þ C 2 s� � ffiffiffiffiffi

g>

<sup>8</sup> <sup>g</sup><sup>&</sup>gt; <sup>þ</sup> <sup>A</sup><sup>2</sup>

<sup>&</sup>gt; þ 1 8 π2

� �

Φ2

<sup>z</sup><sup>2</sup> <sup>þ</sup> π2 2

<sup>1</sup> � <sup>2</sup><sup>x</sup> π � �<sup>2</sup> ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

<sup>1</sup> <sup>þ</sup> cx<sup>2</sup> <sup>p</sup> (34)

<sup>q</sup> (35)

<sup>0</sup> ¼ z (36)

<sup>z</sup> � <sup>1</sup> <sup>16</sup> <sup>π</sup><sup>4</sup> A2

<sup>2</sup> <sup>c</sup> (38)

<sup>&</sup>gt; ¼ 0 (37)

(39)

cos x≃ f xð Þ¼ ; c

The algebraic approximation of the eigenvalue equation, we get with (34) is:

<sup>1</sup> � <sup>2</sup>Φ<sup>0</sup> π � �<sup>2</sup> ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi <sup>1</sup> <sup>þ</sup> <sup>c</sup>Φ<sup>2</sup> 0 <sup>q</sup> <sup>¼</sup> <sup>1</sup>

1 Φ<sup>0</sup>

using one of the formulae proposed in [8] for cos x, namely:

For a detailed discussion on this issue, see [18].

54 Semiconductors - Growth and Characterization

with

notation:

(35) can be written as:

<sup>z</sup><sup>3</sup> <sup>þ</sup>

1 <sup>16</sup> <sup>π</sup><sup>4</sup>

cg<sup>&</sup>gt; � <sup>A</sup><sup>2</sup>

and considering that the well is not too shallow:

<sup>z</sup> <sup>β</sup> <sup>&</sup>gt; <sup>1</sup> � � <sup>¼</sup> <sup>π</sup><sup>2</sup>

<sup>32</sup> <sup>1</sup> <sup>þ</sup> Cg<sup>&</sup>gt; 2

an infinite one, so the first term in ffiffiffiffiffi

þ π5

we obtain for the physically interesting root the expression:

<sup>4</sup> � <sup>π</sup><sup>3</sup> 8

� � ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

1 þ C 2 s� � ffiffiffiffiffi

� �

<sup>&</sup>gt; � <sup>1</sup> 2 π2

$$z(\beta=1) = \frac{\pi^2}{4} - \frac{\pi^3}{8} \sqrt{\left(1 + \frac{\mathbb{C}}{2}\right)} p + \frac{\pi^4}{32} (1 + \mathbb{C}) p^2 + \frac{\pi^5}{32} \sqrt{1 + \frac{\mathbb{C}}{2}\frac{\mathbb{C}}{2}} p^3 + \frac{\pi^6}{128} \frac{\mathbb{C}}{2} (1 + \mathbb{C}) p^4 + \dots \tag{42}$$

It is a simple exercise to check that the first three terms of the previous formula coincides with the first three terms of the power series given by Eq. (17) of [8].

If the parameter A<sup>&</sup>gt; cannot be considered "large," the exact expression of the root can be obtained using the standard approach [19]; they are elementary, but cumbersome, and will be not given here; the interested reader can find them in [18] .

#### 3.1.2. The case β < 1

If β < 1, the eigenvalue equation for the dimensionless wave vector is:

$$\frac{\cos \Phi\_0}{\Phi\_0} = \frac{1}{\sqrt{\mathcal{V}\_{<}} \sqrt{A\_{<}^2 + \Phi\_0^2}}, \quad 0 < \Phi\_0 < \frac{\pi}{2} \tag{43}$$

with the following definitions for the parameters:

$$\mathcal{V}\_{<} = 1 - \beta, \ \frac{1}{\mathcal{V}\_{<}} = \mathcal{g}\_{<}, \ A\_{<}^{2} = \frac{P^{2}\beta}{1 - \beta} = P^{2}\beta\mathcal{V}\_{<} \tag{44}$$

Using the de Alcantara Bonfim-Griffiths algebraization for cos x (34) [8], it gives an algebraic equation, which becomes, with the same substitution

$$
\Phi\_0^2 = z \tag{45}
$$

a cubic equation:

$$z^3 + \left(A\_<^2 - \frac{1}{2}\pi^2 - \frac{1}{16}\pi^4 \text{cg}\_<\right)z^2 + \frac{\pi^2}{2}\left(\frac{1}{8}\pi^2 - A\_<^2 - \frac{\pi^2}{8}\text{g}\_<\right)z + \frac{1}{16}\pi^4 A\_<^2 = 0, \qquad \beta < 1 \tag{46}$$

Following the same steps as in the previous case, we find that the parameters g>, α<sup>&</sup>gt; enter into the various expressions needed for obtaining the cubic roots only through the monoms g>α<sup>2</sup> >, α2 <sup>&</sup>gt; at various powers, and the roots of Eq. (46) can be obtained from the root (39) making the substitution:

$$\mathfrak{g}\_{>} \to -\mathfrak{g}\_{<'} \quad a\_{>}^{2} \to -a\_{<}^{2} \tag{47}$$

of the parameters p, β, a, c: In other words, the main contribution to the errors of our results is given by the approximation of trigonometric functions with algebraic ones, not by the approximation of the exact formulae of the roots of cubic equations with the low order terms of their

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

As already mentioned, one of the physical motivations of the calculation of the energy of bound states in heterostructures is to explain their photoluminescence properties. In several cases (see for instance [20]), the authors use Barker's formula for the energy levels in a square well [15]. Much more precise analytical expressions for these energy are available in the literature [8, 9], for the case of constant mass; in this paper, we propose similar formulas,

In the previous subsections, we analyzed the ground state ð Þ n ¼ 0 and the first excited state ð Þ n ¼ 1 of a square well, with BenDaniel-Duke boundary conditions. For n⩾2, the de Alcantara

> 2 =π<sup>2</sup>

2 <sup>q</sup> , <sup>2</sup>n<sup>π</sup> <sup>&</sup>lt; <sup>Φ</sup> <sup>&</sup>lt; <sup>2</sup>n<sup>π</sup> <sup>þ</sup>

but the eigenvalue equation, obtained in this way, is a sextic equation (which cannot be reduced

the eigenvalue equation the algebraization of tan (see later on, Eq. (73) and (74) of the present

gives a quartic equation for the dimensionless wave vector. Its roots are given by complicated,

In order to illustrate graphically some of our results, let us notice that, using Eqs. (17)–(22), we

where zn is the root of the cubic equations obtained after the algebraization of the transcendental eigenvalue equations for the ground state ð Þ n ¼ 0 and for the first excited state ð Þ n ¼ 1 : According to the Eqs. (39) and (50), for a deep well, the root z can be approximated with a quartic polynomial in p, the inverse of the potential strength P: Let us mention that, if we replace in the definition of P, Eq. (16), mi with the free electron mass, we choose the length of

miL<sup>2</sup>

the well L ¼ 10 nm and we express the potential V<sup>0</sup> in electron volts, we get:

V0

<sup>p</sup> , p <sup>¼</sup> <sup>3</sup>:<sup>9</sup> � <sup>10</sup>�<sup>2</sup> <sup>1</sup>

ffiffiffiffiffiffi V0

<sup>P</sup> <sup>¼</sup> <sup>25</sup>:<sup>616</sup> ffiffiffiffiffiffi

<sup>2</sup>ℏ<sup>2</sup> En <sup>¼</sup> <sup>Φ</sup><sup>2</sup>

), so it cannot be solved. We meet similar difficulties if we try to use in

<sup>2</sup> or sin ð Þ <sup>x</sup>=<sup>x</sup>

2

<sup>n</sup> ¼ znð Þp (52)

p (53)

π

http://dx.doi.org/10.5772/intechopen.73837

57

<sup>2</sup> (51)

, in the sense used in [16],

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 þ cð Þ Φ � 2nπ

series expansions.

3.3. Higher-order states

to a cubic equation in Φ<sup>2</sup>

but still elementary formulas.

considering the case of position-dependent mass.

Bonfim formula (34) can be extended to larger arguments:

paper). Even the "parabolic approximation" for cos ð Þ x=x

3.4. Graphical illustration of our main results

can write the following relations for the energy:

cos <sup>Φ</sup> <sup>≃</sup> <sup>1</sup> � <sup>4</sup>ð Þ <sup>Φ</sup> � <sup>2</sup>n<sup>π</sup>

in Eq. (39). The final result, z β < 1 � �, expressed in terms of p and β, has exactly the form (41).

#### 3.2. The first odd state

#### 3.2.1. The case β > 1

The exact eigenvalue equation for the first odd state, which is also the first excited state, can be written as:

$$\frac{\sin\Phi\_1}{\Phi\_1} = \frac{1}{\sqrt{\mathcal{V}\_>}\sqrt{A\_>^2 - \Phi\_1^2}}, \quad \frac{\pi}{2} < \Phi\_1 < \pi \tag{48}$$

As the shape of the function sin x=x on the interval 0½ � ; π is quite similar with the shape of cos x on the interval 0½ � ; π=2 , we can try an algebraization for sin x=x similar to that proposed by de Alcantara Bonfim and Griffiths for cos x :

$$\frac{\sin \Phi\_1}{\Phi\_1} \simeq \frac{1 - \left(\Phi\_1/\pi\right)^2}{\sqrt{1 + a\Phi\_1^2}} \quad , \quad 0 < \Phi\_1 < \pi, \quad a \simeq 0.2 \tag{49}$$

A detailed discussion of the precision of this approximation is given in [18] (see Fig. 3 and Eq. (88)). Following, exactly the same steps as in the case of the ground state, we find that

$$z(\beta) = \pi^2 - \pi^2 \sqrt{(1+\pi^2 a)} \frac{p}{\sqrt{\beta}} + \frac{\pi^4}{2} a \frac{p^2}{\beta} + \pi^4 \sqrt{(1+\pi^2 a)} (\beta - 1) \left(\frac{p}{\sqrt{\beta}}\right)^3 + \frac{\pi^6}{2} a (\beta - 1) \frac{p^4}{\beta^2} + \dots \tag{50}$$

For β < 1, the expression of the root, in terms of β and p is identical with (50), written in terms of g<, g>, α<, α>, the formulae are different, see Eqs. (99) and (103) in [18].

For both cases—β≶1—in the limit of an infinitely deep root, <sup>z</sup> <sup>β</sup>≶1; <sup>α</sup><sup>&</sup>lt; <sup>¼</sup> <sup>0</sup> � � <sup>¼</sup> <sup>π</sup><sup>2</sup>, <sup>Φ</sup>1ðβ≶1, α<sup>&</sup>lt; ¼ 0Þ ¼ π, as requested, and the first correction to this value is negative.

The relative errors of the formulas (39) and (50), with respect to the exact roots of the corresponding algebraic equations, are very small—of about 10�<sup>4</sup> …10�<sup>6</sup> for physically interesting values of the parameters p, β, a, c: In other words, the main contribution to the errors of our results is given by the approximation of trigonometric functions with algebraic ones, not by the approximation of the exact formulae of the roots of cubic equations with the low order terms of their series expansions.

As already mentioned, one of the physical motivations of the calculation of the energy of bound states in heterostructures is to explain their photoluminescence properties. In several cases (see for instance [20]), the authors use Barker's formula for the energy levels in a square well [15]. Much more precise analytical expressions for these energy are available in the literature [8, 9], for the case of constant mass; in this paper, we propose similar formulas, considering the case of position-dependent mass.

#### 3.3. Higher-order states

<sup>z</sup><sup>3</sup> <sup>þ</sup> <sup>A</sup><sup>2</sup>

substitution:

3.2. The first odd state

3.2.1. The case β > 1

written as:

z β

� � <sup>¼</sup> <sup>π</sup><sup>2</sup> � <sup>π</sup><sup>2</sup>

α2

<sup>&</sup>lt; � <sup>1</sup> 2 <sup>π</sup><sup>2</sup> � <sup>1</sup> <sup>16</sup> <sup>π</sup><sup>4</sup> cg<sup>&</sup>lt; <sup>z</sup><sup>2</sup> <sup>þ</sup> π2 2

sin Φ<sup>1</sup> Φ<sup>1</sup>

sin Φ<sup>1</sup> Φ<sup>1</sup>

> ffiffiffi <sup>β</sup> <sup>p</sup> <sup>þ</sup>

nding algebraic equations, are very small—of about 10�<sup>4</sup>

π4 2 a p2 <sup>β</sup> <sup>þ</sup> <sup>π</sup><sup>4</sup>

of g<, g>, α<, α>, the formulae are different, see Eqs. (99) and (103) in [18].

α<sup>&</sup>lt; ¼ 0Þ ¼ π, as requested, and the first correction to this value is negative.

Alcantara Bonfim and Griffiths for cos x :

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 þ π<sup>2</sup> ð Þa q p <sup>¼</sup> <sup>1</sup> ffiffiffiffiffiffi γ> p ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi A2 <sup>&</sup>gt; � <sup>Φ</sup><sup>2</sup> 1 <sup>q</sup> , <sup>π</sup>

<sup>≃</sup> <sup>1</sup> � ð Þ <sup>Φ</sup>1=<sup>π</sup>

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi <sup>1</sup> <sup>þ</sup> <sup>a</sup>Φ<sup>2</sup> 1

1 8

<sup>g</sup><sup>&</sup>gt; ! �g<, <sup>α</sup><sup>2</sup>

<sup>π</sup><sup>2</sup> � <sup>A</sup><sup>2</sup>

Following the same steps as in the previous case, we find that the parameters g>, α<sup>&</sup>gt; enter into the various expressions needed for obtaining the cubic roots only through the monoms g>α<sup>2</sup>

<sup>&</sup>gt; at various powers, and the roots of Eq. (46) can be obtained from the root (39) making the

in Eq. (39). The final result, z β < 1 � �, expressed in terms of p and β, has exactly the form (41).

The exact eigenvalue equation for the first odd state, which is also the first excited state, can be

As the shape of the function sin x=x on the interval 0½ � ; π is quite similar with the shape of cos x on the interval 0½ � ; π=2 , we can try an algebraization for sin x=x similar to that proposed by de

2

A detailed discussion of the precision of this approximation is given in [18] (see Fig. 3 and Eq. (88)). Following, exactly the same steps as in the case of the ground state, we find that

q

For β < 1, the expression of the root, in terms of β and p is identical with (50), written in terms

For both cases—β≶1—in the limit of an infinitely deep root, <sup>z</sup> <sup>β</sup>≶1; <sup>α</sup><sup>&</sup>lt; <sup>¼</sup> <sup>0</sup> � � <sup>¼</sup> <sup>π</sup><sup>2</sup>, <sup>Φ</sup>1ðβ≶1,

The relative errors of the formulas (39) and (50), with respect to the exact roots of the correspo-

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 1 þ π<sup>2</sup> ð Þa

<sup>&</sup>lt; � <sup>π</sup><sup>2</sup> <sup>8</sup> <sup>g</sup><sup>&</sup>lt;

<sup>&</sup>gt; ! �α<sup>2</sup>

z þ 1 <sup>16</sup> <sup>π</sup><sup>4</sup> A2

<sup>&</sup>lt; ¼ 0, β < 1 (46)

<sup>&</sup>lt; (47)

<sup>2</sup> <sup>&</sup>lt; <sup>Φ</sup><sup>1</sup> <sup>&</sup>lt; <sup>π</sup> (48)

<sup>q</sup> , <sup>0</sup> <sup>&</sup>lt; <sup>Φ</sup><sup>1</sup> <sup>&</sup>lt; <sup>π</sup>, a <sup>≃</sup>0:<sup>2</sup> (49)

<sup>β</sup> � <sup>1</sup> � � <sup>p</sup>

ffiffiffi β p !<sup>3</sup>

þ π6

…10�<sup>6</sup> for physically interesting values

<sup>2</sup> <sup>a</sup> <sup>β</sup> � <sup>1</sup> � � <sup>p</sup><sup>4</sup>

<sup>β</sup><sup>2</sup> <sup>þ</sup> …

(50)

>,

� �

� �

56 Semiconductors - Growth and Characterization

In the previous subsections, we analyzed the ground state ð Þ n ¼ 0 and the first excited state ð Þ n ¼ 1 of a square well, with BenDaniel-Duke boundary conditions. For n⩾2, the de Alcantara Bonfim formula (34) can be extended to larger arguments:

$$\cos\Phi \simeq \frac{1 - 4(\Phi - 2n\pi)^2/\pi^2}{\sqrt{1 + c(\Phi - 2n\pi)}^2}, \quad 2n\pi < \Phi < 2n\pi + \frac{\pi}{2} \tag{51}$$

but the eigenvalue equation, obtained in this way, is a sextic equation (which cannot be reduced to a cubic equation in Φ<sup>2</sup> ), so it cannot be solved. We meet similar difficulties if we try to use in the eigenvalue equation the algebraization of tan (see later on, Eq. (73) and (74) of the present paper). Even the "parabolic approximation" for cos ð Þ x=x <sup>2</sup> or sin ð Þ <sup>x</sup>=<sup>x</sup> 2 , in the sense used in [16], gives a quartic equation for the dimensionless wave vector. Its roots are given by complicated, but still elementary formulas.

#### 3.4. Graphical illustration of our main results

In order to illustrate graphically some of our results, let us notice that, using Eqs. (17)–(22), we can write the following relations for the energy:

$$\frac{m\_i L^2}{2\hbar^2} E\_n = \Phi\_n^2 = z\_n(p) \tag{52}$$

where zn is the root of the cubic equations obtained after the algebraization of the transcendental eigenvalue equations for the ground state ð Þ n ¼ 0 and for the first excited state ð Þ n ¼ 1 : According to the Eqs. (39) and (50), for a deep well, the root z can be approximated with a quartic polynomial in p, the inverse of the potential strength P: Let us mention that, if we replace in the definition of P, Eq. (16), mi with the free electron mass, we choose the length of the well L ¼ 10 nm and we express the potential V<sup>0</sup> in electron volts, we get:

$$P = 25.616\sqrt{V\_0} \cdot p = 3.9 \times 10^{-2} \frac{1}{\sqrt{V\_0}} \tag{53}$$

tan

tan

kin,2nL

and can be solved following exactly the same approach.

<sup>2</sup> ¼ � mi

kin,2nþ<sup>1</sup>L

j j mo

<sup>2</sup> <sup>¼</sup> j j mo mi

kout,2<sup>n</sup> kin,2<sup>n</sup>

¼ � β kout,2<sup>n</sup> kin,2<sup>n</sup>

> ¼ 1 β β

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

kin,2nþ<sup>1</sup> kout,2nþ<sup>1</sup>

As already mentioned, the wave function in the Schrodinger Eq. (3) can be interpreted as an envelope function. This approximation works well when the materials constituting the heterostructures are perfectly lattice-matched and they crystallize in the same crystallographic structure (in the most cases, the zinc blend structure). Its application is restricted to the vicinity of the high-symmetry points in the host's Brillouin zone ð Þ Γ; X; L : Actually, most of the heterostructures' energy levels relevant to actual devices are relatively closed to a symmetry

Figure 2. Schematic representation of the conduction band Ec and of the valence band Ev for type I (a) and type (II) (b)

semiconductors.

, even states (54)

http://dx.doi.org/10.5772/intechopen.73837

59

, odd states (55)

Figure 1. The plot of Φ<sup>2</sup> <sup>0</sup> ð Þ<sup>a</sup> and <sup>Φ</sup><sup>2</sup> <sup>1</sup> ð Þb , which are proportional to the energies E0, E1, as functions of the inverse potential strength p:

We shall plot our main results, that is, the series expansions of the dimensionless wave vectors, Φ2 <sup>0</sup> and Φ<sup>2</sup> <sup>1</sup>, as functions of p, on the range 0 < p < 0:1, when the condition of convergence is satisfactorily fulfilled (Figure 1). The energy is a monotonically increasing function of β; its values, for β ¼ 1, are obtained from Eqs. (42) and (50).

#### 3.5. Applications to other nanostructures

Our calculations can be easily applied to type II semiconductors heterostructures, when one of the effective mass of the charge carrier is negative: mimo < 0 ([5], chapter 3, Eqs. (35) and (36)); a detailed description of such heterointerfaces can be found for instance in [5], p. 66. So, instead of (25) and (26), the eigenvalue equations take the form:

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods http://dx.doi.org/10.5772/intechopen.73837 59

$$\tan\frac{k\_{in,2n}L}{2} = -\frac{m\_i}{|m\_0|}\frac{k\_{out,2n}}{k\_{in,2n}} = -|\beta|\frac{k\_{out,2n}}{k\_{in,2n}}, \quad \text{even} \quad \text{states} \tag{54}$$

$$\tan\frac{k\_{in,2n+1}L}{2} = \frac{|m\_o|}{m\_i}\frac{k\_{in,2n+1}}{k\_{out,2n+1}} = \frac{1}{|\beta|} \left|\beta\right|\_\prime \text{ odd states} \tag{55}$$

and can be solved following exactly the same approach.

We shall plot our main results, that is, the series expansions of the dimensionless wave vectors,

satisfactorily fulfilled (Figure 1). The energy is a monotonically increasing function of β; its

Our calculations can be easily applied to type II semiconductors heterostructures, when one of the effective mass of the charge carrier is negative: mimo < 0 ([5], chapter 3, Eqs. (35) and (36)); a detailed description of such heterointerfaces can be found for instance in [5], p. 66. So, instead

values, for β ¼ 1, are obtained from Eqs. (42) and (50).

<sup>0</sup> ð Þ<sup>a</sup> and <sup>Φ</sup><sup>2</sup>

of (25) and (26), the eigenvalue equations take the form:

3.5. Applications to other nanostructures

<sup>1</sup>, as functions of p, on the range 0 < p < 0:1, when the condition of convergence is

<sup>1</sup> ð Þb , which are proportional to the energies E0, E1, as functions of the inverse

Φ2 <sup>0</sup> and Φ<sup>2</sup>

Figure 1. The plot of Φ<sup>2</sup>

58 Semiconductors - Growth and Characterization

potential strength p:

As already mentioned, the wave function in the Schrodinger Eq. (3) can be interpreted as an envelope function. This approximation works well when the materials constituting the heterostructures are perfectly lattice-matched and they crystallize in the same crystallographic structure (in the most cases, the zinc blend structure). Its application is restricted to the vicinity of the high-symmetry points in the host's Brillouin zone ð Þ Γ; X; L : Actually, most of the heterostructures' energy levels relevant to actual devices are relatively closed to a symmetry

Figure 2. Schematic representation of the conduction band Ec and of the valence band Ev for type I (a) and type (II) (b) semiconductors.

point in the host's Brillouin zone. A popular example is given by the lowest conduction states of GaAs-GaAlAs heterostructures with GaAs layer (typically, its thickness is about 100 Å or larger). A detailed description of the cases in which the envelope function model is successful is given in [5], p. 66 (See Figure 2).

As there are some similarities between QWs and QDs, our results are also relevant for these devices. The simplest remark is that the eigenvalues equations for the first odd state in a QW are identical to that corresponding to the l ¼ 0 state in a QD (see for instance [13], problem 63). Also, the eigenvalue equations for the wave vectors of the energy levels for a finite barrier rectangular shaped QD, Eq. (36) in [21], are quite similar to ours—(29) and (30), but somewhat more complicated. The ground state energy of electrons and holes in a core/shell QD is given by the Eq. (21) of [22], an equation similar to ours, just mentioned previously. Such results are important, inter alia, for the interpretation of photoluminescence spectra and photon harvesting of QDs.

#### 4. The infinite square well with two semiconductor slabs

#### 4.1. The symmetric case

Let us consider an infinite 1D square well, delimited by two rigid walls situated in �L=2, respectively L=2, containing two semiconductor slabs, of equal width, but of different materials. It is a toy model for a Janus nanorod, composed of two different semiconductors, with large work functions. We preferred to choose this particular case (equal width), in order to avoid too cumbersome mathematical calculations. The electron effective mass is position dependent, like in (2):

$$m(\mathbf{x}) = \begin{cases} m\_1 & -L/2 < \mathbf{x} < \mathbf{0} \\ m\_2 & \mathbf{0} < \mathbf{x} < L/2 \end{cases} \tag{56}$$

with:

$$
\beta m\_2 = \beta m\_1 \tag{57}
$$

ψð Þ¼ x

The boundary conditions for the wave function give:

The BenDaniel-Duke boundary condition means:

primary interest, we can combine (64) and (65) to obtain:

and the continuity in the origin:

or:

we get:

With

it can be written as:

(

A<sup>1</sup> sin k1x þ φ<sup>1</sup>

sin � kL

sin ffiffiffi <sup>β</sup> <sup>p</sup> kL

1 m<sup>1</sup> ψ0 ð Þ¼ 0�

ffiffiffi

1 ffiffiffi

kL

<sup>2</sup> ¼ � ffiffiffi

kL

<sup>2</sup> , <sup>φ</sup><sup>2</sup> <sup>¼</sup> <sup>n</sup>2<sup>π</sup> � ffiffiffi

β p tan ffiffiffi

<sup>β</sup> <sup>p</sup> kL

Replacing in (66), the values of φ1, φ<sup>2</sup> obtained from (61) and (62):

φ<sup>1</sup> ¼ n1π þ

tan kL

<sup>β</sup> <sup>p</sup> <sup>k</sup>1<sup>x</sup> <sup>þ</sup> <sup>φ</sup><sup>2</sup>

<sup>2</sup> <sup>þ</sup> <sup>φ</sup><sup>1</sup> � �

<sup>2</sup> <sup>þ</sup> <sup>φ</sup><sup>2</sup> � �

> 1 m<sup>2</sup> ψ0

Together with the orthonormality condition for the wave function, Eqs. (61)–(63) and (65) form a system of five equations for five quantities, k, φ1, φ2, A1, A2: As the amplitudes are not of

A<sup>2</sup> sin ffiffiffi

� �, � <sup>L</sup>=<sup>2</sup> <sup>&</sup>lt; <sup>x</sup> <sup>&</sup>lt; <sup>0</sup>

� �, � <sup>L</sup>=<sup>2</sup> <sup>&</sup>lt; <sup>x</sup> <sup>&</sup>lt; <sup>0</sup>

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

(60)

61

¼ 0 (61)

http://dx.doi.org/10.5772/intechopen.73837

¼ 0 (62)

ð Þ 0<sup>þ</sup> (64)

<sup>2</sup> (67)

<sup>2</sup> (68)

<sup>2</sup> <sup>¼</sup> <sup>K</sup> (69)

A<sup>1</sup> sinφ<sup>1</sup> ¼ A<sup>2</sup> sinφ<sup>2</sup> (63)

<sup>β</sup> <sup>p</sup> <sup>A</sup><sup>1</sup> cosφ<sup>1</sup> <sup>¼</sup> <sup>A</sup><sup>2</sup> cos <sup>φ</sup><sup>2</sup> (65)

<sup>β</sup> <sup>p</sup> tan <sup>φ</sup><sup>1</sup> <sup>¼</sup> tanφ<sup>2</sup> (66)

<sup>β</sup> <sup>p</sup> kL

We want to investigate how the energies of the electronic bound states will be affected, compared to the situation when in the infinite well there is only one slab, with effective electron mass m<sup>1</sup> or m2: As

$$E = \frac{\hbar^2 k\_1^2}{2m\_1} = \frac{\hbar^2 k\_2^2}{2m\_2} \tag{58}$$

we have, with (57):

$$k\_2 = \sqrt{\beta}k, \ k = k\_1\tag{59}$$

The electronic wave function is obtained solving the Schrodinger equation, as in the case of a finite well, studied in Section 2:

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods http://dx.doi.org/10.5772/intechopen.73837 61

$$\psi(\mathbf{x}) = \begin{cases} A\_1 \sin\left(k\_1 \mathbf{x} + \varphi\_1\right), & -L/2 < \mathbf{x} < \mathbf{0} \\ A\_2 \sin\left(\sqrt{\beta} k\_1 \mathbf{x} + \varphi\_2\right), & -L/2 < \mathbf{x} < \mathbf{0} \end{cases} \tag{60}$$

The boundary conditions for the wave function give:

$$\sin\left(-\frac{kL}{2} + \varphi\_1\right) = 0\tag{61}$$

$$\sin\left(\sqrt{\beta}\frac{kL}{2} + \varphi\_2\right) = 0\tag{62}$$

and the continuity in the origin:

$$A\_1 \sin \varphi\_1 = A\_2 \sin \varphi\_2 \tag{63}$$

The BenDaniel-Duke boundary condition means:

$$\frac{1}{m\_1}\psi'(0\_-) = \frac{1}{m\_2}\psi'(0\_+)\tag{64}$$

or:

(56)

(58)

point in the host's Brillouin zone. A popular example is given by the lowest conduction states of GaAs-GaAlAs heterostructures with GaAs layer (typically, its thickness is about 100 Å or larger). A detailed description of the cases in which the envelope function model is successful

As there are some similarities between QWs and QDs, our results are also relevant for these devices. The simplest remark is that the eigenvalues equations for the first odd state in a QW are identical to that corresponding to the l ¼ 0 state in a QD (see for instance [13], problem 63). Also, the eigenvalue equations for the wave vectors of the energy levels for a finite barrier rectangular shaped QD, Eq. (36) in [21], are quite similar to ours—(29) and (30), but somewhat more complicated. The ground state energy of electrons and holes in a core/shell QD is given by the Eq. (21) of [22], an equation similar to ours, just mentioned previously. Such results are important, inter alia, for the interpretation of photoluminescence spectra and photon harvesting of QDs.

Let us consider an infinite 1D square well, delimited by two rigid walls situated in �L=2, respectively L=2, containing two semiconductor slabs, of equal width, but of different materials. It is a toy model for a Janus nanorod, composed of two different semiconductors, with large work functions. We preferred to choose this particular case (equal width), in order to avoid too cumbersome mathematical calculations. The electron effective mass is position dependent, like in (2):

m xð Þ¼ <sup>m</sup>1, � <sup>L</sup>=<sup>2</sup> <sup>&</sup>lt; <sup>x</sup> <sup>&</sup>lt; <sup>0</sup>

We want to investigate how the energies of the electronic bound states will be affected, compared to the situation when in the infinite well there is only one slab, with effective

> ¼ ℏ2 k 2 2 2m<sup>2</sup>

<sup>E</sup> <sup>¼</sup> <sup>ℏ</sup><sup>2</sup> k2 1 2m<sup>1</sup>

<sup>k</sup><sup>2</sup> <sup>¼</sup> ffiffiffi

The electronic wave function is obtained solving the Schrodinger equation, as in the case of a

�

m2, 0 < x < L=2

m<sup>2</sup> ¼ βm<sup>1</sup> (57)

<sup>β</sup> <sup>p</sup> k, k <sup>¼</sup> <sup>k</sup><sup>1</sup> (59)

4. The infinite square well with two semiconductor slabs

is given in [5], p. 66 (See Figure 2).

60 Semiconductors - Growth and Characterization

4.1. The symmetric case

electron mass m<sup>1</sup> or m2: As

finite well, studied in Section 2:

we have, with (57):

with:

$$A\sqrt{\beta}A\_1\cos\varphi\_1 = A\_2\cos\varphi\_2\tag{65}$$

Together with the orthonormality condition for the wave function, Eqs. (61)–(63) and (65) form a system of five equations for five quantities, k, φ1, φ2, A1, A2: As the amplitudes are not of primary interest, we can combine (64) and (65) to obtain:

$$\frac{1}{\sqrt{\beta}}\tan\varphi\_1 = \tan\varphi\_2\tag{66}$$

Replacing in (66), the values of φ1, φ<sup>2</sup> obtained from (61) and (62):

$$
\varphi\_1 = n\_1 \pi + \frac{\text{kL}}{2}, \quad \varphi\_2 = n\_2 \pi - \sqrt{\beta} \frac{\text{kL}}{2} \tag{67}
$$

we get:

$$\tan\frac{kL}{2} = -\sqrt{\beta}\tan\sqrt{\beta}\frac{kL}{2}\tag{68}$$

With

$$\frac{kL}{2} = K \tag{69}$$

it can be written as:

$$
\tan K + \sqrt{\beta} \tan \sqrt{\beta} K = 0 \tag{70}
$$

4.2. The asymmetric case

the relations:

noticing that:

value equation:

with the notations:

function extends outside the wells.

Defining the wave vector k<sup>0</sup> by:

Let us consider now the case of a rectangular infinite asymmetric well, with the potential:

with V<sup>0</sup> > 0, containing, as in the previous example, two semiconductor slabs. It is also a toy model of a Janus nanorod, somewhat more realistic than that discussed in Section 4.1. We also chose a particular geometry (the same width for each slab) to avoid irrelevant mathematical complications. For an electronic bound state of energy E > V0, the wave vectors (and the electronic effective masses) are different in different slabs, according to

; E � <sup>V</sup><sup>0</sup> <sup>¼</sup> <sup>ℏ</sup><sup>2</sup>k<sup>2</sup>

<sup>V</sup><sup>0</sup> <sup>¼</sup> <sup>ℏ</sup><sup>2</sup> k 2 0 2m<sup>1</sup>

> <sup>1</sup> � k 2 2

> > ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi β K<sup>2</sup> <sup>1</sup> � <sup>K</sup><sup>2</sup> 0

<sup>2</sup> , K<sup>0</sup> <sup>¼</sup> <sup>k</sup>0<sup>L</sup>

and following exactly the same steps as in the symmetric case, we obtain the following eigen-

ffiffiffi β p ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi K2 <sup>1</sup> � <sup>K</sup><sup>2</sup> 0 <sup>q</sup> tan

<sup>2</sup> , K<sup>2</sup> <sup>¼</sup> <sup>k</sup>2<sup>L</sup>

If 0 < E < V0, k<sup>2</sup> (and, evidently, K2) become imaginary, and tan K<sup>2</sup> ! �itanhK2: The Eq. (82) and its hyperbolic counterpart are much complicated than (68); even if there are some methods of obtaining approximate analytical solutions, they will be not discussed here. The case of a finite asymmetric well, with two different semiconductor slices, can be studied following exactly the same approach, but now the complications are even more serious, as the wave

k 2 <sup>2</sup> <sup>¼</sup> <sup>β</sup> <sup>k</sup><sup>2</sup>

2 2m<sup>2</sup>

, m<sup>2</sup> ¼ βm<sup>1</sup> (79)

http://dx.doi.org/10.5772/intechopen.73837

� � (81)

<sup>q</sup> � � <sup>¼</sup> <sup>0</sup> (82)

<sup>2</sup> (83)

∞, x < �L=2 0, � L=2 < x < 0 V0, 0 < x < L=2 ∞, x > L=2

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

(78)

63

(80)

V xð Þ¼

<sup>E</sup> <sup>¼</sup> <sup>ℏ</sup><sup>2</sup> k 2 1 2m<sup>1</sup>

1 K1

tan K<sup>1</sup> þ

<sup>K</sup><sup>1</sup> <sup>¼</sup> <sup>k</sup>1<sup>L</sup>

8 >>><

>>>:

If ffiffiffi <sup>β</sup> <sup>p</sup> <sup>¼</sup> <sup>1</sup>, (67) gives:

$$2\tan\frac{kL}{2} = n\pi \to k\_n = \frac{n\pi}{L} \tag{71}$$

$$k\_1 = \frac{\pi}{L} \to K\_1 = \frac{\pi}{2} \tag{72}$$

so the solutions corresponding to the infinite well with an homogenous medium inside the walls.

Eq. (70) is a transcendental one, and its solutions cannot be expressed as a finite combination of elementary functions. A quite popular analytical approximation for the tangent function has been proposed by de Alcantara-Bonfim [8] and generalized by the present author [9]:

$$
\tan x \simeq \frac{0.45\pi(x - n\pi)}{2x - (2n - 1)\pi}, \quad \left(n - \frac{1}{2}\right)\pi < x < n\pi \tag{73}
$$

$$
\tan x \simeq \frac{0.45\pi(x - n\pi)}{(2n + 1)\pi - 2\pi}, \quad n\pi < x < \left(n + \frac{1}{2}\right)\pi \tag{74}
$$

In order to see how this approximation works, let us consider the first two roots of Eq. (68), if β≲ 1: For β ¼ 0:9, we obtain (for instance, using the command FindRoot in Mathematica) K1,exact ¼ 1:65804 and K2,exact ¼ 3:29797, close to π=2, respectively π, that is, to the values corresponding to β ¼ 0: We shall discuss the case of the second root of Eq. (70). As <sup>K</sup>2,exact <sup>¼</sup> <sup>3</sup>:29797<sup>∈</sup> <sup>π</sup>; <sup>3</sup> <sup>2</sup> <sup>π</sup> � � and <sup>K</sup>2,exact ffiffiffi <sup>β</sup> <sup>p</sup> <sup>¼</sup> <sup>2</sup>:9682<sup>∈</sup> <sup>π</sup> <sup>2</sup> ; <sup>π</sup> � �, the two tangent functions appearing in Eq. (70) will be approximated by the two variants of Eqs. (73) and (74), the result being the following:

$$\frac{K - \pi}{3\pi - 2K} + \frac{\sqrt{\beta} \{ \sqrt{\beta}K - \pi \}}{2\sqrt{\beta}K - \pi} = 0\tag{75}$$

So, K can be obtained as a root of a second order equation, namely:

$$K(\beta) = \frac{\pi}{4} \frac{-3\beta + 1 + \sqrt{9\beta^2 - 24\beta^2 + 26\beta - 8\beta^{1/2} + 1}}{\beta^{1/2} \left(1 - \beta^{1/2}\right)} \tag{76}$$

We find that <sup>K</sup> <sup>β</sup> <sup>¼</sup> <sup>0</sup>:<sup>9</sup> � � <sup>¼</sup> <sup>K</sup>2, approx <sup>¼</sup> <sup>3</sup>:2987, so quite close to the exact value, the error being:

$$\frac{K\_{2,exact} - K\_{2, approx}}{K\_{2,exact}} = \frac{3.29797 - 3.2987}{3.29797} = -2.2135 \times 10^{-4} \tag{77}$$

However, due to the rapid variation of the tangent functions near its singularities, this approximation method must be used with utmost care, as it can easily give unacceptable results (this is the case of the first root, for β ¼ 0:9).

#### 4.2. The asymmetric case

tan <sup>K</sup> <sup>þ</sup> ffiffiffi

2 tan kL

<sup>k</sup><sup>1</sup> <sup>¼</sup> <sup>π</sup>

been proposed by de Alcantara-Bonfim [8] and generalized by the present author [9]:

tan <sup>x</sup> <sup>≃</sup> <sup>0</sup>:45πð Þ <sup>x</sup> � <sup>n</sup><sup>π</sup> 2x � ð Þ 2n � 1 π

tan <sup>x</sup> <sup>≃</sup> <sup>0</sup>:45πð Þ <sup>x</sup> � <sup>n</sup><sup>π</sup> ð Þ 2n þ 1 π � 2x

> K � π <sup>3</sup><sup>π</sup> � <sup>2</sup><sup>K</sup> <sup>þ</sup>

�3β þ 1 þ

So, K can be obtained as a root of a second order equation, namely:

K β � � <sup>¼</sup> <sup>π</sup> 4

the case of the first root, for β ¼ 0:9).

K2,exact � K2, approx K2,exact

<sup>2</sup> <sup>π</sup> � � and <sup>K</sup>2,exact

If ffiffiffi

<sup>β</sup> <sup>p</sup> <sup>¼</sup> <sup>1</sup>, (67) gives:

62 Semiconductors - Growth and Characterization

<sup>K</sup>2,exact <sup>¼</sup> <sup>3</sup>:29797<sup>∈</sup> <sup>π</sup>; <sup>3</sup>

being the following:

β p tan ffiffiffi

<sup>2</sup> <sup>¼</sup> <sup>n</sup><sup>π</sup> ! kn <sup>¼</sup> <sup>n</sup><sup>π</sup>

<sup>L</sup> ! <sup>K</sup><sup>1</sup> <sup>¼</sup> <sup>π</sup>

, n � <sup>1</sup>

2 � �

, nπ < x < n þ

ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi

<sup>2</sup> <sup>þ</sup> <sup>26</sup><sup>β</sup> � <sup>8</sup>β<sup>1</sup>=<sup>2</sup> <sup>þ</sup> <sup>1</sup>

3

1 2 � �

so the solutions corresponding to the infinite well with an homogenous medium inside the walls. Eq. (70) is a transcendental one, and its solutions cannot be expressed as a finite combination of elementary functions. A quite popular analytical approximation for the tangent function has

In order to see how this approximation works, let us consider the first two roots of Eq. (68), if β≲ 1: For β ¼ 0:9, we obtain (for instance, using the command FindRoot in Mathematica) K1,exact ¼ 1:65804 and K2,exact ¼ 3:29797, close to π=2, respectively π, that is, to the values corresponding to β ¼ 0: We shall discuss the case of the second root of Eq. (70). As

<sup>β</sup> <sup>p</sup> <sup>¼</sup> <sup>2</sup>:9682<sup>∈</sup> <sup>π</sup>

<sup>9</sup>β<sup>2</sup> � <sup>24</sup><sup>β</sup>

We find that <sup>K</sup> <sup>β</sup> <sup>¼</sup> <sup>0</sup>:<sup>9</sup> � � <sup>¼</sup> <sup>K</sup>2, approx <sup>¼</sup> <sup>3</sup>:2987, so quite close to the exact value, the error being:

<sup>¼</sup> <sup>3</sup>:<sup>29797</sup> � <sup>3</sup>:<sup>2987</sup>

However, due to the rapid variation of the tangent functions near its singularities, this approximation method must be used with utmost care, as it can easily give unacceptable results (this is

appearing in Eq. (70) will be approximated by the two variants of Eqs. (73) and (74), the result

ffiffiffi <sup>β</sup> p ffiffiffi<sup>β</sup> <sup>p</sup> <sup>K</sup> � <sup>π</sup> � � 2 ffiffiffi

ffiffiffi

q

<sup>β</sup> <sup>p</sup> <sup>K</sup> <sup>¼</sup> <sup>0</sup> (70)

<sup>L</sup> (71)

π < x < nπ (73)

<sup>2</sup> ; <sup>π</sup> � �, the two tangent functions

<sup>β</sup> <sup>p</sup> <sup>K</sup> � <sup>π</sup> <sup>¼</sup> <sup>0</sup> (75)

<sup>β</sup><sup>1</sup>=<sup>2</sup> <sup>1</sup> � <sup>β</sup><sup>1</sup>=<sup>2</sup> � � (76)

<sup>3</sup>:<sup>29797</sup> ¼ �2:<sup>2135</sup> � <sup>10</sup>�<sup>4</sup> (77)

π (74)

<sup>2</sup> (72)

Let us consider now the case of a rectangular infinite asymmetric well, with the potential:

$$V(\mathbf{x}) = \begin{cases} \ast \ast & \mathbf{x} < -\mathbf{L}/2 \\ 0, & -\mathbf{L}/2 < \mathbf{x} < 0 \\ V\_{0\prime} & 0 < \mathbf{x} < \mathbf{L}/2 \\ \ast & \mathbf{x} > \mathbf{L}/2 \end{cases} \tag{78}$$

with V<sup>0</sup> > 0, containing, as in the previous example, two semiconductor slabs. It is also a toy model of a Janus nanorod, somewhat more realistic than that discussed in Section 4.1. We also chose a particular geometry (the same width for each slab) to avoid irrelevant mathematical complications. For an electronic bound state of energy E > V0, the wave vectors (and the electronic effective masses) are different in different slabs, according to the relations:

$$E = \frac{\hbar^2 k\_1^2}{2m\_1};\ E - V\_0 = \frac{\hbar^2 k\_2^2}{2m\_2},\ m\_2 = \beta m\_1 \tag{79}$$

Defining the wave vector k<sup>0</sup> by:

$$V\_0 = \frac{\hbar^2 k\_0^2}{2m\_1} \tag{80}$$

noticing that:

$$k\_2^2 = \beta \left(k\_1^2 - k\_2^2\right) \tag{81}$$

and following exactly the same steps as in the symmetric case, we obtain the following eigenvalue equation:

$$\frac{1}{K\_1}\tan K\_1 + \frac{\sqrt{\beta}}{\sqrt{K\_1^2 - K\_0^2}}\tan\sqrt{\beta(K\_1^2 - K\_0^2)} = 0\tag{82}$$

with the notations:

$$K\_1 = \frac{k\_1 L}{2}, \quad K\_2 = \frac{k\_2 L}{2}, \quad K\_0 = \frac{k\_0 L}{2} \tag{83}$$

If 0 < E < V0, k<sup>2</sup> (and, evidently, K2) become imaginary, and tan K<sup>2</sup> ! �itanhK2: The Eq. (82) and its hyperbolic counterpart are much complicated than (68); even if there are some methods of obtaining approximate analytical solutions, they will be not discussed here. The case of a finite asymmetric well, with two different semiconductor slices, can be studied following exactly the same approach, but now the complications are even more serious, as the wave function extends outside the wells.

#### 5. Conclusions

In this chapter, we obtained approximate analytical solutions for the eigenvalue equation of the first two bound states in a semiconductor quantum well, in a particular case of positiondependent mass of the charge carrier—in fact, the simplest one, corresponding to BenDaniel-Duke boundary conditions. This position dependence can be characterized by β, the ratio of the mass inside, to the mass outside the well. Actually, we obtained quite simple expressions for the dimensionless wave vector, in terms of the potential strength and of β: Even if we solved this problem in terms of one-particle quantum mechanics, obtaining the wave function and the eigenvalues of the bound states, our results can be directly applied in the theory of envelope functions in the conduction band at heterointerfaces. Our approach is based on the "algebraization" of trigonometric functions present in the transcendental eigenvalue equations; in this way, they are transformed in tractable algebraic equations.

[2] Lévai G, Özer O. An exactly solvable Schrödinger equation with finite positive positiondependenteffective mass. Journal of Mathematical Physics. 2010;51:092103(13 pp). DOI:

Semiconductor Quantum Wells with BenDaniel-Duke Boundary Conditions and Janus Nanorods

http://dx.doi.org/10.5772/intechopen.73837

65

[3] Nikitin AG, Zasadko TM. Superintegrable systems with position dependent mass. Jour-

[4] Sebawe Abdalla M, Eleuch H. Exact solutions of the position-dependent-effective mass Schrödinger equation. AIP Advances. 2016;6:055011(7 pp). DOI: 10.1063/1.4949567

[5] Bastard G. Wave mechanics applied to semiconductor heterostructures. Les Editions de

[7] Barsan V. New applications of the Lambert and generalized Lambert functions in ferro-

[8] de Alcantara Bonfim OF, Griffiths DJ. Exact and approximate energy spectrum for the finite square well and related potentials. American Journal of Physics. 2006;74:43-48. DOI:

[9] Barsan V. Algebraic approximations for transcendental equations with applications in nanophysics. Philosophical Magazine. 2015;95:3023-3038. DOI: 10.1080/14786435.2015.

[10] Song Y, Chen S. Janus nanoparticles: Preparation, characterization, and applications.

[11] Singh VA, Kumar L. Revisiting elementary quantum mechanics with the Daniel—Duke boundary conditions. American Journal of Physics. 2006;74:412-418. DOI: 10.1119/1.2174

[12] Singh S, Pathak P, Singh VA. Approximate approaches to the one-dimensional finite potential well. European Journal of Physics. 2011;32:1701-1710. DOI: 10.1088/0143-0807/

[13] Fluegge S. Practical Quantum Mechanics (I). Berlin: Springer-Verlag; 1971. 333 p. DOI:

[14] Pitkanen PH. Rectangular potential well problem in quantum mechanics. American Jour-

[15] Barker BI, Rayborn GH, Ioup JJ, Ioup GE. Approximating the finite square well with an infinite well: Energies and eigenfunctions. American Journal of Physics. 1991;59:10381042

[16] Barsan V, Dragomir R. A new approximation for the square well problem. Optoelectron-

[17] Barsan V. A new analytic approximation for the energy eigenvalues of a finite square well.

nal of Physics. 1955;23:111-113. DOI: 10.1119/1.1933912(1955)

ics and Advanced Materials. 2012;6:917-925

Romanian Reports in Physics. 2012;64:685-694

Chemistry, an Asian Journal. 2014;9:418-424. DOI: 10.1002/asia.201301398

[6] Ihn T. Semiconductor Nanostructures. Oxford: Oxford University Press; 2011. 552 p

magnetism and quantum mechanics. arXiv:1611.01014v2. 2016

nal of Mathematical Physics. 2015;56:042101(13 pp). DOI: 10.1063/1.4908107

10.1063/1.3483716

Physique. 1990. 366 p

1081425

031

32/6/023

10.1007/978-3-642-61995-3

https://doi.org/10.1119/1.2140771

We also proposed two models for a semiconductor Janus nanorod—a system, which was not yet treated analytically.

Our results can be easily extended to more realistic (e.g., linear) position dependence of the mass carrier and to other nanosystems. For instance, the eigenvalue equations for the wave vectors of bound energy levels of a finite barrier rectangular-shaped quantum dot, Eq. (36) in [21], are quite similar to ours—(22), (23), but somewhat more complicated. The ground state energy of electrons and holes in a core/shell quantum dot is given by Eq. (21) of [22], an equation similar to ours, just mentioned previously. Such results are important, inter alia, for the interpretation of photoluminescence spectra of heterojunctions.

#### Acknowledgements

The author acknowledges the financial support of the IFIN-HH-ANCSI project PN 16 42 01 01/ 2016 and of the IFIN-HH-JINR grant 04-4-1121-2015/2017.

#### Author details

Victor Barsan

Address all correspondence to: vbarsan@theory.nipne.ro

National Institute for Physics and Nuclear Engineering, Bucharest-Magurele, Romania

#### References

[1] von Roos O. Position-dependent effective masses in semiconductor theory. Physical Review. 1983;B27:7547-7559. DOI: 10.1103/PhysRevB.27.7547

[2] Lévai G, Özer O. An exactly solvable Schrödinger equation with finite positive positiondependenteffective mass. Journal of Mathematical Physics. 2010;51:092103(13 pp). DOI: 10.1063/1.3483716

5. Conclusions

64 Semiconductors - Growth and Characterization

yet treated analytically.

Acknowledgements

Author details

Victor Barsan

References

In this chapter, we obtained approximate analytical solutions for the eigenvalue equation of the first two bound states in a semiconductor quantum well, in a particular case of positiondependent mass of the charge carrier—in fact, the simplest one, corresponding to BenDaniel-Duke boundary conditions. This position dependence can be characterized by β, the ratio of the mass inside, to the mass outside the well. Actually, we obtained quite simple expressions for the dimensionless wave vector, in terms of the potential strength and of β: Even if we solved this problem in terms of one-particle quantum mechanics, obtaining the wave function and the eigenvalues of the bound states, our results can be directly applied in the theory of envelope functions in the conduction band at heterointerfaces. Our approach is based on the "algebraization" of trigonometric functions present in the transcendental eigenvalue equa-

We also proposed two models for a semiconductor Janus nanorod—a system, which was not

Our results can be easily extended to more realistic (e.g., linear) position dependence of the mass carrier and to other nanosystems. For instance, the eigenvalue equations for the wave vectors of bound energy levels of a finite barrier rectangular-shaped quantum dot, Eq. (36) in [21], are quite similar to ours—(22), (23), but somewhat more complicated. The ground state energy of electrons and holes in a core/shell quantum dot is given by Eq. (21) of [22], an equation similar to ours, just mentioned previously. Such results are important, inter alia, for

The author acknowledges the financial support of the IFIN-HH-ANCSI project PN 16 42 01 01/

National Institute for Physics and Nuclear Engineering, Bucharest-Magurele, Romania

Review. 1983;B27:7547-7559. DOI: 10.1103/PhysRevB.27.7547

[1] von Roos O. Position-dependent effective masses in semiconductor theory. Physical

tions; in this way, they are transformed in tractable algebraic equations.

the interpretation of photoluminescence spectra of heterojunctions.

2016 and of the IFIN-HH-JINR grant 04-4-1121-2015/2017.

Address all correspondence to: vbarsan@theory.nipne.ro


[18] Barsan V, Ciornei M-C. Semiconductor quantum wells with BenDaniel-Duke boundary conditions: Approximate analytical results. European Journal of Physics. 2017;38:015407 (22 pp). DOI: 10.1088/0143-0807/38/1/015407

**Section 2**

**Semiconductors for Energy Applications**


**Semiconductors for Energy Applications**

[18] Barsan V, Ciornei M-C. Semiconductor quantum wells with BenDaniel-Duke boundary conditions: Approximate analytical results. European Journal of Physics. 2017;38:015407

[19] Weisstein EW. "Cubic Formula." From MathWorld—A Wolfram Web Resource. http://

[20] Biswas D, Kumar S, Das T. Interdiffusion induced changes in the photoluminescence of InxGa<sup>1</sup>xAs=GaAs quantum dots interpreted. Journal of Applied Physics. 2007;101:026108

[21] Ata E, Demirhan D, Buyukkilic F. 2-d finite barrier rectangular quantum dots: Schroedinger

[22] Ibral A, Zouitine A, Assaid EM, Achouby HE. Polarization effects on spectra of spherical core/shell nanostructures: Perturbation theory against finite difference approach. Physica.

description. Physica. 2014;E62:71-75. DOI: 10.1016/j.physe.2014.11.011

(22 pp). DOI: 10.1088/0143-0807/38/1/015407

mathworld.wolfram.com/CubicFormula.html

2015;B458:73-84. DOI: 10.1016/j.physb.2014.11.009

(3 pp). DOI: 10.1063/1.2430510

66 Semiconductors - Growth and Characterization

**Chapter 4**

**Provisional chapter**

**Te Thin Films** 

= 5 ml.

**Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films**

The paper presents the investigation on the influence of substrate temperature *Ts* and the sprayed initial solution volume *Vs* on structural, substructural, optical properties, and

phase nanocrystalline ZnO films with average crystallite size of *DC* = 25–270 nm and thickness of *d* = 0.8–1.2 μm can be deposited at substrate temperatures of *Ts* > 473 K. The continuous CZTS films with optimal thickness (*d* = 1.3 μm) for application as absorber layers in solar cells were deposited at the sprayed initial precursor volume of *Vs*

The increase of the substrate temperature up to 673 K caused the significant improvements in the stoichiometry of ZnO films. The optimal stoichiometry ratio of CZTS films for application in solar cells was obtained at *Vs* = 3–4 ml. Optical study of ZnO films showed that these films have a high-transmission coefficient values of *T* = 60–80%. To the best of our knowledge, there is the lack of works devoted to the study of CZT films

**ZnSnS4, Cd1−xZn**

**x**

ZnSnS<sup>4</sup> (CZTS) films as well as state-of-the-art of

Te (CZT) films obtained by spray pyrolysis technique. The single-

Te, thin films, pulsed spray pyrolysis

DOI: 10.5772/intechopen.72988

© 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,

© 2018 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.

and reproduction in any medium, provided the original work is properly cited.

ZnO is an *n*-type direct-gap semiconductor with a wide band gap (*Eg* = 3.37 eV at *T* = 300 K) having the highest value of exciton energy (60 meV) among the binary compounds [1]. This material is a perspective for application in microelectronics, nanoelectronics, optoelectronics, sensors, solar cells among others due to its unique physical, electrical, and optical properties, non-toxic nature, and chemical and thermal stability in the ambient

**Obtained by Spray Pyrolysis Method**

**Obtained by Spray Pyrolysis Method**

Oleksandr Dobrozhan, Denys Kurbatov, Petro Danilchenko and Anatoliy Opanasyuk

Oleksandr Dobrozhan, Denys Kurbatov, Petro Danilchenko and Anatoliy Opanasyuk

http://dx.doi.org/10.5772/intechopen.72988

**Abstract**

studying the Cd1−xZnx

**Keywords:** ZnO, Cu<sup>2</sup>

**1. Introduction**

**Nanostructured ZnO, Cu<sup>2</sup>**

Additional information is available at the end of the chapter

elemental composition of ZnO and Cu<sup>2</sup>

obtained by spray pyrolysis technique.

ZnSnS<sup>4</sup>

, Cd1−xZnx

Additional information is available at the end of the chapter

#### **Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method Nanostructured ZnO, Cu<sup>2</sup> ZnSnS4, Cd1−xZn x Te Thin Films Obtained by Spray Pyrolysis Method**

DOI: 10.5772/intechopen.72988

Oleksandr Dobrozhan, Denys Kurbatov, Petro Danilchenko and Anatoliy Opanasyuk Oleksandr Dobrozhan, Denys Kurbatov, Petro Danilchenko and Anatoliy Opanasyuk

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/intechopen.72988

#### **Abstract**

The paper presents the investigation on the influence of substrate temperature *Ts* and the sprayed initial solution volume *Vs* on structural, substructural, optical properties, and elemental composition of ZnO and Cu<sup>2</sup> ZnSnS<sup>4</sup> (CZTS) films as well as state-of-the-art of studying the Cd1−xZnx Te (CZT) films obtained by spray pyrolysis technique. The singlephase nanocrystalline ZnO films with average crystallite size of *DC* = 25–270 nm and thickness of *d* = 0.8–1.2 μm can be deposited at substrate temperatures of *Ts* > 473 K. The continuous CZTS films with optimal thickness (*d* = 1.3 μm) for application as absorber layers in solar cells were deposited at the sprayed initial precursor volume of *Vs* = 5 ml. The increase of the substrate temperature up to 673 K caused the significant improvements in the stoichiometry of ZnO films. The optimal stoichiometry ratio of CZTS films for application in solar cells was obtained at *Vs* = 3–4 ml. Optical study of ZnO films showed that these films have a high-transmission coefficient values of *T* = 60–80%. To the best of our knowledge, there is the lack of works devoted to the study of CZT films obtained by spray pyrolysis technique.

**Keywords:** ZnO, Cu<sup>2</sup> ZnSnS<sup>4</sup> , Cd1−xZnx Te, thin films, pulsed spray pyrolysis

## **1. Introduction**

ZnO is an *n*-type direct-gap semiconductor with a wide band gap (*Eg* = 3.37 eV at *T* = 300 K) having the highest value of exciton energy (60 meV) among the binary compounds [1]. This material is a perspective for application in microelectronics, nanoelectronics, optoelectronics, sensors, solar cells among others due to its unique physical, electrical, and optical properties, non-toxic nature, and chemical and thermal stability in the ambient

and reproduction in any medium, provided the original work is properly cited.

atmosphere [2]. It should be noted that, at present, Ukrainian sector of renewable energy, in particular solar, is developing rapidly. First of all it is made possible, thanks to the government support policy. In turn, it leads to an increased interest in the development of new solar cell designs to create further production of solar modules higher efficiency [3, 4]. Due to the absence of rare and toxic elements in zinc oxide compound and possibility to apply low-cost deposition techniques, this material may be an alternative to the traditional ITO ((In<sup>2</sup> O3 )0.9-(SnO<sup>2</sup> )0.1) and FTO (SnO<sup>2</sup> :F)) transparent conductive layers in thin-film solar cells (SCs) and another optoelectronic device [5]. Nowadays, the perspective substitution of the traditional Si, CdTe, Сu(In,Ga)(S,Se)<sup>2</sup> absorption layers in thin-film SCs is considered Cu<sup>2</sup> ZnSnS<sup>4</sup> (CZTS) semiconductor compound which has the optimal optical properties (*Eg* = 1.5 eV, *α* ~ 104 –105 cm−1) [6].

spray pyrolysis [5, 8–10]) techniques. Typically, the physical methods allow to obtain more perfect films with a higher structural quality, and these methods provide a precise control of the films thickness and low content of defects in deposited material compare to chemical methods, but physical deposition techniques require the usage of more complicated equipment and presence of high level of vacuum, thus they are energy-consuming. In contrary, chemical techniques to deposit ZnO, CZTS, and CZT films are low-cost and energy savers. Among them, spray pyrolysis method is considered as the most promising technique. This technique is simple and non-vacuum providing the deposition of the continuous, porous,

Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method

http://dx.doi.org/10.5772/intechopen.72988

71

Taking into account the increased interest to the nanosized materials with properties, significantly different to bulk materials (caused by quantum-size effect), several scientific groups have obtained the nanocrystalline ZnO and CZTS films [16, 17]. It is important to note that the works dedicated to the study of the nanosized structures used chemical techniques for films deposition. ZnO, CZTS, and CZT films deposited by spray pyrolysis technique are not yet

The image of laboratory setup developed for the deposition of ZnO, CZTS, and CZT films by pulsed spray pyrolysis is showed in **Figure 1**. It consists of a spraying gun with initial precursor volume reservoir (1), spraying nozzle (2), and microcontroller block (3), allowing the control of the number of spraying cycles, time, and pauses between cycles. To the spraying gun, the compressor with pressure regulator (4) is connected with the aim of producing the air flow for transportation of the dispersed precursor onto heated substrate surface. Between the spraying gun and the compressor, an electromagnetic valve (5) is installed, where the "open" and "closed" regimes are controlled by the microcontroller block (3). The heating of substrate (6) is provided by the heating plate (7). During the deposition of films by spray pyrolysis technique, the properties of ZnO, CZTS, and CZT condensates are dependent on the precursor choice and physical, chemical deposition conditions. **Table 1** presents the overview of deposition conditions and precursors typically used to deposit the ZnO, CZTS, and CZT

**Figure 1.** Image of the experimental setup for ZnO, CZTS, and CZT films deposition by pulsed spray pyrolysis: (1) spraying gun with initial precursor volume reservoir, (2) spraying nozzle, (3) microcontroller block, (4) compressor, (5)

nanostructured films, and multilayered structures [15].

well-studied; this fact conditioned the aim of our study.

films by spray pyrolysis technique.

electromagnetic valve, (6) substrate, and (7) heating plate [18, 19].

Cd1–xZnx Te (CZT) solid solutions are perspective alternative absorption materials to Сu(In,Ga) (S,Se)<sup>2</sup> in the tandem solar cells having the band gap value of *Eg* = 1.1 eV. The appealing advantage of CZT compound is the variation of band gap by changing the zinc concentration. The optimal CZT solid solution with *Eg* ~1.7 eV can be obtained at the chemical composition of x ~ 0.2 [7]. To achieve the best working characteristics of devices, ZnO, CZTS, and CZT films must have the single-phase structure with large coherent domain sizes (CDS) *L*, low levels of microdeformations *ε*, microstresses *σ*, dislocation concentrations *ρ*, and well-controlled elemental composition. Unfortunately, typically these films have a high level of defects and secondary phases with different band gaps worsening the performance of the devices based on them. ZnO and CZT films in solar cells should possess high-transmission coefficients and controlled band gaps. Moreover, in order to improve the structural and optical properties of films for application in the low-cost optoelectronic devices, ZnO, CZTS, and CZT films should be deposited by lowcost, non-vacuum methods with optimized physical and technological deposition conditions.

Among the methods to deposit the ZnO, CZTS, and CZT films, special attention is paid to the spray pyrolysis technique having unique advantages: simplicity, efficiency, and cheapness. This technique provides the non-vacuum deposition of a large-area thin film with wellcontrolled properties.

It was shown [8, 9] that the greatest influence on physical properties and elemental composition of ZnO film has a substrate temperature *Ts* , CZTS film—the sprayed initial precursor volume *Vs* . It should be noted that until now CZT films deposited by spray pyrolysis technique are not well-studied, except some works [10, 11].

Thus, the investigation of the influence of deposition conditions on structural, substructural, and optical properties of ZnO, CZTS, and CZT films deposited by spray pyrolysis technique is the perspective in terms of its application in highly efficient optoelectronic devices.

## **2. ZnO, CZTS, and CZT thin films deposition methods. Peculiarities of the spray pyrolysis technique**

The wide range of methods is well developed to deposit ZnO, CZT, and CZTS films which split into physical (for example, magnetron sputtering [12–14]) and chemical (for example, spray pyrolysis [5, 8–10]) techniques. Typically, the physical methods allow to obtain more perfect films with a higher structural quality, and these methods provide a precise control of the films thickness and low content of defects in deposited material compare to chemical methods, but physical deposition techniques require the usage of more complicated equipment and presence of high level of vacuum, thus they are energy-consuming. In contrary, chemical techniques to deposit ZnO, CZTS, and CZT films are low-cost and energy savers. Among them, spray pyrolysis method is considered as the most promising technique. This technique is simple and non-vacuum providing the deposition of the continuous, porous, nanostructured films, and multilayered structures [15].

atmosphere [2]. It should be noted that, at present, Ukrainian sector of renewable energy, in particular solar, is developing rapidly. First of all it is made possible, thanks to the government support policy. In turn, it leads to an increased interest in the development of new solar cell designs to create further production of solar modules higher efficiency [3, 4]. Due to the absence of rare and toxic elements in zinc oxide compound and possibility to apply low-cost deposition techniques, this material may be an alternative to the traditional

cells (SCs) and another optoelectronic device [5]. Nowadays, the perspective substitution

(CZTS) semiconductor compound which has the optimal optical properties

Te (CZT) solid solutions are perspective alternative absorption materials to Сu(In,Ga)

 in the tandem solar cells having the band gap value of *Eg* = 1.1 eV. The appealing advantage of CZT compound is the variation of band gap by changing the zinc concentration. The

x ~ 0.2 [7]. To achieve the best working characteristics of devices, ZnO, CZTS, and CZT films must have the single-phase structure with large coherent domain sizes (CDS) *L*, low levels of microdeformations *ε*, microstresses *σ*, dislocation concentrations *ρ*, and well-controlled elemental composition. Unfortunately, typically these films have a high level of defects and secondary phases with different band gaps worsening the performance of the devices based on them. ZnO and CZT films in solar cells should possess high-transmission coefficients and controlled band gaps. Moreover, in order to improve the structural and optical properties of films for application in the low-cost optoelectronic devices, ZnO, CZTS, and CZT films should be deposited by lowcost, non-vacuum methods with optimized physical and technological deposition conditions.

Among the methods to deposit the ZnO, CZTS, and CZT films, special attention is paid to the spray pyrolysis technique having unique advantages: simplicity, efficiency, and cheapness. This technique provides the non-vacuum deposition of a large-area thin film with well-

It was shown [8, 9] that the greatest influence on physical properties and elemental composi-

Thus, the investigation of the influence of deposition conditions on structural, substructural, and optical properties of ZnO, CZTS, and CZT films deposited by spray pyrolysis technique

The wide range of methods is well developed to deposit ZnO, CZT, and CZTS films which split into physical (for example, magnetron sputtering [12–14]) and chemical (for example,

is the perspective in terms of its application in highly efficient optoelectronic devices.

**2. ZnO, CZTS, and CZT thin films deposition methods. Peculiarities** 

. It should be noted that until now CZT films deposited by spray pyrolysis technique

:F)) transparent conductive layers in thin-film solar

~1.7 eV can be obtained at the chemical composition of

absorption layers in thin-film SCs is considered

, CZTS film—the sprayed initial precursor vol-

ITO ((In<sup>2</sup>

Cd1–xZnx

(S,Se)<sup>2</sup>

ZnSnS<sup>4</sup>

(*Eg* = 1.5 eV, *α* ~ 104

controlled properties.

ume *Vs*

tion of ZnO film has a substrate temperature *Ts*

are not well-studied, except some works [10, 11].

**of the spray pyrolysis technique**

Cu<sup>2</sup>

O3

)0.9-(SnO<sup>2</sup>

70 Semiconductors - Growth and Characterization

optimal CZT solid solution with *Eg*

of the traditional Si, CdTe, Сu(In,Ga)(S,Se)<sup>2</sup>

–105 cm−1) [6].

)0.1) and FTO (SnO<sup>2</sup>

Taking into account the increased interest to the nanosized materials with properties, significantly different to bulk materials (caused by quantum-size effect), several scientific groups have obtained the nanocrystalline ZnO and CZTS films [16, 17]. It is important to note that the works dedicated to the study of the nanosized structures used chemical techniques for films deposition. ZnO, CZTS, and CZT films deposited by spray pyrolysis technique are not yet well-studied; this fact conditioned the aim of our study.

The image of laboratory setup developed for the deposition of ZnO, CZTS, and CZT films by pulsed spray pyrolysis is showed in **Figure 1**. It consists of a spraying gun with initial precursor volume reservoir (1), spraying nozzle (2), and microcontroller block (3), allowing the control of the number of spraying cycles, time, and pauses between cycles. To the spraying gun, the compressor with pressure regulator (4) is connected with the aim of producing the air flow for transportation of the dispersed precursor onto heated substrate surface. Between the spraying gun and the compressor, an electromagnetic valve (5) is installed, where the "open" and "closed" regimes are controlled by the microcontroller block (3). The heating of substrate (6) is provided by the heating plate (7). During the deposition of films by spray pyrolysis technique, the properties of ZnO, CZTS, and CZT condensates are dependent on the precursor choice and physical, chemical deposition conditions. **Table 1** presents the overview of deposition conditions and precursors typically used to deposit the ZnO, CZTS, and CZT films by spray pyrolysis technique.

**Figure 1.** Image of the experimental setup for ZnO, CZTS, and CZT films deposition by pulsed spray pyrolysis: (1) spraying gun with initial precursor volume reservoir, (2) spraying nozzle, (3) microcontroller block, (4) compressor, (5) electromagnetic valve, (6) substrate, and (7) heating plate [18, 19].


**3. Morphological, structural, and substructural properties of ZnO** 

The surface morphology, structural, substructural, optical properties, and chemical composition of ZnO and CZTS films deposited by spray pyrolysis method are determined by its

Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method

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73

SEM images of ZnO films deposited at different substrate temperatures are presented in **Figure 2a**–**d** [19, 20]. It has been shown that at substrate temperatures higher than 473 K, crack-free and continuous nanocrystalline ZnO films with a good adhesion to substrate were

The average grain size in the condensates was in the range of *DC* = (25–270) nm (see inset in **Figure 2d**), increasing under the increase the deposition temperature up to 673 K. Whereas,

One of the main film parameters of CZTS films is its thickness, which is typically controlled

in the range of *d* = 0.244–0.754 μm was studied by authors [9]. These values are not optimal for absorption of nearly 100% solar radiation because of the necessity CZTS films with thickness

Thus, we have studied the CZTS films deposited by spray pyrolysis technique at different sprayed initial precursor volumes which had the higher thickness than studied in Ref. [9].

In **Figure 2e**–**h**, the SEM images of CZTS films and its cross-section deposited at different *Vs* are presented. It can be seen that in the range of studied values continuous films were formed,

Inset (d) shows the film surface with high resolution obtained at 673 K [19] and CZTS films surface deposited at different

, ml: 2 (e), 3 (f), 4 (g), and 5 (h). Inset (h) shows the films' cross-section [20].

. Dependence CZTS films properties vs. films thickness

, K: 473 (а), 573 (b), 623 (c), and 673 (d).

the film thickness determined by the cross-sectional image was *d* = 0.8–1.2 μm.

**and CZTS films obtained by spray pyrolysis technique**

physical, chemical, and technological deposition conditions.

**Figure 2.** SEM images of ZnO films surface deposited at different temperature *Ts*

dispersed initial precursor volumes *Vs*

**3.1. The morphological properties**

by the dispersed precursor volume *Vs*

formed.

of *d* = 1–3 μm [6].

**Table 1.** Precursors and physical-chemical conditions to deposit the ZnO, CZTS, and CZT films by spray pyrolysis method.

It should be noted that in order to obtain initial molecular solution for the deposition, the typical materials are metal salts dissolved in polar solvents, particularly in water, ethanol, etc. The most common substrates used are the non-oriented glass and silicon slides. The average substrate temperature is in the range of 250–823 K. It should be noted that these values are lower in comparison to the substrate temperatures used in physical methods.

## **3. Morphological, structural, and substructural properties of ZnO and CZTS films obtained by spray pyrolysis technique**

The surface morphology, structural, substructural, optical properties, and chemical composition of ZnO and CZTS films deposited by spray pyrolysis method are determined by its physical, chemical, and technological deposition conditions.

#### **3.1. The morphological properties**

It should be noted that in order to obtain initial molecular solution for the deposition, the typical materials are metal salts dissolved in polar solvents, particularly in water, ethanol, etc. The most common substrates used are the non-oriented glass and silicon slides. The average substrate temperature is in the range of 250–823 K. It should be noted that these values are

**Table 1.** Precursors and physical-chemical conditions to deposit the ZnO, CZTS, and CZT films by spray pyrolysis method.

lower in comparison to the substrate temperatures used in physical methods.

**№ Initial precursor Solvent Concentration** 

) H2

H2

H2

H2

H2 O + CH<sup>3</sup>

H2

H2

H2

(CH3 )2

H2 O + C<sup>2</sup> H5

O 0.02

SO 0.010

OH 0.020

0.010 0.010 0.080

(1:1:3)

0.005 0.005 0.040

ZnO films deposition

1 Zinc chloride (ZnCl<sup>2</sup>

72 Semiconductors - Growth and Characterization

(Zn(CH<sup>3</sup>

Zn(CH<sup>3</sup>

4 Cadmium chloride (CdCl2

5 Copper chloride (CuCl<sup>2</sup> ) Zinc chloride (ZnCl<sup>2</sup> ) Tin chloride (SnCl<sup>2</sup>

6 Copper chloride (CuCl<sup>2</sup> ) Zinc chloride (ZnCl<sup>2</sup> ) Tin chloride (SnCl<sup>2</sup>

(ZnCl<sup>2</sup>

Тhiourea (CS(NH<sup>2</sup> )2 )

Тhiourea (CS(NH<sup>2</sup> )2 )

COO)<sup>2</sup> ∙2H2 O)

COO)<sup>2</sup>

) Zinc chloride

)

)

)

) Tellurium chloride (TeCl<sup>4</sup>

2 Zinc acetate

3 Zinc acetate

CZT films deposition

CZTS films deposition

**(М)**

**Substrate type**

O 0.10 Silicon 623–823 [21]

O 0.10 Glass 773 [22]

O 0.10 Glass 523–723 [23]

O 0.04 Glass 573 [24]

O 0.50 Glass 453–723 [26]

O 0.10 Glass 623 [27]

OH 0.20 Glass 693 [25]

Soda-lime glass

Soda-lime glass

**Substrate temperature,** 

Glass 250–325 [10,

623 [28]

553–633 [29]

**Ref.**

11]

*Ts*  **(K)**

> SEM images of ZnO films deposited at different substrate temperatures are presented in **Figure 2a**–**d** [19, 20]. It has been shown that at substrate temperatures higher than 473 K, crack-free and continuous nanocrystalline ZnO films with a good adhesion to substrate were formed.

> The average grain size in the condensates was in the range of *DC* = (25–270) nm (see inset in **Figure 2d**), increasing under the increase the deposition temperature up to 673 K. Whereas, the film thickness determined by the cross-sectional image was *d* = 0.8–1.2 μm.

> One of the main film parameters of CZTS films is its thickness, which is typically controlled by the dispersed precursor volume *Vs* . Dependence CZTS films properties vs. films thickness in the range of *d* = 0.244–0.754 μm was studied by authors [9]. These values are not optimal for absorption of nearly 100% solar radiation because of the necessity CZTS films with thickness of *d* = 1–3 μm [6].

> Thus, we have studied the CZTS films deposited by spray pyrolysis technique at different sprayed initial precursor volumes which had the higher thickness than studied in Ref. [9].

> In **Figure 2e**–**h**, the SEM images of CZTS films and its cross-section deposited at different *Vs* are presented. It can be seen that in the range of studied values continuous films were formed,

**Figure 2.** SEM images of ZnO films surface deposited at different temperature *Ts* , K: 473 (а), 573 (b), 623 (c), and 673 (d). Inset (d) shows the film surface with high resolution obtained at 673 K [19] and CZTS films surface deposited at different dispersed initial precursor volumes *Vs* , ml: 2 (e), 3 (f), 4 (g), and 5 (h). Inset (h) shows the films' cross-section [20].

which had a well adhesion to substrate and characterized by the absence of cracks and holes on their surfaces. The maximal layer thickness is determined by the cross-section method and it was *d* = 1.3 μm at *Vs* = 5 ml (**Figure 2h**).

crystallographic plane. There are also presented lines at angles 47.15–47.50° and 55.55–56.45° which correspond to the reflection from (220) and (312) CZTS planes, respectively. It should be noted that during the increasing of precursor volume the intensity of peaks is increased and its half-width is decreased. It may be caused by the increasing of film thickness and by improvement of the films' crystalline quality. It is well-known that intensities ratio between the number of diffraction reflections from kesterite and stannite crystallographic planes is different [30]. Taking into account this fact, determination of these ratios gives an opportunity to estimate precisely the materials dominate phase. The measured intensity ratio *І*(112)/*І*(220) from (112) and (220) crystallographic planes films was 2.23–2.56. These values are similar to the values, determined for un-doped films with kesterite phase (*І*(112)/*І*(220) ≈ 2.80) [31]; thus, most probably, the investigated films have a kesterite phase. This conclusion is confirmed by the experimental measurements of lattice constants ratio (*с*/*2а* = 0.9970–1.0203), that was similar

Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method

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75

Lattice parameter of the materials is a characteristic which is very sensitive to stoichiometry varying, impurities introduction, oxidation, etc. Thus, the precise determination of these val-

In **Figure 4**, the dependencies of ZnO and CZTS films lattice parameters *а*, *с* vs. deposition conditions are presented. In **Figure 4a**, it can be seen that during the increase of substrate temperature, measured parameters *а*, *с* for ZnO films are approached to the reference data that may be caused by the film stoichiometry improvement. High-temperature condensates composition approaching to the stoichiometric is confirmed by the chemical composition analysis data. In case of CZTS films (**Figure 4b**) most similar to the reference data *а* and *с* values are obtained by us at *Vs* = 4–5 ml, that is well-correlated with elemental composition analysis. It was estimated that lattice parameters are varied in the range of *aZnO* = 0.32477–0.32554 nm, *сZnO* = 0.51507– 0.52111 nm, *c/aZnO* = 1.5822–1.6046, *аCZTS* = 0.5423–0.5480 nm, *сCZTS* = 1.0823–1.1182 nm, and *с*/*2аCZTS* = 0.9970–1.0203, the unit cell volume was in the range of *Vcell(ZnO)* = 0.0427–0.0477 nm<sup>3</sup>

obtained for ZnO and CZTS films deposited by spray pyrolysis technique in Ref. [34].

**Figure 4.** The dependencies of the lattice parameters *а*, *с* on the physical and technological deposition conditions

, K—for ZnO films (a) dispersion solution volume *Vs*

, that is well-correlated to the reference data [31] and values

, ml—for CZTS films (b)]. Horizontal lines

to 1.0. These values are typical for kesterite [32].

ues allows us to study the corresponding processes.

and *Vcell(CZTS)* = 0.3183–0.3358 nm<sup>3</sup>

[substrate temperature *Ts*

are corresponding to the stoichiometric material.

#### **3.2. Structural and substructural properties**

Structural and substructural properties of ZnO and CZTS films have a significant influence on functional characteristics of devices [18–20]. Thus, its study is an important scientific objective. For example, band gap of zinc oxide films can be significantly increased by means of nanostructuring due to the quantum-size effects. At the same time, CZTS films as the absorber layers is SC should have the crystallites with sizes larger that diffusion length of minority charge carriers [6]. However, the films obtained by spray pyrolysis are usually characterized by high levels of microdeformations, microstresses, and density of dislocations in comparison to the values observed in the condensates deposited by physical vacuum methods, e.g., thermal evaporation, magnetron sputtering, etc.

The detailed description of methods applied to study structural, substructural, and optical properties of films is described elsewhere [18–20].

In **Figure 3a**, the XRD patterns of ZnO films deposited at different substrate temperatures are presented. On the diffraction patterns of the low-temperature films is dominated the diffraction line at angles 35.60–36.10° that corresponds to the reflection from (101) plane of ZnO hexagonal phase. On the diffraction patterns of the films deposited at *Ts* > 573 K, the lines at angles 31.80° and 34.80° are dominated, which correspond to the reflections of (100) and (002) crystallographic planes, respectively. X-ray analysis has been shown that deposited samples are single-phase and contain ZnO hexagonal phase. Secondary phases have not been determined by XRD analysis.

In **Figure 3b**, the X-ray patterns of CZTS films deposited at different dispersed solution volumes are shown. As can be seen from **Figure 3b**, on X-ray patterns is dominated the line on angles 28.05–28.50° which corresponds to the reflection from (112) CZTS tetragonal phase

**Figure 3.** XRD patterns of ZnO films obtained at different substrate temperatures *Ts* , K: 473 (1), 523 (2), 573 (3), 623 (4), and 673 (5) (а); and CZTS films deposited at different dispersion solution volume *Vs* , ml: 2 (1), 3 (2), 4 (3), and 5 (4) (b). Vertical lines correspond to the JCPDS cards (ZnO—№ 01–089-1397; CZTS—№ 00–026-0575) [19, 20].

crystallographic plane. There are also presented lines at angles 47.15–47.50° and 55.55–56.45° which correspond to the reflection from (220) and (312) CZTS planes, respectively. It should be noted that during the increasing of precursor volume the intensity of peaks is increased and its half-width is decreased. It may be caused by the increasing of film thickness and by improvement of the films' crystalline quality. It is well-known that intensities ratio between the number of diffraction reflections from kesterite and stannite crystallographic planes is different [30]. Taking into account this fact, determination of these ratios gives an opportunity to estimate precisely the materials dominate phase. The measured intensity ratio *І*(112)/*І*(220) from (112) and (220) crystallographic planes films was 2.23–2.56. These values are similar to the values, determined for un-doped films with kesterite phase (*І*(112)/*І*(220) ≈ 2.80) [31]; thus, most probably, the investigated films have a kesterite phase. This conclusion is confirmed by the experimental measurements of lattice constants ratio (*с*/*2а* = 0.9970–1.0203), that was similar to 1.0. These values are typical for kesterite [32].

Lattice parameter of the materials is a characteristic which is very sensitive to stoichiometry varying, impurities introduction, oxidation, etc. Thus, the precise determination of these values allows us to study the corresponding processes.

In **Figure 4**, the dependencies of ZnO and CZTS films lattice parameters *а*, *с* vs. deposition conditions are presented. In **Figure 4a**, it can be seen that during the increase of substrate temperature, measured parameters *а*, *с* for ZnO films are approached to the reference data that may be caused by the film stoichiometry improvement. High-temperature condensates composition approaching to the stoichiometric is confirmed by the chemical composition analysis data. In case of CZTS films (**Figure 4b**) most similar to the reference data *а* and *с* values are obtained by us at *Vs* = 4–5 ml, that is well-correlated with elemental composition analysis. It was estimated that lattice parameters are varied in the range of *aZnO* = 0.32477–0.32554 nm, *сZnO* = 0.51507– 0.52111 nm, *c/aZnO* = 1.5822–1.6046, *аCZTS* = 0.5423–0.5480 nm, *сCZTS* = 1.0823–1.1182 nm, and *с*/*2аCZTS* = 0.9970–1.0203, the unit cell volume was in the range of *Vcell(ZnO)* = 0.0427–0.0477 nm<sup>3</sup> and *Vcell(CZTS)* = 0.3183–0.3358 nm<sup>3</sup> , that is well-correlated to the reference data [31] and values obtained for ZnO and CZTS films deposited by spray pyrolysis technique in Ref. [34].

**Figure 4.** The dependencies of the lattice parameters *а*, *с* on the physical and technological deposition conditions [substrate temperature *Ts* , K—for ZnO films (a) dispersion solution volume *Vs* , ml—for CZTS films (b)]. Horizontal lines are corresponding to the stoichiometric material.

**Figure 3.** XRD patterns of ZnO films obtained at different substrate temperatures *Ts*

and 673 (5) (а); and CZTS films deposited at different dispersion solution volume *Vs*

Vertical lines correspond to the JCPDS cards (ZnO—№ 01–089-1397; CZTS—№ 00–026-0575) [19, 20].

which had a well adhesion to substrate and characterized by the absence of cracks and holes on their surfaces. The maximal layer thickness is determined by the cross-section method and

Structural and substructural properties of ZnO and CZTS films have a significant influence on functional characteristics of devices [18–20]. Thus, its study is an important scientific objective. For example, band gap of zinc oxide films can be significantly increased by means of nanostructuring due to the quantum-size effects. At the same time, CZTS films as the absorber layers is SC should have the crystallites with sizes larger that diffusion length of minority charge carriers [6]. However, the films obtained by spray pyrolysis are usually characterized by high levels of microdeformations, microstresses, and density of dislocations in comparison to the values observed in the condensates deposited by physical vacuum methods, e.g., ther-

The detailed description of methods applied to study structural, substructural, and optical

In **Figure 3a**, the XRD patterns of ZnO films deposited at different substrate temperatures are presented. On the diffraction patterns of the low-temperature films is dominated the diffraction line at angles 35.60–36.10° that corresponds to the reflection from (101) plane of ZnO hexagonal phase. On the diffraction patterns of the films deposited at *Ts* > 573 K, the lines at angles 31.80° and 34.80° are dominated, which correspond to the reflections of (100) and (002) crystallographic planes, respectively. X-ray analysis has been shown that deposited samples are single-phase and contain ZnO hexagonal phase. Secondary phases have not been deter-

In **Figure 3b**, the X-ray patterns of CZTS films deposited at different dispersed solution volumes are shown. As can be seen from **Figure 3b**, on X-ray patterns is dominated the line on angles 28.05–28.50° which corresponds to the reflection from (112) CZTS tetragonal phase

it was *d* = 1.3 μm at *Vs* = 5 ml (**Figure 2h**).

74 Semiconductors - Growth and Characterization

**3.2. Structural and substructural properties**

mal evaporation, magnetron sputtering, etc.

mined by XRD analysis.

properties of films is described elsewhere [18–20].

, K: 473 (1), 523 (2), 573 (3), 623 (4),

, ml: 2 (1), 3 (2), 4 (3), and 5 (4) (b).

In **Figure 5**, the results of measurements *L* and *ε* parameters in studied films by threefold convolution method are presented. As can be seen from **Figure 5a**, in ZnO films during the increasing of substrate temperature from 473 to 673 K, there is a tendency of the CDS values increasing in direction of [100]—from *L* ~ 14 to ~ 21 nm, in direction of [101]—from *L* ~ 11 to ~ 20 nm and in direction of [102]—from *L* ~ 10 to ~ 63 nm. Similar *L*-*Ts* dependencies are observed in our previous works [33], where II-VI type compounds (CdTe, ZnS, ZnSe, and ZnTe) were obtained by close-spaced vacuum sublimation technique. At the same time, the microdeformations level in ZnO films in direction of [100] is decreased from *ε* ~ 1.6 × 10−5 to ~ 0.5 × 10−5, in direction of [101]—from *ε* ~ 3.5 × 10−5 to ~ 1.2 × 10−5, in direction of [102]—from *ε* ~ 1.0 × 10−5 to ~ 0.7 × 10−5, at the same substrate temperatures range (**Figure 5b**). Similar *ε* decreasing at substrate temperatures of *Ts* > 573 K is observed in CdTe and ZnTe films [35].

same time, the microdeformations level in CZTS films for the directions normal to (112)–(220) crystallographic planes is varied in the range of *ε* ~ (0.93–0.99) × 10−3; for (112)–(312) planes—*ε* ~ (0.76–0.77) × 10−3; for (220)–(312) planes—*ε* ~ (0.65–0.71) × 10−3 (**Figure 5d**). It should be noted that measured microdeformations values in CZTS films are lower than presented in Ref. [37], where *ε* ~ (1.26–6.60) × 10−3. The measured level of microdeformations allowed us to determine the level of microstresses (*σ*) in nanocrystalline ZnO and CZTS films. It was estimated that microstresses levels in ZnO and CZTS films varied in the ranges of *σZnO* = 0.48–1.53 MPa,

Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method

In CZTS films, during the increase of the dispersion solution volume microstresses level is

In **Figure 6**, the results of measurements of the dislocations concentration on the boundaries

in direction normal to (100) crystallographic plane for ZnO films or (112) crystallographic plane for CZTS films are presented. Studied ZnO layers are characterized by rather low

ues. In Ref. [38], authors have estimated that in ZnO nanocrystalline films with thickness of *d* = 0.135–0.392 μm, deposited at *Ts* = 473 K, the dislocation concentration values are higher

dispersion volume of *Vs* = 5 ml. It should be noted that these values *ρ* are smaller than observed earlier for CZTS films deposited by chemical methods (spray pyrolysis – *ρ* = (11.6–80.3) ×

[40]).

 (in case of CZTS films (b)) on dislocations density *ρ*: on the sub-grain boundaries (1), within CDS units (2) and general dislocations concentration (3) for the direction normal to (100)–(200) planes for ZnO and to (112)–(220) planes for

) [35]), and higher compare to the films obtained by vacuum methods (thermal

) CSDs blocks and of the general dislocations concentration (*ρ*)

) than determined here. In Ref. [39], authors have also obtained

(**Figure 6b**) in CZTS films general dislocation density *ρ* does not

compare to the results obtained by other authors. As it

) compared to our results. It has been shown that

(in case of ZnO films (a)) and dispersed initial solution volume

decreased, wherein the smallest *σ* values are observed in layers obtained at *Vs* = 5 ml.

onto *σ* level in ZnO films was also stud-

http://dx.doi.org/10.5772/intechopen.72988

, there is a tendency to decrease *ρ* val-

are obtained in case of film deposited at

from 623 to 723 K the

77

*σCZTS* = 5.2–20.3 MPa, respectively. The influence of *Ts*

can be seen from **Figure 6a**, during the increase of *Ts*

(*ρL*

) and within volume (*ρε*

(*ρ* = (1.29–4.15) × 10<sup>15</sup> lines/m<sup>2</sup>

during the increase of *Vs*

1016 lines/m<sup>2</sup>

*Vs*

values of *ρ* = (1.3–6.1) × 10<sup>13</sup> lines/m<sup>2</sup>

higher values (*ρ* = (2.4–5.8) × 10<sup>13</sup> lines/m<sup>2</sup>

almost change in all investigated directions.

evaporation – *ρ* = (0.64–4.00) × 10<sup>14</sup> lines/m<sup>2</sup>

**Figure 6.** The influence of substrate temperature *Ts*

CZTS. The measurements error was varied in the range of 15–20%.

The smallest values of *ρ* = (17.0–19.3) × 10<sup>15</sup> lines/m<sup>2</sup>

compression stress level *σ* was decreased (1.77–1.47 GPa).

ied in Refs. [8], where authors estimated that during the increase of *Ts*

As can be seen from **Figure 5c**, in CZTS films during the increasing of the dispersion solution volume from 2 to 5 ml, CDS values are almost were not changed: *L* ~ 24–26 nm (for (112)–(220) planes pair), *L* ~ 25–27 nm (for(112)–(312) planes pair), and *L* ~ 39–40 nm (for (220)–(312) planes pair). Consequently, *Vs* varying influence on CDS sizes is negligible. It should be noted that obtained results of *L* measurements are well-correlated to the results presented in Refs. [36], where CZTS films were deposited under the similar experimental conditions. At the

**Figure 5.** Influence of the substrate temperature on CDS values (а) and microdeformations level (b) of ZnO films on direction normal to the (100)–(200) (1), (101)–(202) (2), (102)–(103) (3) crystallographic planes and of the dispersion solution volume *Vs* on *L* (c) and *ε* (d) for CZTS films on direction normal to the (112)–(220) (1), (112)–(312) (2), (220)–(312) (3) planes. The threefold convolution technique was used.

same time, the microdeformations level in CZTS films for the directions normal to (112)–(220) crystallographic planes is varied in the range of *ε* ~ (0.93–0.99) × 10−3; for (112)–(312) planes—*ε* ~ (0.76–0.77) × 10−3; for (220)–(312) planes—*ε* ~ (0.65–0.71) × 10−3 (**Figure 5d**). It should be noted that measured microdeformations values in CZTS films are lower than presented in Ref. [37], where *ε* ~ (1.26–6.60) × 10−3. The measured level of microdeformations allowed us to determine the level of microstresses (*σ*) in nanocrystalline ZnO and CZTS films. It was estimated that microstresses levels in ZnO and CZTS films varied in the ranges of *σZnO* = 0.48–1.53 MPa, *σCZTS* = 5.2–20.3 MPa, respectively. The influence of *Ts* onto *σ* level in ZnO films was also studied in Refs. [8], where authors estimated that during the increase of *Ts* from 623 to 723 K the compression stress level *σ* was decreased (1.77–1.47 GPa).

In **Figure 5**, the results of measurements *L* and *ε* parameters in studied films by threefold convolution method are presented. As can be seen from **Figure 5a**, in ZnO films during the increasing of substrate temperature from 473 to 673 K, there is a tendency of the CDS values increasing in direction of [100]—from *L* ~ 14 to ~ 21 nm, in direction of [101]—from

cies are observed in our previous works [33], where II-VI type compounds (CdTe, ZnS, ZnSe, and ZnTe) were obtained by close-spaced vacuum sublimation technique. At the same time, the microdeformations level in ZnO films in direction of [100] is decreased from *ε* ~ 1.6 × 10−5 to ~ 0.5 × 10−5, in direction of [101]—from *ε* ~ 3.5 × 10−5 to ~ 1.2 × 10−5, in direction of [102]—from *ε* ~ 1.0 × 10−5 to ~ 0.7 × 10−5, at the same substrate temperatures range (**Figure 5b**). Similar *ε* decreasing at substrate temperatures of *Ts* > 573 K is observed

As can be seen from **Figure 5c**, in CZTS films during the increasing of the dispersion solution volume from 2 to 5 ml, CDS values are almost were not changed: *L* ~ 24–26 nm (for (112)–(220) planes pair), *L* ~ 25–27 nm (for(112)–(312) planes pair), and *L* ~ 39–40 nm (for (220)–(312)

that obtained results of *L* measurements are well-correlated to the results presented in Refs. [36], where CZTS films were deposited under the similar experimental conditions. At the

**Figure 5.** Influence of the substrate temperature on CDS values (а) and microdeformations level (b) of ZnO films on direction normal to the (100)–(200) (1), (101)–(202) (2), (102)–(103) (3) crystallographic planes and of the dispersion

on *L* (c) and *ε* (d) for CZTS films on direction normal to the (112)–(220) (1), (112)–(312) (2), (220)–(312)

varying influence on CDS sizes is negligible. It should be noted

dependen-

*L* ~ 11 to ~ 20 nm and in direction of [102]—from *L* ~ 10 to ~ 63 nm. Similar *L*-*Ts*

in CdTe and ZnTe films [35].

76 Semiconductors - Growth and Characterization

planes pair). Consequently, *Vs*

solution volume *Vs*

(3) planes. The threefold convolution technique was used.

In CZTS films, during the increase of the dispersion solution volume microstresses level is decreased, wherein the smallest *σ* values are observed in layers obtained at *Vs* = 5 ml.

In **Figure 6**, the results of measurements of the dislocations concentration on the boundaries (*ρL* ) and within volume (*ρε* ) CSDs blocks and of the general dislocations concentration (*ρ*) in direction normal to (100) crystallographic plane for ZnO films or (112) crystallographic plane for CZTS films are presented. Studied ZnO layers are characterized by rather low values of *ρ* = (1.3–6.1) × 10<sup>13</sup> lines/m<sup>2</sup> compare to the results obtained by other authors. As it can be seen from **Figure 6a**, during the increase of *Ts* , there is a tendency to decrease *ρ* values. In Ref. [38], authors have estimated that in ZnO nanocrystalline films with thickness of *d* = 0.135–0.392 μm, deposited at *Ts* = 473 K, the dislocation concentration values are higher (*ρ* = (1.29–4.15) × 10<sup>15</sup> lines/m<sup>2</sup> ) than determined here. In Ref. [39], authors have also obtained higher values (*ρ* = (2.4–5.8) × 10<sup>13</sup> lines/m<sup>2</sup> ) compared to our results. It has been shown that during the increase of *Vs* (**Figure 6b**) in CZTS films general dislocation density *ρ* does not almost change in all investigated directions.

The smallest values of *ρ* = (17.0–19.3) × 10<sup>15</sup> lines/m<sup>2</sup> are obtained in case of film deposited at dispersion volume of *Vs* = 5 ml. It should be noted that these values *ρ* are smaller than observed earlier for CZTS films deposited by chemical methods (spray pyrolysis – *ρ* = (11.6–80.3) × 1016 lines/m<sup>2</sup> ) [35]), and higher compare to the films obtained by vacuum methods (thermal evaporation – *ρ* = (0.64–4.00) × 10<sup>14</sup> lines/m<sup>2</sup> [40]).

**Figure 6.** The influence of substrate temperature *Ts* (in case of ZnO films (a)) and dispersed initial solution volume *Vs* (in case of CZTS films (b)) on dislocations density *ρ*: on the sub-grain boundaries (1), within CDS units (2) and general dislocations concentration (3) for the direction normal to (100)–(200) planes for ZnO and to (112)–(220) planes for CZTS. The measurements error was varied in the range of 15–20%.

#### **3.3. The study of the stoichiometry**

Energy dispersed analysis of the X-ray spectra (EDAX) allows us to determine the elemental composition of ZnO and CZTS films obtained in present work. Results determined for films deposited at different physical-chemical and technological conditions are presented in **Table 2**. As it can be seen, ZnO films have some oxygen surplus on zinc. Besides, films stoichiometry is increased during the increasing of the substrate temperature. This fact is confirmed by the concentration ratios *CO/CZn*, that are parts of compound (*γZnO* = 1.4 – *Ts* = 473 K, *γZnO* = 1.2 – *Ts* = 623 K). The impurities connected to the films contamination by the precursor materials have not been determined.

**4. Optical properties of ZnO and CZTS films obtained by spray** 

The study and control of the optical properties of ZnO, CZTS, and CZT films is an important task with the aim of their usage in optoelectronic devices, especially for SCs development. It is well-known that optical characteristics of these films heavily dependent on morphological, structural, substructural properties, chemical composition, and physical (chemical) and tech-

In present work, the transmission light coefficient of ZnO films was in the range of *T* = 60–80% at the wavelength range of *λ* = 430–800 nm. The highest transmission values had films obtained

caused by increasing of the grain sizes in films and by improvement of their structural quality dur-

in bulk materials. During the increasing of the grain size, quantum effects are gradually decreased. At the same time, due to the high level of the substructural defects (primarily dislocations) in nanocrystalline films, that have been given the local deformations on the materials lattice, its average *Eg* have been smaller than in bulk materials [16]. At high substrate temperatures, films with sufficient large grain size and low structural defect concentration were formed. As a result, the band gap of

is presented. It should be noted that the smallest *α* values have been obtained for layers

in CZTS films (b). Dashed line corresponds to band gap value in bulk ZnO (*Eg* = 3.37 eV) and bulk CZTS (*Eg* = 1.50 eV).

. It is well known [43] that in nanocrystalline films (*DC* < 100 nm) band gap is

Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method

*.*

temperature is at first increased and in further decreased. This complex dependence of *Eg*

values for ZnO films were in the range of

http://dx.doi.org/10.5772/intechopen.72988

compare to the values observed

changing depending on the

may be

79

of zinc oxide during the increasing of the deposition

on the dispersed solution precursor volume

in ZnO films (а) and on dispersion solution volume *Vs*

**pyrolysis technique**

**4.1. Optical properties**

ing the increasing of *Ts*

*Vs*

**Figure 7.** Bang gap (*Eg*

nological deposition conditions.

at *Ts* = 673 K. It was estimated that measurement *E*<sup>g</sup>

determined by quantum effects, that leads to the increasing of *Eg*

semiconductor is approaching the bulk value. Similar tendencies of *Eg*

) dependencies on substrate temperature *Ts*

deposition temperature were observed in Refs. [44].

In **Figure 7b**, dependence of the materials *Eg*

3.18–3.30 eV and were also dependent on *Ts*

As can be seen from **Figure 7a**, band gap *Eg*

The control of CZTS films elemental composition is a complex and important task because of its probable determination of the phase conditions, crystal structure, optical, and electrical properties of investigated layers. It was estimated that in CZTS films some copper, zinc, and tin are present in surplus and has some sulfur deficiency. Sulfur losses in films during the pyrolytic reaction of the initial precursor near the surface of the heated substrate may be caused by its high volatility [41]. It should be noted that stoichiometry of studied films is some improved during the increasing of dispersed precursor volume. Also, the obtained ratio *γCZTS\_1* = (0.80–0.84) in CZTS films deposited at precursor dispersion with volume of *Vs* = 2–3 ml is close to the optimal values necessary to develop SCs with high solar energy conversion efficiency (*γCZTS\_1* = (0.8–0.9), *γCZTS\_2* = (1.1–1.2)) [40, 42]. For film obtained by dispersion precursor volume of 3 ml for this requirement corresponds the next ratio—*γCZTS\_2* = 1.2. Impurities related to the films' contamination by the precursor's materials have also not been observed in CZTS layers.


**Table 2.** Measurement results of the chemical composition for ZnO and CZTS films obtained at different conditions.

## **4. Optical properties of ZnO and CZTS films obtained by spray pyrolysis technique**

#### **4.1. Optical properties**

**3.3. The study of the stoichiometry**

78 Semiconductors - Growth and Characterization

observed in CZTS layers.

 **(K)** *CZn* **(at. %)** *CO* **(at. %)** *γZnO* 41.8 58.2 1.4 42.3 57.7 1.4 42.6 57.4 1.3 44.3 55.7 1.2 44.0 56.0 1.2 Stoichiometry 50.0 50.0 1.0

(ml) *CCu* (at. %) *CZn* (at. %) *CSn* (at. %) *CS*

 28.6 21.4 14.3 35.8 0.8 1.5 0.8 27.0 17.3 14.7 40.8 0.8 1.2 0.7 27.7 16.3 15.1 40.9 0.9 1.0 0.6 26.4 15.2 15.4 43.0 0.9 1.0 0.6 Stoichiometry 25.0 12.5 12.5 50.0 1.0 1.0 0.5

**Table 2.** Measurement results of the chemical composition for ZnO and CZTS films obtained at different conditions.

(at. %) *γCZTS\_1 γCZTS\_2 γCZTS\_3*

**ZnO** *Ts*

CZTS *Vs*

Energy dispersed analysis of the X-ray spectra (EDAX) allows us to determine the elemental composition of ZnO and CZTS films obtained in present work. Results determined for films deposited at different physical-chemical and technological conditions are presented in **Table 2**. As it can be seen, ZnO films have some oxygen surplus on zinc. Besides, films stoichiometry is increased during the increasing of the substrate temperature. This fact is confirmed by the concentration ratios *CO/CZn*, that are parts of compound (*γZnO* = 1.4 – *Ts* = 473 K, *γZnO* = 1.2 – *Ts* = 623 K). The impurities connected to the films contamination by the precursor materials have not been determined. The control of CZTS films elemental composition is a complex and important task because of its probable determination of the phase conditions, crystal structure, optical, and electrical properties of investigated layers. It was estimated that in CZTS films some copper, zinc, and tin are present in surplus and has some sulfur deficiency. Sulfur losses in films during the pyrolytic reaction of the initial precursor near the surface of the heated substrate may be caused by its high volatility [41]. It should be noted that stoichiometry of studied films is some improved during the increasing of dispersed precursor volume. Also, the obtained ratio *γCZTS\_1* = (0.80–0.84) in CZTS films deposited at precursor dispersion with volume of *Vs* = 2–3 ml is close to the optimal values necessary to develop SCs with high solar energy conversion efficiency (*γCZTS\_1* = (0.8–0.9), *γCZTS\_2* = (1.1–1.2)) [40, 42]. For film obtained by dispersion precursor volume of 3 ml for this requirement corresponds the next ratio—*γCZTS\_2* = 1.2. Impurities related to the films' contamination by the precursor's materials have also not been

The study and control of the optical properties of ZnO, CZTS, and CZT films is an important task with the aim of their usage in optoelectronic devices, especially for SCs development. It is well-known that optical characteristics of these films heavily dependent on morphological, structural, substructural properties, chemical composition, and physical (chemical) and technological deposition conditions.

In present work, the transmission light coefficient of ZnO films was in the range of *T* = 60–80% at the wavelength range of *λ* = 430–800 nm. The highest transmission values had films obtained at *Ts* = 673 K. It was estimated that measurement *E*<sup>g</sup> values for ZnO films were in the range of 3.18–3.30 eV and were also dependent on *Ts .*

As can be seen from **Figure 7a**, band gap *Eg* of zinc oxide during the increasing of the deposition temperature is at first increased and in further decreased. This complex dependence of *Eg* may be caused by increasing of the grain sizes in films and by improvement of their structural quality during the increasing of *Ts* . It is well known [43] that in nanocrystalline films (*DC* < 100 nm) band gap is determined by quantum effects, that leads to the increasing of *Eg* compare to the values observed in bulk materials. During the increasing of the grain size, quantum effects are gradually decreased. At the same time, due to the high level of the substructural defects (primarily dislocations) in nanocrystalline films, that have been given the local deformations on the materials lattice, its average *Eg* have been smaller than in bulk materials [16]. At high substrate temperatures, films with sufficient large grain size and low structural defect concentration were formed. As a result, the band gap of semiconductor is approaching the bulk value. Similar tendencies of *Eg* changing depending on the deposition temperature were observed in Refs. [44].

In **Figure 7b**, dependence of the materials *Eg* on the dispersed solution precursor volume *Vs* is presented. It should be noted that the smallest *α* values have been obtained for layers

**Figure 7.** Bang gap (*Eg* ) dependencies on substrate temperature *Ts* in ZnO films (а) and on dispersion solution volume *Vs* in CZTS films (b). Dashed line corresponds to band gap value in bulk ZnO (*Eg* = 3.37 eV) and bulk CZTS (*Eg* = 1.50 eV).

deposited at volume of *Vs* = 2 ml, the highest values—at *Vs* = 5 ml, respectively. It is quite typical because of the smallest and highest thickness values of the corresponding layers. During the increasing of dispersed initial precursor solution volume, the values of band gap were varied in the range of *Eg* = 1.06–1.30 eV and were approximately approached to the values typical for bulk stoichiometric material (*Eg* = 1.5 eV). It indicates on increasing of grain sizes and decreasing of films deficiency during the increasing of their thickness. Similar tendencies have been observed in Ref. [9].

where the light absorption and transmission in films are performed, allows us to determine the functional links between chemical elements which are part of the studied materials.

Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method

http://dx.doi.org/10.5772/intechopen.72988

81

In **Figure 8b**, FTIR reflection spectra of ZnO films deposited at different substrate temperatures are presented. Although that thin films were deposited in air by chemical technique

At low frequencies (460–475 cm−1), there is observed minima, which due to the reference data [48], correspond to Zn-O vibrational mode. It should be noted that FTIR spectra obtained on films deposited in all range of substrate temperatures have a C-Cl vibrational mode [50]. The presence of this connection may be caused by the usage of HCl acid, which was added as a precursor during its preparation. The acid paths are also observed in films. In FTIR spectra of ZnO films deposited at *Ts* < 573 K, peaks on the frequencies 1405 and 1560 cm−1 are presented; they were interpreted by us as symmetric and asymmetric С-О vibrational modes [50]. The absence of C-O connections in films deposited at *Ts* > 573 K indicates about the total precursor decomposition near the substrate surface at these temperatures. It eliminates the possibility of adsorption of the acetate elements on ZnO films surface during the pyrolysis, and it leads

obtained spectra were comparatively pure.

, ZnO, and Znx

2 3 4 5

Snx Sy , Cux SnS<sup>y</sup>

*Vs*

Raman shift (cm−1)

Green-laser (*λ* = 514.5 nm)

Red laser (*λ* = 632.8 nm)

UV-laser (*λ* = 325 nm)

to the formation of single-phase zinc oxide polycrystalline films.

SnO<sup>y</sup>

**Experimental data Literature data**

**Table 3.** Peaks interpretation presented on Raman spectra of CZTS films.

It is well known that in CZTS films, the presence of secondary phases, such as Cux

lattices, and they indicate on XRD patterns refractions on similar angles. It complicates the phase analysis by XRD technique. Thus, for precise identification of the secondary phases in CZTS compound, the researchers often use Raman spectroscopy in addition to XRD analysis [54]. It allows to identify not only secondary phases, but also kesterite and stannite. In **Table 3**,

, ml Raman shift, cm−1 Symmetry Mode Reference

142 143–144 E *CZTS E* [56] 340 338–339 A *CZTS A* [54] 664 672 A *2a CZTS A* [56]

339 338–339 A *CZTS A* [52] 663 672 A *2a CZTS A* [57]

340 341 A *CZTS A* [56] — 560 — — 541 — *ZnO* [45] 664 672 A *2a CZTS A* [57]

, is available [39, 51–53]. They are characterized by affiliated

Sy , Znx Sy ,

#### **4.2. Raman and Fourier transform IR (FTIR) spectra**

Raman spectroscopy is an additional to X-ray diffraction analysis method of studying the phase composition and quality of ZnO, CZTS, and CZT thin films.

Raman spectra of ZnO films measured in the range of frequencies 90–800 cm−1 are presented in **Figure 8a**. In spectra, a number of different intensity lines on the next frequencies: 95–98, 333–336, 415, 439–442, 572, and 578–587 cm−1 are observed. Using the reference data, these lines were interpreted by us as the next phonon modes: *E2 low*(*Zn*) [43–45], *E2 high-E2 low* [46], *E1* (*TO*) [45], *E2 high*(*O*) [43–47], *A1* (*LO*) [43] and *E1* (*LO*) [45–46]. In **Figure 8a**, two intensive peaks, which correspond to *E2* mode, are also observed: peak *E2 high*, which is relative to the oxygen anions, is localized at frequency of 439–442 cm−1 and peak *E2 low*, which is relative to zinc cations, is localized at frequency of 95–98 cm−1. It is well known [49] that the crystalline quality of ZnO films has a direct influence on the mode *E2* intensity. Besides, *E2 high*(*О*) peak is very sensitive to the presence of inner defects of material. The deviation of the frequency *E2 high*(*О*) peak from the value typical for bulk ZnO (437 cm−1), that is observed by us in low-temperature condensates, indicates about the presence in zinc oxide high level of microstresses and stretched defects (dislocations) density of the lattice. It should be noted that during the increase of the substrate temperature, the *E2 high*(*O*) peak position is some red-shifted from 442 cm−1 to the typical bulk ZnO values—439 cm−1, which indicates the decrease of *σ* and *ρ* levels.

FTIR spectroscopy is an addition to X-ray diffraction analysis and Raman spectroscopy technique, which allows to obtain an information about the elemental composition of the studied material and its contamination by the precursor impurities. The number of frequencies,

**Figure 8.** Raman (a) and FTIR (b) spectra of ZnO films deposited at different substrate temperatures *Ts* , *K*: 473 (1), 523 (2), 573 (3), 623 (4), and 673 (5).

where the light absorption and transmission in films are performed, allows us to determine the functional links between chemical elements which are part of the studied materials.

deposited at volume of *Vs* = 2 ml, the highest values—at *Vs* = 5 ml, respectively. It is quite typical because of the smallest and highest thickness values of the corresponding layers. During the increasing of dispersed initial precursor solution volume, the values of band gap were varied in the range of *Eg* = 1.06–1.30 eV and were approximately approached to the values typical for bulk stoichiometric material (*Eg* = 1.5 eV). It indicates on increasing of grain sizes and decreasing of films deficiency during the increasing of their thickness. Similar tendencies

Raman spectroscopy is an additional to X-ray diffraction analysis method of studying the

Raman spectra of ZnO films measured in the range of frequencies 90–800 cm−1 are presented in **Figure 8a**. In spectra, a number of different intensity lines on the next frequencies: 95–98, 333–336, 415, 439–442, 572, and 578–587 cm−1 are observed. Using the reference data, these

ized at frequency of 95–98 cm−1. It is well known [49] that the crystalline quality of ZnO films

value typical for bulk ZnO (437 cm−1), that is observed by us in low-temperature condensates, indicates about the presence in zinc oxide high level of microstresses and stretched defects (dislocations) density of the lattice. It should be noted that during the increase of the substrate

FTIR spectroscopy is an addition to X-ray diffraction analysis and Raman spectroscopy technique, which allows to obtain an information about the elemental composition of the studied material and its contamination by the precursor impurities. The number of frequencies,

**Figure 8.** Raman (a) and FTIR (b) spectra of ZnO films deposited at different substrate temperatures *Ts*

intensity. Besides, *E2*

*high*(*O*) peak position is some red-shifted from 442 cm−1 to the typical bulk

*low*(*Zn*) [43–45], *E2*

(*LO*) [45–46]. In **Figure 8a**, two intensive peaks, which

*high*, which is relative to the oxygen anions, is

*low*, which is relative to zinc cations, is local-

*high-E2*

*high*(*О*) peak is very sensitive to the

*low* [46], *E1*

*high*(*О*) peak from the

, *K*: 473 (1), 523

(*TO*)

have been observed in Ref. [9].

80 Semiconductors - Growth and Characterization

*high*(*O*) [43–47], *A1*

has a direct influence on the mode *E2*

[45], *E2*

correspond to *E2*

temperature, the *E2*

(2), 573 (3), 623 (4), and 673 (5).

**4.2. Raman and Fourier transform IR (FTIR) spectra**

lines were interpreted by us as the next phonon modes: *E2*

localized at frequency of 439–442 cm−1 and peak *E2*

(*LO*) [43] and *E1*

mode, are also observed: peak *E2*

presence of inner defects of material. The deviation of the frequency *E2*

ZnO values—439 cm−1, which indicates the decrease of *σ* and *ρ* levels.

phase composition and quality of ZnO, CZTS, and CZT thin films.

In **Figure 8b**, FTIR reflection spectra of ZnO films deposited at different substrate temperatures are presented. Although that thin films were deposited in air by chemical technique obtained spectra were comparatively pure.

At low frequencies (460–475 cm−1), there is observed minima, which due to the reference data [48], correspond to Zn-O vibrational mode. It should be noted that FTIR spectra obtained on films deposited in all range of substrate temperatures have a C-Cl vibrational mode [50]. The presence of this connection may be caused by the usage of HCl acid, which was added as a precursor during its preparation. The acid paths are also observed in films. In FTIR spectra of ZnO films deposited at *Ts* < 573 K, peaks on the frequencies 1405 and 1560 cm−1 are presented; they were interpreted by us as symmetric and asymmetric С-О vibrational modes [50]. The absence of C-O connections in films deposited at *Ts* > 573 K indicates about the total precursor decomposition near the substrate surface at these temperatures. It eliminates the possibility of adsorption of the acetate elements on ZnO films surface during the pyrolysis, and it leads to the formation of single-phase zinc oxide polycrystalline films.

It is well known that in CZTS films, the presence of secondary phases, such as Cux Sy , Znx Sy , Snx Sy , Cux SnS<sup>y</sup> , ZnO, and Znx SnO<sup>y</sup> , is available [39, 51–53]. They are characterized by affiliated lattices, and they indicate on XRD patterns refractions on similar angles. It complicates the phase analysis by XRD technique. Thus, for precise identification of the secondary phases in CZTS compound, the researchers often use Raman spectroscopy in addition to XRD analysis [54]. It allows to identify not only secondary phases, but also kesterite and stannite. In **Table 3**,


**Table 3.** Peaks interpretation presented on Raman spectra of CZTS films.

the results of study the Raman spectra of CZTS films using as an excitation source the radiation of several lasers are presented. At all spectra regardless on the precursor volume and excitation laser type, the main peak on frequencies of (339–340) cm−1 is presented. It is well correlated to the results of previous studies [52, 54, 55]. In Raman spectra obtained using the green laser, lines on the next frequencies: 142, 340, and 664 cm−1 are also observed, which correspond to *CZTS E*, *CZTS A*, and *2а CZTS A* (*CZTS A* mode phonon replica) phonon modes, respectively [54–56].

**5. Conclusions**

a kesterite structure.

weakly dependent on *Vs*

obtained at *Vs* = 3–4 ml.

sponded to A1

and also LO<sup>2</sup>

have a weak mode A<sup>1</sup>

quality.

depended on the dispersed solution volume.

It has been estimated that during the increase of *Ts*

.

sion coefficient *T* = 60–80%. Measured *E*<sup>g</sup>

CZT film spectra (х = 0.32) had a mode LO<sup>2</sup>

of 3.18–3.30 eV and had a complex dependence on *Ts*

namely decreasing *ε*, *σ*, and *ρ* values during the increase of *Ts*

absence of precursor impurities in ZnO films obtained at *Ts* > 573 K.

(Te), peaks of LO<sup>1</sup>

(ZnTe) mode resonant replica.

As a results of the complex study of structure, substructure, optical properties, and elemental composition of ZnO, CZTS, and CZT films obtained by pulsed spray pyrolysis technique dependent on the physical (chemical) and technological deposition conditions, it was determined that ZnO nanocrystalline films have an average grain size of *DC* = 25–270 nm and their thickness was *d* = 0.8–1.2 μm, and were formed at *Ts* > 473 K. CZTS continuous films with optimal thickness of *d* = 1.3 μm were deposited at dispersed initial precursor volume *Vs* = 5 ml. It was found that ZnO and CZTS films were polycrystalline in nature, single-phase, and had hexagonal and tetragonal phases, respectively. CZTS samples had

Nanostructured ZnO, Cu2ZnSnS4, Cd1−xZnxTe Thin Films Obtained by Spray Pyrolysis Method

It has been shown that in ZnO during the increasing of substrate temperature there is a tendency to the increasing of the CDS; however, in CZTS films, their CSD values were weakly

Lattice parameters values in ZnO and CZTS films deposited at *Ts* = 623 K, *Vs* = 4 ml were well-correlated to the reference data that confirms their optimal stoichiometry and crystalline

and dislocation density in ZnO films were decreased; in CZTS films, these parameters were

It has been determined that during the increasing of substrate temperature to 623 K stoichiometry of ZnO layers was improved (*γZnO* = 1.2). It has been shown that optimal for usage in SCs CZTS films, their stoichiometry ratios *γCZTS\_1* = 0.8–0.9, *γCZTS\_2* = 1.1–1.2, *γCZTS\_3* = 0.7 were

Study of the optical characteristics of ZnO films allow to estimate the high values of transmis-

of *Eg* = 1.06–1.30 eV of CZTS layers were approximately approached to the reference data *Eg* = 1.5 eV. Raman spectra analysis of ZnO films confirmed the results of the XRD study,

tra analysis has confirmed the single-phase nature of condensates. FTIR study indicated the

(CdTe), TO<sup>1</sup>

The results of a research study of the ZnO, CZTS, and CZT thin films will be used for the development of the devices, primarily, in third generation high-efficiency thin-film solar cells.

(Te) and ETO(Te) tellure modes were also determined. CZT film spectra (х = 0.75)

(CdTe), TO<sup>2</sup>

the microdeformations level, microstresses,

http://dx.doi.org/10.5772/intechopen.72988

83

values of ZnO layers were determined in the range

. During the increase of *Vs*

(ZnTe). In these spectra, intensive peaks corre-

(ZnTe), and LO<sup>2</sup>

, the values

(ZnTe) modes,

. CZTS films' Raman spec-

Usage of the red- and UV-lasers as phonons excitation source allows us to increase the method's sensitivity onto the revealing of compounds with optical band gap close to *Е<sup>g</sup>* ~ 1.96 and ~ 3.81 eV (excitation radiations energies of corresponding lasers). On spectra, obtained using the red- and UV-lasers, are presented lines on frequencies 339–340 cm−1, 663–664 cm−1 which correspond to the *CZTS A* and *2a CZTS A* phonon modes [52, 56, 57]. The usage of UV-laser in one of the studied films revealed a negligible number of ZnO secondary phase. In addition, these results are supported by the phonon excitation in Raman spectra on the frequency at 560 cm−1 on film obtained from the precursor volume dispersion of 3 ml. Other secondary phases in studied films are not revealed. Raman spectra of CZT films measured during the influence of green-laser excitation radiation (*λ* = 514 nm, *Е* = 2.41 eV) are presented in **Figure 9**. In the spectra of the CZT sample (*х* = 0.32), peak which corresponds to LO<sup>2</sup> (ZnTe) mode is observed. In these spectra, intensive peaks which correspond to A<sup>1</sup> (Te) ETO(Te) telluric modes are also detected. In the spectra of the CZT sample (*х* = 0.75), is observed a weak peak that corresponds to A<sup>1</sup> (Te) mode, peaks of the next modes: LO<sup>1</sup> (CdTe), TO<sup>1</sup> (CdTe), TO<sup>2</sup> (ZnTe), LO<sup>2</sup> (ZnTe), and also detected the LO<sup>2</sup> (ZnTe) mode resonant replica.

**Figure 9.** Raman spectra of CZT films measured during the impact of excitation irradiation of the wavelength 785 nm at room temperature (RT) [7].

## **5. Conclusions**

the results of study the Raman spectra of CZTS films using as an excitation source the radiation of several lasers are presented. At all spectra regardless on the precursor volume and excitation laser type, the main peak on frequencies of (339–340) cm−1 is presented. It is well correlated to the results of previous studies [52, 54, 55]. In Raman spectra obtained using the green laser, lines on the next frequencies: 142, 340, and 664 cm−1 are also observed, which correspond to *CZTS E*, *CZTS A*, and *2а CZTS A* (*CZTS A* mode phonon replica) phonon modes, respectively [54–56]. Usage of the red- and UV-lasers as phonons excitation source allows us to increase the meth-

~ 3.81 eV (excitation radiations energies of corresponding lasers). On spectra, obtained using the red- and UV-lasers, are presented lines on frequencies 339–340 cm−1, 663–664 cm−1 which correspond to the *CZTS A* and *2a CZTS A* phonon modes [52, 56, 57]. The usage of UV-laser in one of the studied films revealed a negligible number of ZnO secondary phase. In addition, these results are supported by the phonon excitation in Raman spectra on the frequency at 560 cm−1 on film obtained from the precursor volume dispersion of 3 ml. Other secondary phases in studied films are not revealed. Raman spectra of CZT films measured during the influence of green-laser excitation radiation (*λ* = 514 nm, *Е* = 2.41 eV) are presented in

~ 1.96 and

(ZnTe)

(ZnTe),

(Te) ETO(Te) telluric

(CdTe), TO<sup>2</sup>

(CdTe), TO<sup>1</sup>

od's sensitivity onto the revealing of compounds with optical band gap close to *Е<sup>g</sup>*

**Figure 9**. In the spectra of the CZT sample (*х* = 0.32), peak which corresponds to LO<sup>2</sup>

(Te) mode, peaks of the next modes: LO<sup>1</sup>

modes are also detected. In the spectra of the CZT sample (*х* = 0.75), is observed a weak peak

**Figure 9.** Raman spectra of CZT films measured during the impact of excitation irradiation of the wavelength 785 nm at

(ZnTe) mode resonant replica.

mode is observed. In these spectra, intensive peaks which correspond to A<sup>1</sup>

that corresponds to A<sup>1</sup>

room temperature (RT) [7].

(ZnTe), and also detected the LO<sup>2</sup>

82 Semiconductors - Growth and Characterization

LO<sup>2</sup>

As a results of the complex study of structure, substructure, optical properties, and elemental composition of ZnO, CZTS, and CZT films obtained by pulsed spray pyrolysis technique dependent on the physical (chemical) and technological deposition conditions, it was determined that ZnO nanocrystalline films have an average grain size of *DC* = 25–270 nm and their thickness was *d* = 0.8–1.2 μm, and were formed at *Ts* > 473 K. CZTS continuous films with optimal thickness of *d* = 1.3 μm were deposited at dispersed initial precursor volume *Vs* = 5 ml. It was found that ZnO and CZTS films were polycrystalline in nature, single-phase, and had hexagonal and tetragonal phases, respectively. CZTS samples had a kesterite structure.

It has been shown that in ZnO during the increasing of substrate temperature there is a tendency to the increasing of the CDS; however, in CZTS films, their CSD values were weakly depended on the dispersed solution volume.

Lattice parameters values in ZnO and CZTS films deposited at *Ts* = 623 K, *Vs* = 4 ml were well-correlated to the reference data that confirms their optimal stoichiometry and crystalline quality.

It has been estimated that during the increase of *Ts* the microdeformations level, microstresses, and dislocation density in ZnO films were decreased; in CZTS films, these parameters were weakly dependent on *Vs* .

It has been determined that during the increasing of substrate temperature to 623 K stoichiometry of ZnO layers was improved (*γZnO* = 1.2). It has been shown that optimal for usage in SCs CZTS films, their stoichiometry ratios *γCZTS\_1* = 0.8–0.9, *γCZTS\_2* = 1.1–1.2, *γCZTS\_3* = 0.7 were obtained at *Vs* = 3–4 ml.

Study of the optical characteristics of ZnO films allow to estimate the high values of transmission coefficient *T* = 60–80%. Measured *E*<sup>g</sup> values of ZnO layers were determined in the range of 3.18–3.30 eV and had a complex dependence on *Ts* . During the increase of *Vs* , the values of *Eg* = 1.06–1.30 eV of CZTS layers were approximately approached to the reference data *Eg* = 1.5 eV. Raman spectra analysis of ZnO films confirmed the results of the XRD study, namely decreasing *ε*, *σ*, and *ρ* values during the increase of *Ts* . CZTS films' Raman spectra analysis has confirmed the single-phase nature of condensates. FTIR study indicated the absence of precursor impurities in ZnO films obtained at *Ts* > 573 K.

CZT film spectra (х = 0.32) had a mode LO<sup>2</sup> (ZnTe). In these spectra, intensive peaks corresponded to A1 (Te) and ETO(Te) tellure modes were also determined. CZT film spectra (х = 0.75) have a weak mode A<sup>1</sup> (Te), peaks of LO<sup>1</sup> (CdTe), TO<sup>1</sup> (CdTe), TO<sup>2</sup> (ZnTe), and LO<sup>2</sup> (ZnTe) modes, and also LO<sup>2</sup> (ZnTe) mode resonant replica.

The results of a research study of the ZnO, CZTS, and CZT thin films will be used for the development of the devices, primarily, in third generation high-efficiency thin-film solar cells.

### **Acknowledgements**

This work was supported by the Ministry of the Education and Science of Ukraine (Grants numbers: 0116U002619, 0115U000665c, 0116U006813, and 0117U003929).

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85

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line thin-film Cu<sup>2</sup>

## **Author details**

Oleksandr Dobrozhan, Denys Kurbatov\*, Petro Danilchenko and Anatoliy Opanasyuk

\*Address all correspondence to: dkurbatov@sumdu.edu.ua

Department of Electronics and Computer Technology, Sumy State University, Sumy, Ukraine

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**Acknowledgements**

84 Semiconductors - Growth and Characterization

**Author details**

Ukraine

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**Chapter 5**

Provisional chapter

**E-ALD: Tailoring the Optoeletronic Properties of Metal**

DOI: 10.5772/intechopen.71014

E-ALD: Tailoring the Optoeletronic Properties of Metal

Technological development in nanoelectronics and solar energy devices demands nanostructured surfaces with controlled geometries and composition. Electrochemical atomic layer deposition (E-ALD) is recognized as a valid alternative to vacuum and chemical bath depositions in terms of growth control, quality and performance of semiconducting systems, such as single 2D semiconductors and multilayered materials. This chapter is specific to the E-ALD of metal chalcogenides on Ag single crystals and highlights the electrochemistry for the layer-by-layer deposition of thin films through surface limited reactions (SLRs). Also discussed herein is the theoretical framework of the under potential deposition (UPD), whose thermodynamic treatment open questions to the correct interpretation of the experimental data. Careful design of the E-ALD process allows fine control over both thickness and composition of the deposited layers, thus tailoring the optoelectronic properties of semiconductor compounds. Specifically, the possibility to tune the band gap by varying either the number of deposition cycles or the growth sequence of ternary compounds paves the way

**Chalcogenides on Ag Single Crystals**

Chalcogenides on Ag Single Crystals

Maurizio Passaponti, Francesco Di Benedetto and

toward the formation of advanced photovoltaic materials.

Keywords: E-ALD, Ag(111), thin films, metal chalcogenides, UPD, photovoltaics

Improving the efficiency of the deposition techniques of compound semiconductors could be crucial for the future sustainable micro and nano-electronic industry. A common challenge

> © 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 eproduction in any medium, provided the original work is properly cited.

© 2018 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.

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Maurizio Passaponti, Francesco Di Benedetto

http://dx.doi.org/10.5772/intechopen.71014

Emanuele Salvietti, Andrea Giaccherini, Filippo Gambinossi, Maria Luisa Foresti,

Emanuele Salvietti, Andrea Giaccherini, Filippo Gambinossi, Maria Luisa Foresti,

Massimo Innocenti

Abstract

1. Introduction

and Massimo Innocenti


#### **E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals** E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals

DOI: 10.5772/intechopen.71014

Emanuele Salvietti, Andrea Giaccherini, Filippo Gambinossi, Maria Luisa Foresti, Maurizio Passaponti, Francesco Di Benedetto and Massimo Innocenti Emanuele Salvietti, Andrea Giaccherini, Filippo Gambinossi, Maria Luisa Foresti, Maurizio Passaponti, Francesco Di Benedetto

Additional information is available at the end of the chapter and Massimo Innocenti

http://dx.doi.org/10.5772/intechopen.71014 Additional information is available at the end of the chapter

#### Abstract

[51] Kishore Kumar Y, Uday Bhaskar P, Suresh Babu G. Effect of copper salt and thiourea

ZnSnS<sup>4</sup>

scattering. Journal of Alloys and Compounds. 2011;**509**:7600-7606. DOI: 10.1016/j.

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solar cell applications. Applied Physics Letters. 2011;**98**:181905. DOI: 10.1063/1.3587614

gle crystals by using polarization dependent Raman spectroscopy. Optical Materials.

sis for the identification of secondary phases: Characterization of Cu<sup>2</sup>

[56] Dumcenco D, Huang Y. The vibrational properties study of kesterite Cu<sup>2</sup>

thin films. Thin Solid Films. 2009;**517**:2519-2523. DOI: 10.1016/j.tsf.2008.11.031

thin films and their solar cells. London: Newnes; 2013. 190p.

thin films by spray pyrolysis. Physica

ZnSnS<sup>4</sup>

ZnSnS<sup>4</sup>

layers for

sin-

ZnSnS<sup>4</sup>

films by Raman

concentrations on the formation of Cu<sup>2</sup>

jallcom.2011.04.097

88 Semiconductors - Growth and Characterization

Cu<sup>2</sup>

ZnSnS<sup>4</sup>

ization of Cu<sup>2</sup>

ISBN: 9780123944290

Status Solidi. 2010;**207**:149-156. DOI: 10.1002/pssa.200925194 [52] Fernandes P, Salome P, da Cunha A. Study of polycrystalline Cu<sup>2</sup>

D. 2012;**45**:445103. DOI: 10.1088/0022-3727/45/44/445103

2013;**35**:419-425. DOI: 10.1016/j.optmat.2012.09.031

(Zn,Sn)(S,Se)<sup>4</sup>

Technological development in nanoelectronics and solar energy devices demands nanostructured surfaces with controlled geometries and composition. Electrochemical atomic layer deposition (E-ALD) is recognized as a valid alternative to vacuum and chemical bath depositions in terms of growth control, quality and performance of semiconducting systems, such as single 2D semiconductors and multilayered materials. This chapter is specific to the E-ALD of metal chalcogenides on Ag single crystals and highlights the electrochemistry for the layer-by-layer deposition of thin films through surface limited reactions (SLRs). Also discussed herein is the theoretical framework of the under potential deposition (UPD), whose thermodynamic treatment open questions to the correct interpretation of the experimental data. Careful design of the E-ALD process allows fine control over both thickness and composition of the deposited layers, thus tailoring the optoelectronic properties of semiconductor compounds. Specifically, the possibility to tune the band gap by varying either the number of deposition cycles or the growth sequence of ternary compounds paves the way toward the formation of advanced photovoltaic materials.

Keywords: E-ALD, Ag(111), thin films, metal chalcogenides, UPD, photovoltaics

#### 1. Introduction

Improving the efficiency of the deposition techniques of compound semiconductors could be crucial for the future sustainable micro and nano-electronic industry. A common challenge

© 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 eproduction in any medium, provided the original work is properly cited.

when reviewing current technologies is the lack of reliable compositional control as well as conformal coating of complex geometries. Electrochemical deposition techniques are a lowcost alternative to vacuum evaporation and chemical bath deposition for the direct fabrication of thin films from molecular precursors. Among electrodeposition processes, electrochemical atomic layer deposition (E-ALD) has the advantage to fabricate semiconductors one atomic layer at a time, thus requiring very low energy consumption, diluted solutions, and operating at standard environmental conditions. Fine control of the thin film growth is achieved through the surface limited reaction (SLR), usually referred as under potential deposition (UPD), of the ionic precursors with the electrode substrate or a preceding layer. When the thin film growth can be rigorously considered epitaxial, the E-ALD method is referred as electrochemical atomic layer epitaxy (ECALE) [1]. E-ALD is recognized as a very effective method for the electrodeposition of ultra-thin films of semiconducting materials. In recent years, thin films of binary [2–5] and ternary semiconductors [6–9] were successfully obtained on silver electrodes. Figure 1 illustrates the steps involved in the E-ALD process of a ternary compound. The chosen sequence for the alternate deposition of different elements dictates both the structure and the stoichiometry of the resulting semiconductor compound.

2. Thermodynamics of the UPD process

defined by means of the Nernst equation:

between solvent, surfaces and anions.

energy, Gj

area) and N<sup>j</sup>

For the sake of simplicity and clarity, we will develop the thermodynamics of an epitaxial E-ALD process (ECALE) with the steps constituted by UPD. As a reference we consider the ECALE growth of CdS on Ag(111). From a purely chemical perspective, ECALE exploits UPD to grow the semiconductor compounds on a well-behaving surface. For instance, the oxidative

E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals

where Sad refers to the sulfur adlayer. The UPD provides a surface limited deposition process, which is characterized by a potential more cathodic than the bulk deposition. The latter is

<sup>S</sup>2�=<sup>S</sup> þ

Considering the experimental observation of the potential shift provided by a UPD process, a

where aSad(θ) is the activity of S atoms adsorbed on Ag(111) as a function of the coverage θ.

which is, incorrectly, independent from the activity of the sulfide anions. This heuristic approach, though enabling an intuitive description of the UPD process, does not take into account other terms such the local defects of the electrode surface, and the mutual interactions

A more accurate description of the UPD thermodynamics can be formulated in the framework of the ideal polarized electrode. Following this approach, the substrate (i.e. Ag(111)) on which UPD takes place is in contact with the electrolyte solution (sol) containing S2�, Na+ and X� ions and solvent (H2O). From a theoretical standpoint, the electrolyte solution can be considered in contact with a reference electrode reversible with respect to a general anion X�. The phases relevant for the present analysis are the Ag(111) surface, the solution (sol), and the interphase (IP) between the substrate and sol. Each phase, j, is described thermodynamically by the Gibbs

th species):

; Pj

; AIP; N<sup>j</sup> i

<sup>G</sup><sup>j</sup> <sup>¼</sup> Gj <sup>T</sup><sup>j</sup>

RT <sup>2</sup><sup>F</sup> ln aS aS<sup>2</sup>�

RT

� ES<sup>2</sup>�=<sup>S</sup> <sup>¼</sup> RT

<sup>2</sup><sup>F</sup> ln aSad ð Þ <sup>θ</sup> aS<sup>2</sup>�

> <sup>2</sup><sup>F</sup> ln aSad ð Þ <sup>θ</sup> aS

(temperature), P<sup>j</sup>

(5)

� ⇆ Ag 111 ð Þþ <sup>S</sup><sup>2</sup>� (1)

(2)

http://dx.doi.org/10.5772/intechopen.71014

91

(3)

(4)

(pressure), AIP (interfacial

deposition of an atomic layer of sulfur atoms on top of Ag(111) surface is:

Ag 111 ð Þ� Sad þ 2e

ES<sup>2</sup>�=<sup>S</sup> <sup>¼</sup> <sup>E</sup><sup>0</sup>

<sup>¼</sup> <sup>E</sup><sup>0</sup> <sup>S</sup>2�=<sup>S</sup> þ

heuristic extension of the Nernst equation can be proposed:

ES<sup>2</sup>� <sup>=</sup>SAd

<sup>Δ</sup>EUPD <sup>¼</sup> ES<sup>2</sup>� <sup>=</sup>SAd

, which is a function of the variables T<sup>j</sup>

<sup>i</sup> (amount of matter of the i

The total derivative of the Gibbs energy gives [10]:

Rearranging Eqs. (2) and (3) we obtain the underpotential shift, ΔEUPD [10]:

This chapter highlights the current progress in E-ALD of binary and ternary chalcogenides on Ag(111) from an electrochemical perspective.

Figure 1. Schematic of the E-ALD growth of a ternary semiconductor. The solutions used in each step are indicated in the boxes below.

#### 2. Thermodynamics of the UPD process

when reviewing current technologies is the lack of reliable compositional control as well as conformal coating of complex geometries. Electrochemical deposition techniques are a lowcost alternative to vacuum evaporation and chemical bath deposition for the direct fabrication of thin films from molecular precursors. Among electrodeposition processes, electrochemical atomic layer deposition (E-ALD) has the advantage to fabricate semiconductors one atomic layer at a time, thus requiring very low energy consumption, diluted solutions, and operating at standard environmental conditions. Fine control of the thin film growth is achieved through the surface limited reaction (SLR), usually referred as under potential deposition (UPD), of the ionic precursors with the electrode substrate or a preceding layer. When the thin film growth can be rigorously considered epitaxial, the E-ALD method is referred as electrochemical atomic layer epitaxy (ECALE) [1]. E-ALD is recognized as a very effective method for the electrodeposition of ultra-thin films of semiconducting materials. In recent years, thin films of binary [2–5] and ternary semiconductors [6–9] were successfully obtained on silver electrodes. Figure 1 illustrates the steps involved in the E-ALD process of a ternary compound. The chosen sequence for the alternate deposition of different elements dictates both the structure

This chapter highlights the current progress in E-ALD of binary and ternary chalcogenides on

Figure 1. Schematic of the E-ALD growth of a ternary semiconductor. The solutions used in each step are indicated in the

and the stoichiometry of the resulting semiconductor compound.

Ag(111) from an electrochemical perspective.

90 Semiconductors - Growth and Characterization

boxes below.

For the sake of simplicity and clarity, we will develop the thermodynamics of an epitaxial E-ALD process (ECALE) with the steps constituted by UPD. As a reference we consider the ECALE growth of CdS on Ag(111). From a purely chemical perspective, ECALE exploits UPD to grow the semiconductor compounds on a well-behaving surface. For instance, the oxidative deposition of an atomic layer of sulfur atoms on top of Ag(111) surface is:

$$\text{Ag}(111) - \text{S}\_{ul} + 2\text{e}^- \Leftrightarrow \text{Ag}(111) + \text{S}^{2-} \tag{1}$$

where Sad refers to the sulfur adlayer. The UPD provides a surface limited deposition process, which is characterized by a potential more cathodic than the bulk deposition. The latter is defined by means of the Nernst equation:

$$E\_{s^2-/s} = E\_{s^2-/s}^0 + \frac{RT}{2F} \ln\left(\frac{a\_S}{a\_{S^{2-}}}\right) \tag{2}$$

Considering the experimental observation of the potential shift provided by a UPD process, a heuristic extension of the Nernst equation can be proposed:

$$E\_{s^2-/\_{S\_{Ad}}} = E\_{s^2-/s}^0 + \frac{RT}{2F} \ln\left[\frac{a\_{S\_{ul}}(\theta)}{a\_{S^{2-}}}\right] \tag{3}$$

where aSad(θ) is the activity of S atoms adsorbed on Ag(111) as a function of the coverage θ. Rearranging Eqs. (2) and (3) we obtain the underpotential shift, ΔEUPD [10]:

$$
\Delta E\_{\rm LPD} = E\_{\rm s2-/\_{S\_{\rm AL}}} - E\_{\rm s2-/\_{\rm S}} = \frac{RT}{2F} \ln \left[ \frac{a\_{S\_{\rm al}}(\theta)}{a\_{\rm S}} \right] \tag{4}
$$

which is, incorrectly, independent from the activity of the sulfide anions. This heuristic approach, though enabling an intuitive description of the UPD process, does not take into account other terms such the local defects of the electrode surface, and the mutual interactions between solvent, surfaces and anions.

A more accurate description of the UPD thermodynamics can be formulated in the framework of the ideal polarized electrode. Following this approach, the substrate (i.e. Ag(111)) on which UPD takes place is in contact with the electrolyte solution (sol) containing S2�, Na+ and X� ions and solvent (H2O). From a theoretical standpoint, the electrolyte solution can be considered in contact with a reference electrode reversible with respect to a general anion X�. The phases relevant for the present analysis are the Ag(111) surface, the solution (sol), and the interphase (IP) between the substrate and sol. Each phase, j, is described thermodynamically by the Gibbs energy, Gj , which is a function of the variables T<sup>j</sup> (temperature), P<sup>j</sup> (pressure), AIP (interfacial area) and N<sup>j</sup> <sup>i</sup> (amount of matter of the i th species):

$$\mathbf{G}^{\dagger} = \mathbf{G}^{\dagger} \left( \mathbf{T}^{\dagger}, \mathbf{P}^{\dagger}, \mathbf{A}^{\text{IP}}, \mathbf{N}\_{i}^{j} \right) \tag{5}$$

The total derivative of the Gibbs energy gives [10]:

$$dG = \sum\_{j} dG^{j} = \gamma^{I\mathcal{P}} dA^{I\mathcal{P}} + \sum\_{j,i} S^{j} dT^{j} - V^{j} dP^{j} + \mu^{j} dN\_{i}^{j} \tag{6}$$

where γIP is the surface tension of the interphase, Sj the entropy, Vj the volume and μ<sup>j</sup> the chemical potential.

Let us now consider a solution of Na2S in NaOH buffer solution (i.e. pH = 13) at constant T<sup>j</sup> , Pj . The main charge transfer equilibrium is:

$$\mathcal{S}\_{sol}^{2-} \nrightarrow \mathcal{S}\_{cl} \tag{7}$$

ΓIP <sup>S</sup>2� <sup>¼</sup> <sup>1</sup> RT

ΓIP OH� <sup>¼</sup> <sup>1</sup> RT

> ΓIP <sup>S</sup> <sup>¼</sup> <sup>1</sup> RT

substituting Eqs. (11)–(13) and (18):

ln aS<sup>2</sup>� þ

ΓIP OH� RT qion

ΓIP <sup>S</sup>2� RT qion

<sup>E</sup> <sup>¼</sup> <sup>E</sup><sup>0</sup> <sup>þ</sup>

where

• <sup>Γ</sup>IP <sup>S</sup>2� RT qion

• <sup>Γ</sup>IP OH� RT qion

• <sup>Γ</sup>IP <sup>S</sup> RT qion

• <sup>1</sup> qion γ asol <sup>S</sup>2� ; <sup>a</sup>sol

coverage of the adlayer ( θS).

OH� ; <sup>θ</sup>el <sup>S</sup> ; qion � � � <sup>γ</sup><sup>0</sup>

∂γIP ∂ ln aS<sup>2</sup>� � �

∂γIP ∂ ln aOH� � �

∂γIP ∂ ln θ<sup>S</sup> � �

<sup>E</sup> <sup>¼</sup> <sup>∂</sup>γIP ∂qion � �

<sup>γ</sup> <sup>¼</sup> <sup>γ</sup> asol

ln aOH� þ

• E<sup>0</sup> is the standard potential of the electrode with respect to the reference electrode,

ln aS<sup>2</sup>� takes into account the effect of the activity,

the fractional amount for compounds and the coverage of adlayers,

Hence, the surface tension depends on the activity of the chemical species:

asol OH� ;θel <sup>S</sup> ;<sup>q</sup> ð Þ ion

E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals

asol <sup>S</sup>2� ;θel

asol <sup>S</sup>2� ;asol

asol <sup>s</sup>2� ;asol OH� ;θel S

<sup>S</sup>2� ;asol OH� ;θel

Therefore, upon the integration of Eq. (8), an extension of the Nernst equation can be obtained

ΓIP <sup>S</sup> RT qion

ln aOH� takes into account the effect of the pH, constant in a buffer solution,

h i takes into account the interaction with Ag(111),

Eventually, we should state that Eq. (19), in principle, takes into account all the effects neglected by Eq. (3), thus describing the equilibrium of a system defined by three components, as predicted by the phase rule. Experimental observations suggest that the potential of the UPD process of sulfur depend on the concentration of sulfur atoms (aS<sup>2</sup>�), the pH (aOH�) and the

ln θ<sup>S</sup> takes into account the effect of the surface coverage: θ<sup>S</sup> = 1 for bulk phases, θ<sup>S</sup>

ln θ<sup>S</sup> þ

1 qion γ asol <sup>S</sup>2� ;asol OH� ;θel <sup>S</sup> ;qion � � � <sup>γ</sup><sup>0</sup>

2 4

<sup>S</sup> ;qion � � (15)

http://dx.doi.org/10.5772/intechopen.71014

OH� ;qion � � (16)

� � (17)

<sup>S</sup> ;qion � � (18)

(14)

93

3 5 (19)

where solrefers to the liquid phase andelrefers to the electrode. The corresponding electrocapillary equation is [11, 12]:

$$-d\gamma^{\text{IP}} = \Gamma^{\text{IP}}\_{\text{S}^{2-}} d\mu^{\text{sd}}\_{\text{S}^{2-}} + \Gamma^{\text{IP}}\_{\text{OH}^{-}} d\mu^{\text{sd}}\_{\text{OH}^{-}} + \Gamma^{\text{IP}}\_{\text{S}} d\mu^{\text{el}}\_{\text{S}} - q\_{im} dE \tag{8}$$

where E is the potential referred to the reference electrode and it is defined by means of the chemical potential ΓIP Na<sup>þ</sup> <sup>d</sup>μsol Na<sup>þ</sup> . According to this latter definition, qion is the surface excess of the electric charge density:

$$\eta\_{ion} = \left(\frac{\partial \gamma^{ll}}{\partial E}\right)\_{\left(\mu\_{\mu\_{-}}^{\text{sd}}, \mu\_{\text{OH}^{-1}}^{\text{sd}}, \mu\_{\text{S}}^{\text{d}}\right)}\tag{9}$$

also known as Lippmann equation. ΓIP OH� , ΓIP <sup>S</sup>2� and <sup>Γ</sup>IP <sup>S</sup> are the relative excess of Na<sup>+</sup> and <sup>S</sup>2� with respect to H2O (solvent):

$$
\Gamma\_{S^{2-}}^{IP} = \frac{1}{A^{IP}} \left[ N\_{S^{2-}}^{IP} - \frac{N\_{S^{2-}}^{sol} N\_{H\_2O}^{IP}}{N\_{H\_2O}^{sol}} \right] \tag{10}
$$

It is important to notice that Eq. (10), in the Gibbs theoretical framework, is not a mere change of variable. In fact, the amounts of matter (NIP <sup>i</sup> ) depend on the arbitrary definition of the interface while the relative excess of matter (ΓIP <sup>i</sup> ) is independent of the area and the thickness of the interface. Moreover, the contribution of the formation of the S adlayer is described by the term dγIP. In this context, using the following well-known equation for the chemical potential in the case of complete dissociation:

$$
\mu\_{\mathcal{S}^{2-}}^{\text{sol}} = \mu\_{\mathcal{S}^{2-}}^{0,\text{sol}} + RT \ln a\_{\mathcal{S}^{2-}} \tag{11}
$$

$$
\mu\_{OH^{-}}^{sol} = \mu\_{OH^{-}}^{0,sol} + RT \ln a\_{OH^{-}} \tag{12}
$$

$$
\mu\_{\rm S}^{el} = \mu\_{\rm S}^{el} + RT \ln \Theta\_{\rm S} \tag{13}
$$

where a is the activity in the liquid phase and θ is the fractional amount in the solid phase. In this framework the relationship between activities and surface tension is well represented by the following:

E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals http://dx.doi.org/10.5772/intechopen.71014 93

$$T\_{S^{2-}}^{IP} = \frac{1}{RT} \left(\frac{\partial \boldsymbol{\gamma}^{IP}}{\partial \ln \boldsymbol{a}\_{S^{2-}}}\right)\_{\left(\boldsymbol{a}\_{\partial \mathcal{H}^{-}}^{ul}, \boldsymbol{\theta}\_{S}^{ul}, \boldsymbol{q}\_{\partial \mathcal{u}}\right)}\tag{14}$$

$$\Gamma\_{OH^{-}}^{IP} = \frac{1}{RT} \left( \frac{\partial \gamma^{IP}}{\partial \ln a\_{OH^{-}}} \right)\_{\left(d\_{S^{2-}}^{\text{self}}, \partial\_{S}^{d}, q\_{in}\right)} \tag{15}$$

$$\Gamma\_S^{\rm IP} = \frac{1}{RT} \left( \frac{\partial \gamma^{\rm IP}}{\partial \ln \Theta\_S} \right)\_{\left(\mathcal{A}\_{S^{2-}}^{\rm{cal}}, \mathcal{A}\_{\rm OI^{-}}^{\rm{cal}}, \eta\_{\rm in}\right)} \tag{16}$$

$$E = \left(\frac{\partial \gamma^{l\mathcal{P}}}{\partial q\_{im}}\right)\_{\left(a\_{s^{2-}}^{\text{self}}, a\_{OH^{-}}^{\text{self}}, \mathcal{O}\_{S}^{d}\right)}\tag{17}$$

Hence, the surface tension depends on the activity of the chemical species:

$$\mathcal{V} = \mathcal{V} \begin{pmatrix} & & \\ a\_{\mathcal{S}^2 -}^{sl}, a\_{\mathcal{O}H^-}^{sl}, \theta\_{\mathcal{S}}^l, q\_{im} \end{pmatrix} \tag{18}$$

Therefore, upon the integration of Eq. (8), an extension of the Nernst equation can be obtained substituting Eqs. (11)–(13) and (18):

$$E = E^{0} + \frac{\Gamma\_{S^{3-}}^{\rm IP}RT}{q\_{im}} \ln a\_{S^{-}} + \frac{\Gamma\_{OH^{-}}^{\rm IP}RT}{q\_{im}} \ln a\_{OH^{-}} + \frac{\Gamma\_{S}^{\rm IP}RT}{q\_{im}} \ln \Theta\_{S} + \frac{1}{q\_{im}} \left[ \gamma \left( \frac{1}{a\_{S^{-}}^{\rm el} \cdot a\_{OH^{-}}^{\rm el} \cdot \theta\_{S}^{d} q\_{im}} \right) - \gamma^{\rm O} \right] \tag{19}$$

where

dG <sup>¼</sup> <sup>X</sup>

The main charge transfer equilibrium is:

92 Semiconductors - Growth and Characterization

chemical potential.

equation is [11, 12]:

chemical potential ΓIP

electric charge density:

j

�dγIP <sup>¼</sup> <sup>Γ</sup>IP

Na<sup>þ</sup> <sup>d</sup>μsol

of variable. In fact, the amounts of matter (NIP

interface while the relative excess of matter (ΓIP

in the case of complete dissociation:

the following:

also known as Lippmann equation. ΓIP

with respect to H2O (solvent):

<sup>S</sup>2� <sup>d</sup>μsol

ΓIP <sup>S</sup>2� <sup>¼</sup> <sup>1</sup>

μsol

μsol

<sup>S</sup>2� <sup>¼</sup> <sup>μ</sup><sup>0</sup>,sol

OH� <sup>¼</sup> <sup>μ</sup><sup>0</sup>,sol

μel <sup>S</sup> <sup>¼</sup> <sup>μ</sup>el

dGj <sup>¼</sup> <sup>γ</sup>IPdAIP <sup>þ</sup><sup>X</sup>

j,i Sj

where γIP is the surface tension of the interphase, Sj the entropy, Vj the volume and μ<sup>j</sup> the

Let us now consider a solution of Na2S in NaOH buffer solution (i.e. pH = 13) at constant T<sup>j</sup>

S<sup>2</sup>�

<sup>S</sup>2� <sup>þ</sup> <sup>Γ</sup>IP

qion <sup>¼</sup> <sup>∂</sup>γIP ∂E � �

OH� , ΓIP

AIP <sup>N</sup>IP

where solrefers to the liquid phase andelrefers to the electrode. The corresponding electrocapillary

OH� <sup>d</sup>μsol

where E is the potential referred to the reference electrode and it is defined by means of the

μsol <sup>s</sup>2� ;μsol OH� ;μel S

<sup>S</sup>2� and <sup>Γ</sup>IP

<sup>S</sup>2� � <sup>N</sup>sol

It is important to notice that Eq. (10), in the Gibbs theoretical framework, is not a mere change

of the interface. Moreover, the contribution of the formation of the S adlayer is described by the term dγIP. In this context, using the following well-known equation for the chemical potential

where a is the activity in the liquid phase and θ is the fractional amount in the solid phase. In this framework the relationship between activities and surface tension is well represented by

" #

<sup>S</sup>2� <sup>N</sup>IP H2O Nsol H2O

OH� <sup>þ</sup> <sup>Γ</sup>IP

<sup>S</sup> <sup>d</sup>μel

Na<sup>þ</sup> . According to this latter definition, qion is the surface excess of the

dT<sup>j</sup> � <sup>V</sup><sup>j</sup>

dPj <sup>þ</sup> <sup>μ</sup><sup>j</sup>

dN<sup>j</sup>

sol ⇆ Sel (7)

� � (9)

<sup>i</sup> ) depend on the arbitrary definition of the

<sup>i</sup> ) is independent of the area and the thickness

<sup>S</sup>2� þ RT ln aS<sup>2</sup>� (11)

OH� þ RT ln aOH� (12)

<sup>S</sup> þ RT ln θ<sup>S</sup> (13)

<sup>S</sup> are the relative excess of Na<sup>+</sup> and <sup>S</sup>2�

<sup>S</sup> � qiondE (8)

<sup>i</sup> (6)

, Pj .

(10)


Eventually, we should state that Eq. (19), in principle, takes into account all the effects neglected by Eq. (3), thus describing the equilibrium of a system defined by three components, as predicted by the phase rule. Experimental observations suggest that the potential of the UPD process of sulfur depend on the concentration of sulfur atoms (aS<sup>2</sup>�), the pH (aOH�) and the coverage of the adlayer ( θS).

#### 3. E-ALD of metal chalcogenides

E-ALD of metal chalcogenides involves sequential SLR of metal and chalcogenides ions. Except for few cases, the choice of chalcogenides as first layer is preferred for its higher affinity for silver and to avoid alloy formation with the electrode substrate [2]. In the E-ALD methodology a monolayer of the compound is obtained by alternating underpotential deposition of the metallic element with the underpotential deposition of the non-metallic element in a cycle, so the thickness of deposited film is a function of the numbers of deposition cycles. Figure 2a shows the cyclic voltammogramm of 0.5 mM Na2SeO3 on Ag(111) and of ZnSO4 on Se-covered Ag(111). The cathode peak CSe at 0.95 V corresponds to the bulk reduction of the Se. At more negative potentials hydrogen evolution prevents the detection of the Sead reduction. Zn adlayers deposited on a Se-covered Ag(111) show a CV profile shifted to more positive potential with a reduction peak CZn at 0.65 V. To prevent Se dissolution, the operating potential for the surface limited deposition of Zn on Ag/Se should be more positive than CSe. Anodic stripping analysis as a function of the applied potential allow to choosing the optimal conditions for the UPD process, which in this case is in the range 0.95 < E < 0.8 V, where a plateau at Q = 60 μC cm<sup>2</sup> is observed (see Figure 2b). Alternate depositions of both chalcogenide and metal can be repeated as many times as necessary to form a compound with the desired thickness and composition. Figure 2c shows the charges involved in the layer-by-layer growth of Zn and Se in a 1:1 ratio.

outlet, in the front. The solution is pushed to the distribution valve by the N2 overpressure. The inlet of the solution is blocked by a piston of the distribution valve, which is tightly held in place by a spring. The piston can be raised by opening the solenoid valve and sending compressed air at 6 atm, which is higher than the pressure exerted by the spring. The piston is raised enough to allow the solution to flow into the cell. Different solutions are pushed to the cell following the desired sequence by acting on the corresponding solenoid valves. The nitrogen pressure exerted on the solutions determines a flow rate of about 1 mL s<sup>1</sup>

E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals

http://dx.doi.org/10.5772/intechopen.71014

The electrochemical cell (Figure 3b) is a Teflon cylinder with a height of 40 mm, an internal diameter of 10 mm and external diameter of 50 mm. The working electrode at the bottom of the cylinder and the counter electrode on top delimit the electrochemical cell volume (1.6 mL). The solution inlet and outlet are placed on the side walls of the cylinder; for hydrodynamic reasons, the inlet is inclined. The working electrode was silver disks cut according to the Bridgman technique [13]. The counter electrode was a gold foil, and the reference electrode was an Ag/AgCl (saturated KCl) placed in the outlet tubing. Leakage is avoided by pressing both the working and the counter electrode against a suitable Viton® O-ring. All potentials reported in the paper are

E-ALD technique has been successfully used to fabricate ultrathin films of metal sulfides on silver electrodes by alternating the underpotential deposition of metal and sulfur. These compounds include cadmium sulfide (CdS) [2, 14, 15], zinc sulfide (ZnS) [2], nickel sulfide (NiS) [4], lead sulfide (PbS) [16], copper sulfides (CuxS) [17, 18] and tin sulfides (SnxSy) [19]. A typical E-ALD cycle includes the underpotential deposition of sulfur followed by the surface limited reaction (SLR) of metal on S-covered Ag. The UPD of sulfur on crystalline and polycrystalline silver have been extensively investigated in the past [20–22]. Electrochemical measurements on Ag(110), Ag(100) and Ag(111) show that sulfur UPD deposition processes differ significantly on

operations are carried out under computer control.

Figure 3. (a) Distribution valve and (b) electrochemical cell.

quoted with respect to the Ag/AgCl (saturated KCl) reference electrode.

5. E-ALD of binary MxSy semiconductors on Ag(111)

. All

95

Figure 2. (a) Cyclic voltammograms of (—) Se on Ag(111) and (—) Zn on Se-covered Ag(111) in ammonia buffer solutions (pH 9.2); (b) charges involved in the stripping of Zn, underpotentially deposited on Ag(111)/Sead, as a function of the deposition potential; (c) charges involved in the striping of (◯) Se and (□) Zn as a function of the number of E-ALD cycles. The solid lines represent the linear fit to the data.

#### 4. Flow cell apparatus and experimental conditions

E-ALD thin films of metal chalcogenides were grown on silver single crystal disks using an automated deposition apparatus consisting of Pyrex solution reservoirs, solenoid valves, a distribution valve and a flow-cell [2]. The solutions contained in the Pyrex reservoirs are previously degassed and then constantly kept under a nitrogen pressure p(N2) = 0.3 atm. Figure 3a shows the distribution valve with seven solution inlets, in the top, and one solution E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals http://dx.doi.org/10.5772/intechopen.71014 95

Figure 3. (a) Distribution valve and (b) electrochemical cell.

3. E-ALD of metal chalcogenides

94 Semiconductors - Growth and Characterization

growth of Zn and Se in a 1:1 ratio.

The solid lines represent the linear fit to the data.

4. Flow cell apparatus and experimental conditions

E-ALD thin films of metal chalcogenides were grown on silver single crystal disks using an automated deposition apparatus consisting of Pyrex solution reservoirs, solenoid valves, a distribution valve and a flow-cell [2]. The solutions contained in the Pyrex reservoirs are previously degassed and then constantly kept under a nitrogen pressure p(N2) = 0.3 atm. Figure 3a shows the distribution valve with seven solution inlets, in the top, and one solution

Figure 2. (a) Cyclic voltammograms of (—) Se on Ag(111) and (—) Zn on Se-covered Ag(111) in ammonia buffer solutions (pH 9.2); (b) charges involved in the stripping of Zn, underpotentially deposited on Ag(111)/Sead, as a function of the deposition potential; (c) charges involved in the striping of (◯) Se and (□) Zn as a function of the number of E-ALD cycles.

E-ALD of metal chalcogenides involves sequential SLR of metal and chalcogenides ions. Except for few cases, the choice of chalcogenides as first layer is preferred for its higher affinity for silver and to avoid alloy formation with the electrode substrate [2]. In the E-ALD methodology a monolayer of the compound is obtained by alternating underpotential deposition of the metallic element with the underpotential deposition of the non-metallic element in a cycle, so the thickness of deposited film is a function of the numbers of deposition cycles. Figure 2a shows the cyclic voltammogramm of 0.5 mM Na2SeO3 on Ag(111) and of ZnSO4 on Se-covered Ag(111). The cathode peak CSe at 0.95 V corresponds to the bulk reduction of the Se. At more negative potentials hydrogen evolution prevents the detection of the Sead reduction. Zn adlayers deposited on a Se-covered Ag(111) show a CV profile shifted to more positive potential with a reduction peak CZn at 0.65 V. To prevent Se dissolution, the operating potential for the surface limited deposition of Zn on Ag/Se should be more positive than CSe. Anodic stripping analysis as a function of the applied potential allow to choosing the optimal conditions for the UPD process, which in this case is in the range 0.95 < E < 0.8 V, where a plateau at Q = 60 μC cm<sup>2</sup> is observed (see Figure 2b). Alternate depositions of both chalcogenide and metal can be repeated as many times as necessary to form a compound with the desired thickness and composition. Figure 2c shows the charges involved in the layer-by-layer

outlet, in the front. The solution is pushed to the distribution valve by the N2 overpressure. The inlet of the solution is blocked by a piston of the distribution valve, which is tightly held in place by a spring. The piston can be raised by opening the solenoid valve and sending compressed air at 6 atm, which is higher than the pressure exerted by the spring. The piston is raised enough to allow the solution to flow into the cell. Different solutions are pushed to the cell following the desired sequence by acting on the corresponding solenoid valves. The nitrogen pressure exerted on the solutions determines a flow rate of about 1 mL s<sup>1</sup> . All operations are carried out under computer control.

The electrochemical cell (Figure 3b) is a Teflon cylinder with a height of 40 mm, an internal diameter of 10 mm and external diameter of 50 mm. The working electrode at the bottom of the cylinder and the counter electrode on top delimit the electrochemical cell volume (1.6 mL). The solution inlet and outlet are placed on the side walls of the cylinder; for hydrodynamic reasons, the inlet is inclined. The working electrode was silver disks cut according to the Bridgman technique [13]. The counter electrode was a gold foil, and the reference electrode was an Ag/AgCl (saturated KCl) placed in the outlet tubing. Leakage is avoided by pressing both the working and the counter electrode against a suitable Viton® O-ring. All potentials reported in the paper are quoted with respect to the Ag/AgCl (saturated KCl) reference electrode.

#### 5. E-ALD of binary MxSy semiconductors on Ag(111)

E-ALD technique has been successfully used to fabricate ultrathin films of metal sulfides on silver electrodes by alternating the underpotential deposition of metal and sulfur. These compounds include cadmium sulfide (CdS) [2, 14, 15], zinc sulfide (ZnS) [2], nickel sulfide (NiS) [4], lead sulfide (PbS) [16], copper sulfides (CuxS) [17, 18] and tin sulfides (SnxSy) [19]. A typical E-ALD cycle includes the underpotential deposition of sulfur followed by the surface limited reaction (SLR) of metal on S-covered Ag. The UPD of sulfur on crystalline and polycrystalline silver have been extensively investigated in the past [20–22]. Electrochemical measurements on Ag(110), Ag(100) and Ag(111) show that sulfur UPD deposition processes differ significantly on the three silver facets and in situ STM experiments have evidenced the presence of differently ordered sulfur structures depending on substrate orientation. The formation of the first layer of S on Ag(100) and Ag(110) occurs at a potential of E = 0.8 V in pH 13 solutions, whereas on Ag (111) it was obtained at E = 0.68 V in ammonia buffer (pH 9.6). Proceeding toward more positive potentials in the presence of sulfide ions resulted in bulk sulfur deposition. Cyclic voltammogramms performed in Na2S solutions revealed two distinct behaviors. While on Ag (100) only two broad anodic peaks at 1.32 and 1.15 V were observed, cycling the potential on the other two faces resulted in a more complex behavior, with a sharp anodic peak occurring at E = 1.06 and E = 0.78 on Ag(110) and Ag(111), respectively. The charges associated with the UPD deposition of S in the first layer, as calculated from chronocoulometry experiments, were 137μC cm<sup>2</sup> for Ag(100), 163μC cm<sup>2</sup> for Ag(110) and 189μC cm<sup>2</sup> for Ag(111). Although the growth of MxSy on silver electrodes follows an oxidative-reductive behavior in all cases, SLR of metal layers depends on the semiconductor type and solution conditions.

5.2. E-ALD of CdS, ZnS, SnxSy and CuxS

[28, 29].

for Zn and 75 μC cm<sup>2</sup> for S [2].

thus confirming a surface limited process.

subsequent conproportionation reaction:

of E-ALD due to favorable electrochemical characteristics.

Cadmium, tin, zinc, and copper are metals that form strong interactions with sulfur ions, thus generating binary semiconductors with peculiar transport and electrical properties. Cadmium and zinc sulfides have been among the first binary semiconductors to be deposited by means

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UPD deposition of CdS has been investigated on silver single crystals exhibiting different orientations, that is, Ag(111), Ag(110) and Ag(111) [14, 15].The well-defined structures of SUPD formed on silver electrodes in Na2S ammonia buffer solutions drive the epitaxial deposition of the subsequent Cd layer. After washing out the sulfur ions in excess, UPD of cadmium is performed by holding the potential at 0.68 V on S-Ag(111) and at E = 0.6 V on the other two faces. The procedure used to obtain further alternate S and Cd layers is identical with that used for the Ag/S/Cd structure. The average values of the charge deposited in each E-ALD cycle were 103 μC cm<sup>2</sup> for Ag(100), 85 μC cm<sup>2</sup> for Ag(110) and 70 μC cm<sup>2</sup> for Ag(111). Ex situ XPS analysis confirmed the presence of cadmium and sulfur in a 1:1 stoichiometric ratio

Similarly to cadmium sulfide, E-ALD of ZnS thin film were grown on silver electrodes by first depositing sulfur at E = 0.70 V from Na2S solution, and then injecting ZnSO4 solution while keeping the electrode at the same potential to deposit Zn underpotentially. Plots of the charges for Zn and S measured in the stripping of ZnS deposits were linear, with a slope of 67 μC cm<sup>2</sup>

Differently from CdS and ZnS, the other two metal sulfides (CuxS and SnxSy) exist in different stoichiometric ratios and morphologies; their interest in E-ALD growth relies on the tunable

The surface limited reaction of SnxSy has been thoroughly studied by Innocenti et al. [19] both on bare and on S-covered silver substrates. Electrodeposition of tin on bare silver showed two anodic peaks, Ec1 = 0.70 V and Ec2 = 0.48 V, the latter ascribed to the formation of Sn(IV) hydroxides. Differently, on S-covered Ag(111) the reduction peak was seen at lower potentials, Ec = 0.61 V, thus suggesting true underpotential deposition mechanism. Independently from the deposition time, the charge involved in the oxidative process remained nearly constant,

Thin films of copper sulfides were fabricated on silver substrates through E-ALD. Although surface limited the layer-by-layer growth of CuxS was found not to be a true UPD process like for the other metal sulfides. As reported by Innocenti et al. [17] the electrochemistry of copper on S-covered Ag(111) is quite complex. Cyclic voltammograms as obtained by sweeping the potential between 0.05 and 0.55 V in 1 mM Cu(II) solutions in ammonia buffer revealed the presence of two cathodic peaks, E1 = 0.39 V and E2 = 0.42 V. While the latter was easily associated to the bulk reduction of Cu(II) to Cu(0), the nature of the first cathodic peak is still under debate, except for the fact that it precedes bulk deposition. More-in-depth electrochemical analysis suggests the process is surface limited in the range 0.3 to 0.38 V and it involves the formation of Cu(0) through the two-step reduction of cupric tetra-amino complex and the

transport and electronic properties by changes in composition, x and y [30, 31].

#### 5.1. E-ALD of PbS and NiS

Lead sulfide (PbS) and nickel sulfide (NiS) are binary semiconductors that have received considerable attention for a variety of applications, such as detectors [23], sensing materials [24] and solar cells [25]. A complete electrochemical study of PbS multilayers has been reported by Fernandes et al. [16]. The Pbad cannot be formed on S-covered silver electrodes because of partial redissolution of sulfur. So the first layer is the PbUPD deposited on Ag(111), this was obtained from 5.0 mM Pb(NO3)2 solutions in acetic buffer at pH 5 by scanning the potential from 0.2 to 0.45 V. Two well-defined peaks were observed at 0.35 and 0.29 V for the deposition and the dissolution of Pb monolayer, respectively. Next, the underpotential deposition of S on Pb-covered Ag(111) was obtained by scanning the potential from 1.0 to 0.70 V in 2.5 mM Na2S solutions in ammonia buffer. The constancy of the anodic stripping of Pb deposited at 0.45 V and at different accumulation times ensures the process is surface limited, giving rise to a charge of 332 μC cm<sup>2</sup> for the first layer. Successive sulfur and lead layers resulted in a linear growth with an average charge per cycle of approximately 83 μC cm<sup>2</sup> . Morphological analysis by ex-situ AFM measurements revealed the deposits consisted of homogeneous films of PbS small clusters.

Differently from lead, underpotential deposition of nickel on bare Ag(111) is not possible due to weak adhesion with the electrode substrate and competing surface phase transformations [26]. On the contrary nickel presents a well-defined surface limited reaction on S-covered Ag (111), showing cathodic and anodic peaks at 0.52 and 0.22 V, respectively. UPD layers of nickel were obtained from NiCl2 in boric acid solutions (pH 6.5) at E = 0.6 V. The amount of sulfur, as determined by separate anodic stripping experiments, was found to increase linearly with the number of the deposition cycles. Conversely, the stripping of nickel was less precise due to the formation of oxide and hydroxide films with the increase in layer number [27]. Despite these limitations the charges of both Ni and S showed a quasi-linear layer-by-layer growth with slopes of 53 μC cm<sup>2</sup> for Ni and 58 μC cm<sup>2</sup> for S. Morphological analysis by AFM indicated a decrease of average roughness with aging, thus suggesting the formation of a passivation layer with time, which was later confirmed by XPS analysis.

## 5.2. E-ALD of CdS, ZnS, SnxSy and CuxS

the three silver facets and in situ STM experiments have evidenced the presence of differently ordered sulfur structures depending on substrate orientation. The formation of the first layer of S on Ag(100) and Ag(110) occurs at a potential of E = 0.8 V in pH 13 solutions, whereas on Ag (111) it was obtained at E = 0.68 V in ammonia buffer (pH 9.6). Proceeding toward more positive potentials in the presence of sulfide ions resulted in bulk sulfur deposition. Cyclic voltammogramms performed in Na2S solutions revealed two distinct behaviors. While on Ag (100) only two broad anodic peaks at 1.32 and 1.15 V were observed, cycling the potential on the other two faces resulted in a more complex behavior, with a sharp anodic peak occurring at E = 1.06 and E = 0.78 on Ag(110) and Ag(111), respectively. The charges associated with the UPD deposition of S in the first layer, as calculated from chronocoulometry experiments, were 137μC cm<sup>2</sup> for Ag(100), 163μC cm<sup>2</sup> for Ag(110) and 189μC cm<sup>2</sup> for Ag(111). Although the growth of MxSy on silver electrodes follows an oxidative-reductive behavior in all cases, SLR of

Lead sulfide (PbS) and nickel sulfide (NiS) are binary semiconductors that have received considerable attention for a variety of applications, such as detectors [23], sensing materials [24] and solar cells [25]. A complete electrochemical study of PbS multilayers has been reported by Fernandes et al. [16]. The Pbad cannot be formed on S-covered silver electrodes because of partial redissolution of sulfur. So the first layer is the PbUPD deposited on Ag(111), this was obtained from 5.0 mM Pb(NO3)2 solutions in acetic buffer at pH 5 by scanning the potential from 0.2 to 0.45 V. Two well-defined peaks were observed at 0.35 and 0.29 V for the deposition and the dissolution of Pb monolayer, respectively. Next, the underpotential deposition of S on Pb-covered Ag(111) was obtained by scanning the potential from 1.0 to 0.70 V in 2.5 mM Na2S solutions in ammonia buffer. The constancy of the anodic stripping of Pb deposited at 0.45 V and at different accumulation times ensures the process is surface limited, giving rise to a charge of 332 μC cm<sup>2</sup> for the first layer. Successive sulfur and lead layers resulted in a linear growth with an average charge per cycle of approximately 83 μC cm<sup>2</sup>

Morphological analysis by ex-situ AFM measurements revealed the deposits consisted of homo-

Differently from lead, underpotential deposition of nickel on bare Ag(111) is not possible due to weak adhesion with the electrode substrate and competing surface phase transformations [26]. On the contrary nickel presents a well-defined surface limited reaction on S-covered Ag (111), showing cathodic and anodic peaks at 0.52 and 0.22 V, respectively. UPD layers of nickel were obtained from NiCl2 in boric acid solutions (pH 6.5) at E = 0.6 V. The amount of sulfur, as determined by separate anodic stripping experiments, was found to increase linearly with the number of the deposition cycles. Conversely, the stripping of nickel was less precise due to the formation of oxide and hydroxide films with the increase in layer number [27]. Despite these limitations the charges of both Ni and S showed a quasi-linear layer-by-layer growth with slopes of 53 μC cm<sup>2</sup> for Ni and 58 μC cm<sup>2</sup> for S. Morphological analysis by AFM indicated a decrease of average roughness with aging, thus suggesting the formation of a

passivation layer with time, which was later confirmed by XPS analysis.

.

metal layers depends on the semiconductor type and solution conditions.

5.1. E-ALD of PbS and NiS

96 Semiconductors - Growth and Characterization

geneous films of PbS small clusters.

Cadmium, tin, zinc, and copper are metals that form strong interactions with sulfur ions, thus generating binary semiconductors with peculiar transport and electrical properties. Cadmium and zinc sulfides have been among the first binary semiconductors to be deposited by means of E-ALD due to favorable electrochemical characteristics.

UPD deposition of CdS has been investigated on silver single crystals exhibiting different orientations, that is, Ag(111), Ag(110) and Ag(111) [14, 15].The well-defined structures of SUPD formed on silver electrodes in Na2S ammonia buffer solutions drive the epitaxial deposition of the subsequent Cd layer. After washing out the sulfur ions in excess, UPD of cadmium is performed by holding the potential at 0.68 V on S-Ag(111) and at E = 0.6 V on the other two faces. The procedure used to obtain further alternate S and Cd layers is identical with that used for the Ag/S/Cd structure. The average values of the charge deposited in each E-ALD cycle were 103 μC cm<sup>2</sup> for Ag(100), 85 μC cm<sup>2</sup> for Ag(110) and 70 μC cm<sup>2</sup> for Ag(111). Ex situ XPS analysis confirmed the presence of cadmium and sulfur in a 1:1 stoichiometric ratio [28, 29].

Similarly to cadmium sulfide, E-ALD of ZnS thin film were grown on silver electrodes by first depositing sulfur at E = 0.70 V from Na2S solution, and then injecting ZnSO4 solution while keeping the electrode at the same potential to deposit Zn underpotentially. Plots of the charges for Zn and S measured in the stripping of ZnS deposits were linear, with a slope of 67 μC cm<sup>2</sup> for Zn and 75 μC cm<sup>2</sup> for S [2].

Differently from CdS and ZnS, the other two metal sulfides (CuxS and SnxSy) exist in different stoichiometric ratios and morphologies; their interest in E-ALD growth relies on the tunable transport and electronic properties by changes in composition, x and y [30, 31].

The surface limited reaction of SnxSy has been thoroughly studied by Innocenti et al. [19] both on bare and on S-covered silver substrates. Electrodeposition of tin on bare silver showed two anodic peaks, Ec1 = 0.70 V and Ec2 = 0.48 V, the latter ascribed to the formation of Sn(IV) hydroxides. Differently, on S-covered Ag(111) the reduction peak was seen at lower potentials, Ec = 0.61 V, thus suggesting true underpotential deposition mechanism. Independently from the deposition time, the charge involved in the oxidative process remained nearly constant, thus confirming a surface limited process.

Thin films of copper sulfides were fabricated on silver substrates through E-ALD. Although surface limited the layer-by-layer growth of CuxS was found not to be a true UPD process like for the other metal sulfides. As reported by Innocenti et al. [17] the electrochemistry of copper on S-covered Ag(111) is quite complex. Cyclic voltammograms as obtained by sweeping the potential between 0.05 and 0.55 V in 1 mM Cu(II) solutions in ammonia buffer revealed the presence of two cathodic peaks, E1 = 0.39 V and E2 = 0.42 V. While the latter was easily associated to the bulk reduction of Cu(II) to Cu(0), the nature of the first cathodic peak is still under debate, except for the fact that it precedes bulk deposition. More-in-depth electrochemical analysis suggests the process is surface limited in the range 0.3 to 0.38 V and it involves the formation of Cu(0) through the two-step reduction of cupric tetra-amino complex and the subsequent conproportionation reaction:

$$\text{Cu(NH}\_3\text{)}\_4^{2+} + e^- \rightarrow \text{Cu(NH}\_3\text{)}\_2^+ + 2\text{NH}\_3 \tag{20}$$

$$\text{Cu(NH}\_3\text{)}\_2^+ + e^- \rightarrow \text{Cu(0)} + 2\text{NH}\_3\tag{21}$$

by AFM measurements, were found to be homogeneous and similar to the bare Ag(111). Conversely, the band gap values, carried out by diffuse reflectance spectroscopy, decreased with the increase in thickness and Cu/Sn ratio ranging from 2.12 for n = 60, j = 2–2.43 eV for

As for the Cu-Sn-S system the experimental conditions to grow ternary cadmium and zinc sulfides on Ag(111) with the E-ALD were by alternating the underpotential deposition of the corresponding binaries (CdS and ZnS) [36]. The potential chosen for the deposition of S and Zn was 0.65 V, whereas different series of experiments were carried out depositing Cd in correspondence to the first (0.5 V) or to the second UPD (0.65 V). The authors investigated the electrodeposition of CdxZn1-xS thin films exploiting different sequences of E-ALD cycles, that is,

increases linearly with the number of deposition cycles. Yet, the slope of the plot decreases while increasing the number of ZnS cycles per CdS cycles, that is, 89, 67 and 52 μC cm<sup>2</sup> for j = 1, 2 and 5, which reflect the higher percentage of Zn in the deposit. The chemical composition of CdxZn1-xS thin films, as analyzed by XPS have highlighted the nominal stoichiometry is not respected leading to Zn/Cu ratio equal to 1/3, 1/2 and 2 for j = 1, 2 and 5, respectively. However, regardless of the stoichiometry of the obtained ternary compound, the charge involved in cathodic stripping was equal to the charge involved in the anodic one, thus indicating the right

The low contribution of Zn in CdxZn1-xS compounds was also found to occur in the Cu-Zn-S system. Ternary CuxZnyS compounds were obtained through the E-ALD method by alternating deposition of CuxS and ZnS layers. Innocenti et al. [8, 38] studied the electrochemical and compositional behavior of CdxZn1-xS thin films by applying the general sequence Ag/S/

compounds UPD sulfur layers were obtained by keeping the electrode potential at E = 0.68 V for either on bare Ag(111) or on Ag(111) already covered by a metal layer. Deposition of Cu and Zn occurred in ammonia buffer at 0.37 and 0.85 V, respectively. Stripping analysis revealed the charge involved in (Cu + Zn) deposition followed a nonlinear increase with the number of cycles. Moreover, samples containing higher percentage of Zn had lower slope, thus indicating a lower extent of deposition in each cycle; a Cu/Zn ratio of about 6 was found for j = 1, k = 1, and n = 40. From the plot of charges against the number of deposition cycles (n) we expect a ZnS:CuS equimolar ratio for n = 13. Zn deficiency upon increasing the cycles is due to partial redissolution of zinc during the E-ALD of copper, thus causing a rearrangement in the adlayers. Independently from the Cu/Zn ratio, ex-situ investigations highlighted at least two prevailing morphologies, the first one homogeneously covering the Ag(111) surface, and the second one consisting of random network of

The E-ALD of metal selenides (MxSey) on monocrystalline surfaces usually starts with the deposition of a selenium monolayer by using selenite, Se(IV), as a precursor salt. Differently

/(Zn/S)k)]n with j = 1, k = 1, 5, 9, and 1 < n < 60. As for other binary and ternary

]n with k = 1, j = 1, 2, 5 and 1 < n < 20. The charge involved in the stripping

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n = 20, j = 1.

Ag/S[(Cd/S)k/(Zn/S)j

1:1 stoichiometric ratio [35].

nanowires of variable length [8].

6. E-ALD of binary MxSey semiconductors on Ag(111)

[(Cu/S)j

$$\text{Cu(NH}\_3\text{)}\_4^{2+} + \text{Cu(0)} \rightarrow 2\text{Cu(NH}\_3\text{)}\_2^{+} \tag{22}$$

The amount of Cu deposited in a given number of cycles was determined by measuring the charge involved in the anodic stripping. The authors found the charges were linearly increasing with the number of deposition cycles with a slope of 44 μC cm�<sup>2</sup> . XPS results confirmed the valence states of copper and sulfur as Cu(I) and S(�II), respectively, although a possible fraction of S(�I) in the form of disulfide anion was not excluded. The experimental Cu/S ratio observed in the XPS characterization was later attributed to the covellite phase, where positive holes allow Cu ions to be stabilized in their monovalent state [32]. The morphological analysis by AFM was able to evidence the low roughness values of the deposits, thus confirming the high homogeneity and good quality of the thin film obtained.

#### 5.3. E-ALD of ternary MxN1-xS

The interest in cadmium, zinc, tin, and copper sulfides has increased in the last 10 years due to the possibility to fabricate ternary semiconductor compounds, thus allowing fine control over the band gap energy of solar cell devices [32–34]. E-ALD method has been successfully employed to grow ternary materials such as CdxZn1-xS [35, 36], CuxZn1-xS [5] and CuxSnySz [33]. These semiconductors were prepared by sequential deposition of the corresponding binaries; for instance, alternate deposition of CuxS and SnxSy was carried out to form CuxSnySz. Because of the large variety of the possible (x:y:z) combinations research on E-ALD of multinary kesterite group thin films, although challenging, is quite promising for the development of non-linear electro-optic devices and photovoltaic cells. Depending on the adopted sequence profile only certain combinations were attainable, thus limiting the possible metalto-metal stoichiometries. The electrochemical behavior of CuxSnySz, CuxZn1-xS and CdxZn1-xS thin films is separately discussed below.

Di Benedetto et al. [33] investigated the electrodeposition of CuxSnySz thin films exploiting different sequences of E-ALD cycles, that is Ag/S[(Cu/S)k/(Sn/S)j ]n with k = 1, j = 1, 2 and 1 < n < 60. Surface limited deposition of Cu and Sn occurred, respectively, at E = �0.37 V and E = �0.68 V in ammonia buffer containing EDTA. As already reported for binary sulfides SUPD layers on Ag(111) and on metal were obtained by keeping the potential at E = �0.68 V in Na2S ammonia solutions. Stripping analysis of the ternary sulfides yielded a large and well-defined peak centered at E = �0.22 V (Cu stripping), preceded by a broader peak at �0.43 V (Sn stripping). Charges involved in the stripping of both metals (Sn + Cu) and S allowed defining the effective layer-by-layer formation of a ternary compound with a slope of 42 μC cm�<sup>2</sup> , which is very close to the value obtained from the stripping of the binary CuS compound [19]. The chemical composition of Ag/S[(Cu/S)k/(Sn/S)j ]n deposits were analyzed by means of SEM, XPS and TOF-SIMS [37]; the ex-situ characterizations have highlighted the nominal stoichiometry is not respected leading to Sn/Cu ratio equal to 1/13 and 1/9 for j = 1, and 2, respectively. Independently from Sn/Cu ratio the morphology of the growing films, as revealed by AFM measurements, were found to be homogeneous and similar to the bare Ag(111). Conversely, the band gap values, carried out by diffuse reflectance spectroscopy, decreased with the increase in thickness and Cu/Sn ratio ranging from 2.12 for n = 60, j = 2–2.43 eV for n = 20, j = 1.

Cu NH ð Þ<sup>3</sup>

Cu NH ð Þ<sup>3</sup>

Cu NH ð Þ<sup>3</sup>

ing with the number of deposition cycles with a slope of 44 μC cm�<sup>2</sup>

high homogeneity and good quality of the thin film obtained.

different sequences of E-ALD cycles, that is Ag/S[(Cu/S)k/(Sn/S)j

[19]. The chemical composition of Ag/S[(Cu/S)k/(Sn/S)j

5.3. E-ALD of ternary MxN1-xS

98 Semiconductors - Growth and Characterization

thin films is separately discussed below.

2þ <sup>4</sup> þ e

> þ <sup>2</sup> þ e

2þ

� ! Cu NH ð Þ<sup>3</sup>

<sup>4</sup> þ Cuð Þ!0 2Cu NH ð Þ<sup>3</sup>

The amount of Cu deposited in a given number of cycles was determined by measuring the charge involved in the anodic stripping. The authors found the charges were linearly increas-

the valence states of copper and sulfur as Cu(I) and S(�II), respectively, although a possible fraction of S(�I) in the form of disulfide anion was not excluded. The experimental Cu/S ratio observed in the XPS characterization was later attributed to the covellite phase, where positive holes allow Cu ions to be stabilized in their monovalent state [32]. The morphological analysis by AFM was able to evidence the low roughness values of the deposits, thus confirming the

The interest in cadmium, zinc, tin, and copper sulfides has increased in the last 10 years due to the possibility to fabricate ternary semiconductor compounds, thus allowing fine control over the band gap energy of solar cell devices [32–34]. E-ALD method has been successfully employed to grow ternary materials such as CdxZn1-xS [35, 36], CuxZn1-xS [5] and CuxSnySz [33]. These semiconductors were prepared by sequential deposition of the corresponding binaries; for instance, alternate deposition of CuxS and SnxSy was carried out to form CuxSnySz. Because of the large variety of the possible (x:y:z) combinations research on E-ALD of multinary kesterite group thin films, although challenging, is quite promising for the development of non-linear electro-optic devices and photovoltaic cells. Depending on the adopted sequence profile only certain combinations were attainable, thus limiting the possible metalto-metal stoichiometries. The electrochemical behavior of CuxSnySz, CuxZn1-xS and CdxZn1-xS

Di Benedetto et al. [33] investigated the electrodeposition of CuxSnySz thin films exploiting

1 < n < 60. Surface limited deposition of Cu and Sn occurred, respectively, at E = �0.37 V and E = �0.68 V in ammonia buffer containing EDTA. As already reported for binary sulfides SUPD layers on Ag(111) and on metal were obtained by keeping the potential at E = �0.68 V in Na2S ammonia solutions. Stripping analysis of the ternary sulfides yielded a large and well-defined peak centered at E = �0.22 V (Cu stripping), preceded by a broader peak at �0.43 V (Sn stripping). Charges involved in the stripping of both metals (Sn + Cu) and S allowed defining the effective layer-by-layer formation of a ternary compound with a slope of 42 μC cm�<sup>2</sup>

which is very close to the value obtained from the stripping of the binary CuS compound

SEM, XPS and TOF-SIMS [37]; the ex-situ characterizations have highlighted the nominal stoichiometry is not respected leading to Sn/Cu ratio equal to 1/13 and 1/9 for j = 1, and 2, respectively. Independently from Sn/Cu ratio the morphology of the growing films, as revealed

þ

<sup>2</sup> þ 2NH<sup>3</sup> (20)

<sup>2</sup> (22)

. XPS results confirmed

]n with k = 1, j = 1, 2 and

]n deposits were analyzed by means of

,

� ! Cuð Þþ 0 2NH<sup>3</sup> (21)

þ

As for the Cu-Sn-S system the experimental conditions to grow ternary cadmium and zinc sulfides on Ag(111) with the E-ALD were by alternating the underpotential deposition of the corresponding binaries (CdS and ZnS) [36]. The potential chosen for the deposition of S and Zn was 0.65 V, whereas different series of experiments were carried out depositing Cd in correspondence to the first (0.5 V) or to the second UPD (0.65 V). The authors investigated the electrodeposition of CdxZn1-xS thin films exploiting different sequences of E-ALD cycles, that is, Ag/S[(Cd/S)k/(Zn/S)j ]n with k = 1, j = 1, 2, 5 and 1 < n < 20. The charge involved in the stripping increases linearly with the number of deposition cycles. Yet, the slope of the plot decreases while increasing the number of ZnS cycles per CdS cycles, that is, 89, 67 and 52 μC cm<sup>2</sup> for j = 1, 2 and 5, which reflect the higher percentage of Zn in the deposit. The chemical composition of CdxZn1-xS thin films, as analyzed by XPS have highlighted the nominal stoichiometry is not respected leading to Zn/Cu ratio equal to 1/3, 1/2 and 2 for j = 1, 2 and 5, respectively. However, regardless of the stoichiometry of the obtained ternary compound, the charge involved in cathodic stripping was equal to the charge involved in the anodic one, thus indicating the right 1:1 stoichiometric ratio [35].

The low contribution of Zn in CdxZn1-xS compounds was also found to occur in the Cu-Zn-S system. Ternary CuxZnyS compounds were obtained through the E-ALD method by alternating deposition of CuxS and ZnS layers. Innocenti et al. [8, 38] studied the electrochemical and compositional behavior of CdxZn1-xS thin films by applying the general sequence Ag/S/ [(Cu/S)j /(Zn/S)k)]n with j = 1, k = 1, 5, 9, and 1 < n < 60. As for other binary and ternary compounds UPD sulfur layers were obtained by keeping the electrode potential at E = 0.68 V for either on bare Ag(111) or on Ag(111) already covered by a metal layer. Deposition of Cu and Zn occurred in ammonia buffer at 0.37 and 0.85 V, respectively. Stripping analysis revealed the charge involved in (Cu + Zn) deposition followed a nonlinear increase with the number of cycles. Moreover, samples containing higher percentage of Zn had lower slope, thus indicating a lower extent of deposition in each cycle; a Cu/Zn ratio of about 6 was found for j = 1, k = 1, and n = 40. From the plot of charges against the number of deposition cycles (n) we expect a ZnS:CuS equimolar ratio for n = 13. Zn deficiency upon increasing the cycles is due to partial redissolution of zinc during the E-ALD of copper, thus causing a rearrangement in the adlayers. Independently from the Cu/Zn ratio, ex-situ investigations highlighted at least two prevailing morphologies, the first one homogeneously covering the Ag(111) surface, and the second one consisting of random network of nanowires of variable length [8].

## 6. E-ALD of binary MxSey semiconductors on Ag(111)

The E-ALD of metal selenides (MxSey) on monocrystalline surfaces usually starts with the deposition of a selenium monolayer by using selenite, Se(IV), as a precursor salt. Differently from sulfur deposition the oxidative UPD of Se is not allowed due to the low stability of selenite solutions. According to the work by Rajeshwar et al. [39] the electrochemistry of Se (IV) is quite complex. The direct reduction of Se(IV) to Se(0) is attributed to the formation of the electroinactive gray Se, which forms a stable deposit:

$$4H\_2\text{SeO}\_3 + 4H^+ + 4e^- \rightarrow \text{Se} + \text{3}\, H\_2\text{O}\tag{23}$$

For the growth of ZnSe, Sead was deposited from a 0.5 mM Se(IV) in a pH 9.2 ammonia buffer solution for 60 s at potential 0.95 V; then, the solution was replaced with the supporting electrolyte alone, and a potential E = 0.95 V was applied for 60 s to reduce the bulk Se(0). The Zn UPD on a Se-covered Ag(111) is obtained at E = 0.8 V for 30 s from 1 mM ZnSO4 in a pH 9.2 ammonia buffer solution. Once the deposit was formed, the amount of elements deposited in a given number of cycles was estimated from the charge involved in their stripping. The plots of the charges for Zn and Se stripping as a function of the number of cycles are linear, with a charge per cycle of 61 μC cm<sup>2</sup> for Zn and

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The electrochemical conditions necessary to form CdZnSe deposits by E-ALD on Ag(111) are described in the works by Loglio et al. [36, 50]. The first atomic layer on Ag(111) is obtained by SLR of selenium, as previously described. The UPD deposition on Se-covered Ag(111) from ammonia buffer of Cd would cause Zn redissolution. To shift CdUPD toward more negative potentials it is necessary to use a stronger complexing agent, that is, 0.1 M pyrophosphate in 0.01 M NaOH, and slightly shifting the Zn UPD toward more positive potentials using acetic buffer at pH 5.0. The experimental procedure to obtain CdxZn1 xSe with different x values consists of alternating ZnSe and CdSe ECALE cycles with different deposition sequences. The reductive UPD of Cd on a Se-covered Ag(111) substrate, from a 1 mM Cd(II) solution, was obtained by applying a potential E = 0.7 V for 60 s. In an analogous way, ZnSe was obtained by depositing Zn, from a 1 mM Zn(II) solutions, at E = 0.8 V. The stripping analysis of the compound obtained with 100 deposition (CdSe + ZnSe) cycles revealed a very high percentage

The most effective way to increase the amount of Zn in the deposit consists of depositing more ZnSe cycles per CdSe cycle. The x stoichiometric parameter is a linear function of the ZnSe/CdSe sequence; x = 0.5 was obtained with the ZnSe/CdSe sequence equal

The charge involved in stripping the cations equals the charge involved in stripping the anions, thus confirming the right 1:1 stoichiometric ratio calculated from XPS data. The charge per cycle changes while changing the ZnSe/CdSe sequence, a minimum at x = 0.5 is observed, with two symmetrical branches around the minimum. When the percentage of Zn in the compound is approximately equal to that of Cd, the different structures (Zincblend and Wurtzite) are present in a comparable amount, thus suggesting a structural disorder that could be responsi-

The E-ALD of metal tellurides, MxTey, on monocrystalline surfaces starts with the deposition of tellurium. As in the case of selenite, telluride solutions are not stable, so the oxidative UPD is

63 μC cm<sup>2</sup> for Se.

of Cd (74%).

ble for a reduced amount of deposition.

7. E-ALD of MxTey semiconductors on Ag(111)

to 5.

6.2. E-ALD of CdxZn1-XSe

Instead, the electroactive form of Se(0), usually referred as red Se, is obtained through a first reduction of Se(IV) to Se(�II) followed by a comproportionation reaction [40]:

$$\rm H\_2SeO\_3 + 6H^+ + 6e^- \rightarrow H\_2Se + 3\ H\_2O \tag{24}$$

$$\rm H\_2SeO\_3 + H\_2Se \to 3Se + 3 \, H\_2O \tag{25}$$

In ammonia buffer (pH 9.3), the presence of the electroactive red Se is evident for E < �0.95 V through the reduction Se(0) ! Se(�II), and around E = �0.8 V through the oxidation Se (0) ! Se(IV). Adlayers of Se on the electrode surface, Sead, are formed through a two-step procedure involving the deposition of an excess of Se(0) from Se(IV) solutions, followed by the reduction of bulk Se at sufficiently negative potential (E ≈ �0.95 V). The reduction must be performed in the absence of Se(IV) to avoid the comproportionation reaction with Se(�II) leading to a massive formation of Se(0). The STM investigation at potential more negative than the bulk reduction peak has shown two distinct structures. The layer at more positive potentials has a (√<sup>7</sup> � <sup>√</sup>7) R19.1� structure, with an associated charge of 65 <sup>μ</sup>C cm�<sup>2</sup> , whereas at more negative potentials it has a (2 √7 � 2√7) R19.1� structure, with an associated charge of 48 μC cm�<sup>2</sup> [41].

#### 6.1. E-ALD of CdSe and ZnSe

Cadmiun selenide is a promising material for application as thin film solar cells [42, 43], quantum dots [44, 45] and p-n junctions [46, 47]. ZnSe is an ideal candidate for optoelectronic devices, especially blue laser and blue emitters.

The E-ALD methodology has been successfully used for the growth of CdSe and ZnSe on Ag (111) [48, 49]. For the growth of CdSe, Sead was deposited from a 0.5 mM Se(IV) in a pH 8.5 ammonia buffer solution for 30s at a fixed potential E = �0.9 V; then the solution was replaced with the supporting electrolyte alone, and a potential E = �0.9 V was applied for 60s to reduce the bulk Se(0). The reductive underpotential deposition of Cd from a 1 mM CdSO4 in a pH 8.5 ammonia buffer solution on a Se-covered Ag(111) substrate, has 2 peaks. A well-defined UPD peak is observed at E = �0.41 V, whereas the beginning of a second UPD peak is observed at E = �0.69 V. The second UPD peak cannot be completely recorded, since it overlaps bulk Cd deposition and it can never be isolated from bulk Cd deposition. The optimal conditions for the adlayer formation of Cd on Se-covered Ag(111) are by keeping the electrode at E = �0.55 V for 30 s. The plots of the charges for Cd and Se stripping as a function of the number of ECALE cycles are linear, with an average charge per cycle of approximately 75 μC cm�<sup>2</sup> . The coincidence of the charges associated with each layer of Cd and Se gives a 1:1 stoichiometric ratio between the elements as expected for CdSe.

For the growth of ZnSe, Sead was deposited from a 0.5 mM Se(IV) in a pH 9.2 ammonia buffer solution for 60 s at potential 0.95 V; then, the solution was replaced with the supporting electrolyte alone, and a potential E = 0.95 V was applied for 60 s to reduce the bulk Se(0). The Zn UPD on a Se-covered Ag(111) is obtained at E = 0.8 V for 30 s from 1 mM ZnSO4 in a pH 9.2 ammonia buffer solution. Once the deposit was formed, the amount of elements deposited in a given number of cycles was estimated from the charge involved in their stripping. The plots of the charges for Zn and Se stripping as a function of the number of cycles are linear, with a charge per cycle of 61 μC cm<sup>2</sup> for Zn and 63 μC cm<sup>2</sup> for Se.

#### 6.2. E-ALD of CdxZn1-XSe

from sulfur deposition the oxidative UPD of Se is not allowed due to the low stability of selenite solutions. According to the work by Rajeshwar et al. [39] the electrochemistry of Se (IV) is quite complex. The direct reduction of Se(IV) to Se(0) is attributed to the formation of the

Instead, the electroactive form of Se(0), usually referred as red Se, is obtained through a first

In ammonia buffer (pH 9.3), the presence of the electroactive red Se is evident for E < �0.95 V through the reduction Se(0) ! Se(�II), and around E = �0.8 V through the oxidation Se (0) ! Se(IV). Adlayers of Se on the electrode surface, Sead, are formed through a two-step procedure involving the deposition of an excess of Se(0) from Se(IV) solutions, followed by the reduction of bulk Se at sufficiently negative potential (E ≈ �0.95 V). The reduction must be performed in the absence of Se(IV) to avoid the comproportionation reaction with Se(�II) leading to a massive formation of Se(0). The STM investigation at potential more negative than the bulk reduction peak has shown two distinct structures. The layer at more positive poten-

more negative potentials it has a (2 √7 � 2√7) R19.1� structure, with an associated charge of

Cadmiun selenide is a promising material for application as thin film solar cells [42, 43], quantum dots [44, 45] and p-n junctions [46, 47]. ZnSe is an ideal candidate for optoelectronic

The E-ALD methodology has been successfully used for the growth of CdSe and ZnSe on Ag (111) [48, 49]. For the growth of CdSe, Sead was deposited from a 0.5 mM Se(IV) in a pH 8.5 ammonia buffer solution for 30s at a fixed potential E = �0.9 V; then the solution was replaced with the supporting electrolyte alone, and a potential E = �0.9 V was applied for 60s to reduce the bulk Se(0). The reductive underpotential deposition of Cd from a 1 mM CdSO4 in a pH 8.5 ammonia buffer solution on a Se-covered Ag(111) substrate, has 2 peaks. A well-defined UPD peak is observed at E = �0.41 V, whereas the beginning of a second UPD peak is observed at E = �0.69 V. The second UPD peak cannot be completely recorded, since it overlaps bulk Cd deposition and it can never be isolated from bulk Cd deposition. The optimal conditions for the adlayer formation of Cd on Se-covered Ag(111) are by keeping the electrode at E = �0.55 V for 30 s. The plots of the charges for Cd and Se stripping as a function of the number of ECALE

� ! Se þ 3 H2O (23)

� ! H2Se þ 3 H2O (24)

, whereas at

. The coinci-

H2SeO<sup>3</sup> þ H2Se ! 3Se þ 3 H2O (25)

H2SeO<sup>3</sup> þ 4H<sup>þ</sup> þ 4e

reduction of Se(IV) to Se(�II) followed by a comproportionation reaction [40]:

tials has a (√<sup>7</sup> � <sup>√</sup>7) R19.1� structure, with an associated charge of 65 <sup>μ</sup>C cm�<sup>2</sup>

cycles are linear, with an average charge per cycle of approximately 75 μC cm�<sup>2</sup>

dence of the charges associated with each layer of Cd and Se gives a 1:1 stoichiometric ratio

H2SeO<sup>3</sup> þ 6H<sup>þ</sup> þ 6e

electroinactive gray Se, which forms a stable deposit:

100 Semiconductors - Growth and Characterization

48 μC cm�<sup>2</sup> [41].

6.1. E-ALD of CdSe and ZnSe

devices, especially blue laser and blue emitters.

between the elements as expected for CdSe.

The electrochemical conditions necessary to form CdZnSe deposits by E-ALD on Ag(111) are described in the works by Loglio et al. [36, 50]. The first atomic layer on Ag(111) is obtained by SLR of selenium, as previously described. The UPD deposition on Se-covered Ag(111) from ammonia buffer of Cd would cause Zn redissolution. To shift CdUPD toward more negative potentials it is necessary to use a stronger complexing agent, that is, 0.1 M pyrophosphate in 0.01 M NaOH, and slightly shifting the Zn UPD toward more positive potentials using acetic buffer at pH 5.0. The experimental procedure to obtain CdxZn1 xSe with different x values consists of alternating ZnSe and CdSe ECALE cycles with different deposition sequences. The reductive UPD of Cd on a Se-covered Ag(111) substrate, from a 1 mM Cd(II) solution, was obtained by applying a potential E = 0.7 V for 60 s. In an analogous way, ZnSe was obtained by depositing Zn, from a 1 mM Zn(II) solutions, at E = 0.8 V. The stripping analysis of the compound obtained with 100 deposition (CdSe + ZnSe) cycles revealed a very high percentage of Cd (74%).

The most effective way to increase the amount of Zn in the deposit consists of depositing more ZnSe cycles per CdSe cycle. The x stoichiometric parameter is a linear function of the ZnSe/CdSe sequence; x = 0.5 was obtained with the ZnSe/CdSe sequence equal to 5.

The charge involved in stripping the cations equals the charge involved in stripping the anions, thus confirming the right 1:1 stoichiometric ratio calculated from XPS data. The charge per cycle changes while changing the ZnSe/CdSe sequence, a minimum at x = 0.5 is observed, with two symmetrical branches around the minimum. When the percentage of Zn in the compound is approximately equal to that of Cd, the different structures (Zincblend and Wurtzite) are present in a comparable amount, thus suggesting a structural disorder that could be responsible for a reduced amount of deposition.

### 7. E-ALD of MxTey semiconductors on Ag(111)

The E-ALD of metal tellurides, MxTey, on monocrystalline surfaces starts with the deposition of tellurium. As in the case of selenite, telluride solutions are not stable, so the oxidative UPD is not allowed. Te(IV) reduction occurs following two possible schemes [51, 52]. A stable Te deposit is formed upon direct reduction of Te(IV)

$$2H\_2TeO\_2^+ + 3H^+ + 4e^- \to Te + 2\,H\_2O \tag{26}$$

8. E-ALD of CdSxSe1-x on Ag(111)

different sequences of E-ALD cycles, that is Ag/[(Cd/Se)j

E = 1.72 eV for x = 0 [54, 55].

neous phase.

9. Conclusions

Cadmium chalcogenides such as CdSe and CdS are excellent materials for the development of high efficient and low-cost photovoltaic devices. The small lattice mismatch between CdS and CdSe allows the formation of cadmium sulfoselenides CdSxSe1-x over a wide range of compositions (0 < x < 1), thus covering the visible solar spectrum from E = 2.44 eV for x = 1 to

E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals

The first ECALE study on ternary CdSxSe1-x compounds has been reported by Foresti et al. on Ag(111) substrates by alternating the underpotential deposition of CdSe and CdS layers [56]. To obtain CdSe, cadmium was deposited on Se-covered silver substrates from 1 mM Cd(II) solution in ammonia buffer (pH 9.6) by applying a potential E = 0.5 V for 60s. The formation of the first layer of Se on Ag(111) were obtained by a two-step procedure as described before [36]. After Cd deposition the cell was rinsed with ammonia buffer and the ECALE cycle was then completed by depositing CdS. The SUPD layer was obtained by applying a potential E = 0.68 V for 60s followed by washing with ammonia buffer. Finally, the reductive underpotential deposition of cadmium was attained on a S layer from 1 mM Cd(II) solutions by keeping the electrode at E = 0.65 V. The alternated deposition of CdSe and CdS can be repeated as many times as desired to obtained deposits of variable thickness and composition. The authors investigated

1 < n < 100. Independent from the sequence, CdSe deposition appeared to be favored with respect to CdS growth leading to the formation of sulfoselenides with x = 0.2 and 0.4 for j:k = 1:1 and 2:3, respectively. The charges involved in the stripping increase linearly with the number of deposition cycles, thus supporting a layer-by-layer growth. Regardless of the stoichiometry of the ternary compound obtained, the two charges are equal, thus confirming the right 1:1 stoichiometric ratio between Cd and (S + Se). Ex situ AFM measurements as a function of composition indicated the roughness decreased while increasing the S percentage, which determines a better deposit. Photoelectrochemical measurements on CdSxSe1-x ternary compounds revealed a monotonic band gap dependence with x, thus confirming the formation of a single homoge-

The E-ALD is a very inexpensive method for the production of high-crystalline thin film semiconductors. Exploiting Surface Limited Reactions (SLRs) on electrode surfaces allows the layer-by-layer deposition of different atomic elements. The E-ALD has been successfully used to grow ultra-thin films of metal chalcogenides on silver single crystals. The electrochemistry of binary compounds indicates that, with only the exception of the first layer, the charge associated with each layer of either metal or chalcogenide has the same average value. The layer in direct contact with the silver substrate can be regarded as an interface between the metal and the semiconductor electrodeposited on it. Semiconductors grown on Ag(111) by means of E-ALD are characterized by a high crystallinity of the materials and good photoconversion efficiencies. The ability to control the thickness of the deposited layers allows

/(Cd/S)k)]n with j:k = 1:1, j:k = 2:3 and

http://dx.doi.org/10.5772/intechopen.71014

103

or through a two-step process involving the reduction of Te(IV) to Te(�II) followed by a comproportionation reaction

$$2H\_2TeO\_2^+ + 5H^+ + 6e^- \to H\_2Te + 2H\_2O \tag{27}$$

$$2\text{ }\text{H}\_2\text{TeO}\_2^+ + 2\text{H}\_2\text{Te} \rightarrow 3\text{Te} + 2\text{ }\text{H}\_2\text{O} + \text{H}^+ \tag{28}$$

The TeO2 in ammonia buffer solutions (pH 8.5) on Ag(111) shows a large reduction peak at E = �0.4 V, which is only observed during the first scan from �0.1 to �0.9 V. Integration of the peak yields a charge of about 370 μC cm�<sup>2</sup> . Due to the high irreversibility of the system, the reoxidation of the underpotentially deposited Te is prevented by silver oxidation, so the Te UPD peak disappears in the successive CV scans. At about �1.1 V only bulk Te(0) reduction occurs. Te UPD layer starts the reduction at potentials more negative than �1.5 V because of higher bonding energy with the silver substrate.

So, the Tead can be obtained through UPD reduction before bulk deposition or in a two-step process: deposition at a potential of �0.6 V of small excess of bulk Te, followed by the reduction of bulk Te (but not Tead) at a potential of �1.4 V.

#### 7.1. E-ALD of CdTe

CdTe deposition on Ag(111) was obtained by E-ALD method [53], alternating UPD of tellurium and cadmium. The formation of the first layer of Te on Ag(111) is obtained by a two-step procedure instead of the direct deposition before the bulk reduction, previously described. This choice allows to having the same standard sequence for all cycles, which can be repeated as many times as desired.

The second step of the ECALE cycle is the underpotential deposition of the metallic element on the silver substrate covered by the non-metallic element. UPD of cadmium on Te-covered electrode occurs from 0.5 mM Cd(II) in ammonia buffer solution at a potential of �0.6 V, more negative than on bare Ag(111). Cadmium cyclic voltammograms, obtained from 0.5 mM Cd(II) in an ammonia buffer solution, do not exhibit narrow and sharp peaks as in the case of Tead. This finding could be attributed to a partial overlap with a second UPD peak, which in turn cannot be easily isolated from the concomitant beginning of bulk deposition. The second Tead layer cannot be obtained by the direct deposition before the bulk reduction, since underpotential deposition of Te occurs at a more positive potential than Cd stripping. Therefore, the two-step procedure described above has been adopted. The plots of the charges for Cd and Te stripping as a function of the number of cycles are linear, with an average charge per cycle equal to about 175 μC cm�<sup>2</sup> for Cd and 155 μC cm�<sup>2</sup> for Te. The ratio Cd/Te determined on the basis of electrochemical measurements is very close to the 1:1 stoichiometric ratio, which is indicative of a compound formation. The linear behavior suggests a layer-by-layer growth.

## 8. E-ALD of CdSxSe1-x on Ag(111)

not allowed. Te(IV) reduction occurs following two possible schemes [51, 52]. A stable Te

or through a two-step process involving the reduction of Te(IV) to Te(�II) followed by a

The TeO2 in ammonia buffer solutions (pH 8.5) on Ag(111) shows a large reduction peak at E = �0.4 V, which is only observed during the first scan from �0.1 to �0.9 V. Integration of the

reoxidation of the underpotentially deposited Te is prevented by silver oxidation, so the Te UPD peak disappears in the successive CV scans. At about �1.1 V only bulk Te(0) reduction occurs. Te UPD layer starts the reduction at potentials more negative than �1.5 V because of

So, the Tead can be obtained through UPD reduction before bulk deposition or in a two-step process: deposition at a potential of �0.6 V of small excess of bulk Te, followed by the

CdTe deposition on Ag(111) was obtained by E-ALD method [53], alternating UPD of tellurium and cadmium. The formation of the first layer of Te on Ag(111) is obtained by a two-step procedure instead of the direct deposition before the bulk reduction, previously described. This choice allows to having the same standard sequence for all cycles, which can be repeated

The second step of the ECALE cycle is the underpotential deposition of the metallic element on the silver substrate covered by the non-metallic element. UPD of cadmium on Te-covered electrode occurs from 0.5 mM Cd(II) in ammonia buffer solution at a potential of �0.6 V, more negative than on bare Ag(111). Cadmium cyclic voltammograms, obtained from 0.5 mM Cd(II) in an ammonia buffer solution, do not exhibit narrow and sharp peaks as in the case of Tead. This finding could be attributed to a partial overlap with a second UPD peak, which in turn cannot be easily isolated from the concomitant beginning of bulk deposition. The second Tead layer cannot be obtained by the direct deposition before the bulk reduction, since underpotential deposition of Te occurs at a more positive potential than Cd stripping. Therefore, the two-step procedure described above has been adopted. The plots of the charges for Cd and Te stripping as a function of the number of cycles are linear, with an average charge per cycle equal to about 175 μC cm�<sup>2</sup> for Cd and 155 μC cm�<sup>2</sup> for Te. The ratio Cd/Te determined on the basis of electrochemical measurements is very close to the 1:1 stoichiometric ratio, which is indicative of a compound

� ! Te þ 2 H2O (26)

� ! H2Te þ 2 H2O (27)

. Due to the high irreversibility of the system, the

<sup>2</sup> þ 2H2Te ! 3Te þ 2 H2O þ H<sup>þ</sup> (28)

<sup>2</sup> þ 3H<sup>þ</sup> þ 4e

<sup>2</sup> þ 5H<sup>þ</sup> þ 6e

deposit is formed upon direct reduction of Te(IV)

peak yields a charge of about 370 μC cm�<sup>2</sup>

higher bonding energy with the silver substrate.

reduction of bulk Te (but not Tead) at a potential of �1.4 V.

formation. The linear behavior suggests a layer-by-layer growth.

comproportionation reaction

102 Semiconductors - Growth and Characterization

7.1. E-ALD of CdTe

as many times as desired.

H2TeO<sup>þ</sup>

H2TeO<sup>þ</sup>

H2TeO<sup>þ</sup>

Cadmium chalcogenides such as CdSe and CdS are excellent materials for the development of high efficient and low-cost photovoltaic devices. The small lattice mismatch between CdS and CdSe allows the formation of cadmium sulfoselenides CdSxSe1-x over a wide range of compositions (0 < x < 1), thus covering the visible solar spectrum from E = 2.44 eV for x = 1 to E = 1.72 eV for x = 0 [54, 55].

The first ECALE study on ternary CdSxSe1-x compounds has been reported by Foresti et al. on Ag(111) substrates by alternating the underpotential deposition of CdSe and CdS layers [56]. To obtain CdSe, cadmium was deposited on Se-covered silver substrates from 1 mM Cd(II) solution in ammonia buffer (pH 9.6) by applying a potential E = 0.5 V for 60s. The formation of the first layer of Se on Ag(111) were obtained by a two-step procedure as described before [36]. After Cd deposition the cell was rinsed with ammonia buffer and the ECALE cycle was then completed by depositing CdS. The SUPD layer was obtained by applying a potential E = 0.68 V for 60s followed by washing with ammonia buffer. Finally, the reductive underpotential deposition of cadmium was attained on a S layer from 1 mM Cd(II) solutions by keeping the electrode at E = 0.65 V. The alternated deposition of CdSe and CdS can be repeated as many times as desired to obtained deposits of variable thickness and composition. The authors investigated different sequences of E-ALD cycles, that is Ag/[(Cd/Se)j /(Cd/S)k)]n with j:k = 1:1, j:k = 2:3 and 1 < n < 100. Independent from the sequence, CdSe deposition appeared to be favored with respect to CdS growth leading to the formation of sulfoselenides with x = 0.2 and 0.4 for j:k = 1:1 and 2:3, respectively. The charges involved in the stripping increase linearly with the number of deposition cycles, thus supporting a layer-by-layer growth. Regardless of the stoichiometry of the ternary compound obtained, the two charges are equal, thus confirming the right 1:1 stoichiometric ratio between Cd and (S + Se). Ex situ AFM measurements as a function of composition indicated the roughness decreased while increasing the S percentage, which determines a better deposit. Photoelectrochemical measurements on CdSxSe1-x ternary compounds revealed a monotonic band gap dependence with x, thus confirming the formation of a single homogeneous phase.

#### 9. Conclusions

The E-ALD is a very inexpensive method for the production of high-crystalline thin film semiconductors. Exploiting Surface Limited Reactions (SLRs) on electrode surfaces allows the layer-by-layer deposition of different atomic elements. The E-ALD has been successfully used to grow ultra-thin films of metal chalcogenides on silver single crystals. The electrochemistry of binary compounds indicates that, with only the exception of the first layer, the charge associated with each layer of either metal or chalcogenide has the same average value. The layer in direct contact with the silver substrate can be regarded as an interface between the metal and the semiconductor electrodeposited on it. Semiconductors grown on Ag(111) by means of E-ALD are characterized by a high crystallinity of the materials and good photoconversion efficiencies. The ability to control the thickness of the deposited layers allows

[5] Innocenti M, Becucci L, Bencistà I, Carretti E, Cinotti S, Dei L, Di Benedetto F, Lavacchi A, Marinelli F, Salvietti E, Vizza F, Foresti ML. Electrochemical growth of Cu-Zn sulfides.

E-ALD: Tailoring the Optoeletronic Properties of Metal Chalcogenides on Ag Single Crystals

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Figure 4. Experimental band gap values (markers) of ternary compounds as a function of (A) thickness at fixed composition and (B) Cu ratio at fixed number of S layers, n = 121. The dotted lines are the linear fit to the data.

fine tuning the semiconductor transition energy by varying the number of deposition cycles (Figure 4A). Moreover, careful design of the growth sequence of ternary compounds direct the formation of advanced photovoltaic materials having compositionally controlled band gaps (Figure 4B).

#### Author details

Emanuele Salvietti<sup>1</sup> , Andrea Giaccherini<sup>1</sup> , Filippo Gambinossi<sup>1</sup> , Maria Luisa Foresti<sup>1</sup> , Maurizio Passaponti<sup>1</sup> , Francesco Di Benedetto<sup>2</sup> and Massimo Innocenti<sup>1</sup> \*

\*Address all correspondence to: m.innocenti@unifi.it


#### References


fine tuning the semiconductor transition energy by varying the number of deposition cycles (Figure 4A). Moreover, careful design of the growth sequence of ternary compounds direct the formation of advanced photovoltaic materials having compositionally controlled band gaps

Figure 4. Experimental band gap values (markers) of ternary compounds as a function of (A) thickness at fixed compo-

sition and (B) Cu ratio at fixed number of S layers, n = 121. The dotted lines are the linear fit to the data.

, Filippo Gambinossi<sup>1</sup>

, Francesco Di Benedetto<sup>2</sup> and Massimo Innocenti<sup>1</sup>

[1] Gregory BW, Stickney JL. Electrochemical atomic layer epitaxy (ECALE). Journal of

[2] Innocenti M, Pezzatini G, Forni F, Foresti ML. CdS and ZnS deposition on Ag (111) by electrochemical atomic layer epitaxy. Journal of the Electrochemical Society. 2001;148:C357-C362 [3] Innocenti M, Forni F, Pezzatini G, Raiteri R, Loglio F, Foresti ML. Electrochemical behavior of As on silver single crystals and experimental conditions for InAs growth by

[4] Loglio F, Innocenti M, Jarek A, Caporali S, Pasquini I, Foresti ML. Nickel sulfur thin films deposited by ECALE: Electrochemical, XPS and AFM characterization. Journal of Elec-

, Maria Luisa Foresti<sup>1</sup>

\*

,

(Figure 4B).

Author details

104 Semiconductors - Growth and Characterization

Emanuele Salvietti<sup>1</sup>

References

Maurizio Passaponti<sup>1</sup>

, Andrea Giaccherini<sup>1</sup>

1 Department of Chemistry, University of Florence, Sesto Fiorentino, Italy

2 Department of Earth Sciences, University of Florence, Florence, Italy

\*Address all correspondence to: m.innocenti@unifi.it

Electroanalytical Chemistry. 1991;300:543-561

troanalytical Chemistry. 2010;638(10):15-20

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[35] Innocenti M, Cattarin S, Loglio F, Cecconi T, Seravalli G, Foresti ML. Ternary cadmium and zinc sulfides: Composition, morphology and photoelectrochemistry. Electrochimica

[36] Loglio F, Innocenti M, Pezzatini G, Foresti ML. Ternary cadmium and zinc sulfides and selenides: Electrodeposition by ECALE and electrochemical characterization. Journal of

[37] Caporali S, Tolstogouzov A, Teodoro OMND, Innocenti M, Di Benedetto F, Cinotti S, Picca RA, Sportelli MC, Cioffi N. Sn-deficiency in the electrodeposited ternary CuxSnySz

[38] Di Benedetto F, Cinotti S, D'Acapito F, Vizza F, Foresti ML, Guerri A, Lavacchi A, Montegrossi G, Romanelli M, Cioffi N, Innocenti M. Electrodeposited semiconductors at room temperature: An X-ray absorption spectroscopy study of cu-, Zn-, S-bearing thin

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**Chapter 6**

**Provisional chapter**

 **and** 

**Pulsed Electrochemical Deposition of CuInSe2 and**

 **Semiconductor Thin Films**

materials for thin films solar cells due to their direct bandgap and large absorption coefficient. The highly efficient CIGS devices are often fabricated using expensive vacuum based technologies; however, recently electrodeposition has been demonstrated to produce CIGS devices with high efficiencies and it is easily amenable for large area films of high quality with effective material use and high deposition rate. In this context, this chapter discusses the recent developments in CIS and CIGS technologies using electrodeposition. In addition, the fundamental features of electrodeposition such as direct current, pulse and pulse-reverse plating and their application in the fabrication of CIS and CIGS films are discussed. In conclusion, the chapter summarizes the utilization of pulse electrodeposition for fabrication of CIS and CIGS films while making a recommendation

**Pulsed Electrochemical Deposition of CuInSe<sup>2</sup>**

DOI: 10.5772/intechopen.71857

© 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,

© 2018 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.

(CIGS) semiconductors are the most studied absorber

and reproduction in any medium, provided the original work is properly cited.

The present day's global energy production is mostly accomplished from the fossil fuels; however, the inherent problems associated with the use of fossil fuels such as their limited availability and the environmental issues force the mankind to look for alternative solutions for future energy supply. The need to develop renewable energy resources has come to the forefront of discussion. Photovoltaics (PV) is an emerging field and one of the choices for major means of future energy-harvesting. The efficiencies of PV conversion depend on the properties

**Cu(In,Ga)Se2 Semiconductor Thin Films**

Sreekanth Mandati, Bulusu V. Sarada, Suhash R. Dey and

for exploring the group's unique pulse electroplating method.

gallium selenide, photoelectrochemical cells, solar cells

**Keywords:** pulse electrodeposition, semiconductors, thin films, copper indium

Sreekanth Mandati, Bulusu V. Sarada, Suhash R. Dey and Shrikant V. Joshi

**Cu(In,Ga)Se<sup>2</sup>**

Shrikant V. Joshi

**Abstract**

CuInSe2

**1. Introduction**

http://dx.doi.org/10.5772/intechopen.71857

Additional information is available at the end of the chapter

(CIS) and Cu(In,Ga)Se2

Additional information is available at the end of the chapter


#### **Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films Pulsed Electrochemical Deposition of CuInSe<sup>2</sup> and Cu(In,Ga)Se<sup>2</sup> Semiconductor Thin Films**

DOI: 10.5772/intechopen.71857

Sreekanth Mandati, Bulusu V. Sarada, Suhash R. Dey and Shrikant V. Joshi Sreekanth Mandati, Bulusu V. Sarada, Suhash R. Dey and Shrikant V. Joshi

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/intechopen.71857

#### **Abstract**

[44] Kim S, Fisher B, Eisler H-J, Bawendi M. Type-II quantum dots: CdTe/CdSe(core/shell) and CdSe/ZnTe(core/shell) heterostructures. Journal of the American Chemical Society. 2003;

[45] Dabbousi BO, Rodriguez-Viejo J, Mikulec FV, Heine JR, Mattoussi H, Ober R, Jensen KF, Bawendi MG. (CdSe)ZnS coreshell quantum dots: Synthesis and characterization of a size series of highly luminescent Nanocrystallites. The Journal of Physical Chemistry B.

[46] Gur I, Fromer NA, Geier ML, Alivisatos AP. Air-stable all-inorganic nanocrystal solar

[47] Huang Y, Duan X, Lieber CM. Nanowires for integrated multicolor nanophotonics.

[48] Loglio F, Innocenti M, D'acapito F, Felici R, Pezzatini G, Salvietti E, Foresti ML. Cadmium selenide electrodeposited by ECALE: Electrochemical characterization and preliminary

[49] Pezzatini G, Caporali S, Innocenti M, Foresti ML. Formation of ZnSe on Ag(111) by electrochemical atomic layer epitaxy. Journal of Electroanalytical Chemistry. 1999;475(2):

[50] Loglio F, Telford AM, Salvietti E, Innocenti M, Pezzatini G, Cammelli S, D'Acapito F, Felici R, Pozzi A, Foresti ML. Ternary CdxZn1xSe deposited on Ag (111) by ECALE: Electrochemical and EXAFS characterization. Electrochimica Acta. 2008;53:6978-6987 [51] Mori E, Baker CK, Reynolds JR, Rajeshwar K. Aqueous electrochemistry of tellurium at glassy carbon and gold: A combined voltammetry-oscillating quartz crystal microgravimetry study.

[52] Traore M, Modolo R, Vittori O. Electrochemical behaviour of tellurium and silver telluride at rotating glassy carbon electrode. Electrochimica Acta. 1988;33:991-996

[53] Forni F, Innocenti M, Pezzatini G, Foresti ML. Electrochemical aspects of CdTe growth on

[54] Liang YQ, Zhai L, Zhao XS, Xu DS. Band-gap engineering of semiconductor nanowires through composition modulation. The Journal of Physical Chemistry B. 2005;109(15):

[55] Takahashi T, Nichols P, Takei K, Ford AC, Jamshidi A, Wu MC, Ning CZ, Javey A. Contact printing of compositionally graded CdSxSe1-x nanowire parallel arrays for tun-

[56] Foresti ML, Milani S, Loglio F, Innocenti M, Pezzatini G, Cattarin S. Ternary CdSxSe(1-x) deposited on Ag(111) by ECALE: Synthesis and characterization. Langmuir. 2005;21(15):

the face (111) of silver by ECALE. Electrochimica Acta. 2000;45:3225-3231

results by EXAFS. Journal of Electroanalytical Chemistry. 2005;575(1):161-167

cells processed from solution. Science. 2005;310(5747):462-465

Journal of Electroanalytical Chemistry. 1988;253:441-451

able photodetectors. Nanotechnology. 2012;23(4):045201

125(38):11466-11467

108 Semiconductors - Growth and Characterization

1997;101(46):9463-9475

Small. 2005;1(1):142-147

164-170

7120-7123

6900-6907

CuInSe2 (CIS) and Cu(In,Ga)Se2 (CIGS) semiconductors are the most studied absorber materials for thin films solar cells due to their direct bandgap and large absorption coefficient. The highly efficient CIGS devices are often fabricated using expensive vacuum based technologies; however, recently electrodeposition has been demonstrated to produce CIGS devices with high efficiencies and it is easily amenable for large area films of high quality with effective material use and high deposition rate. In this context, this chapter discusses the recent developments in CIS and CIGS technologies using electrodeposition. In addition, the fundamental features of electrodeposition such as direct current, pulse and pulse-reverse plating and their application in the fabrication of CIS and CIGS films are discussed. In conclusion, the chapter summarizes the utilization of pulse electrodeposition for fabrication of CIS and CIGS films while making a recommendation for exploring the group's unique pulse electroplating method.

**Keywords:** pulse electrodeposition, semiconductors, thin films, copper indium gallium selenide, photoelectrochemical cells, solar cells

#### **1. Introduction**

The present day's global energy production is mostly accomplished from the fossil fuels; however, the inherent problems associated with the use of fossil fuels such as their limited availability and the environmental issues force the mankind to look for alternative solutions for future energy supply. The need to develop renewable energy resources has come to the forefront of discussion. Photovoltaics (PV) is an emerging field and one of the choices for major means of future energy-harvesting. The efficiencies of PV conversion depend on the properties

© 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. © 2018 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.

of absorber layer, which is the heart of the solar cell. Silicon has been the foundation of the photovoltaics industry due to its well-known properties, high abundance and well-studied technological aspects of the growth. However, it requires expensive manufacturing technologies such as growing and sawing of ingots. Hence, current trend in photovoltaics requires the development of high performance inexpensive solar absorber materials that can serve in the long term as viable alternatives to the single crystal silicon technology. Among various technologies, CuInSe<sup>2</sup> based solar cells are the most studied and leading candidates to realize commercialization with efficiencies close to 23% [1].

Higher Ga content of 40% has a detrimental effect on the device performance, because it negatively impacts the transport properties of the CIGS absorber film. The current, high-efficiency devices are prepared with bandgaps in the range 1.20–1.25 eV, which corresponds to a Ga/

Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films

http://dx.doi.org/10.5772/intechopen.71857

111

The fabrication methods often are the big difference in industrializing a technology based on a materials system. To determine the most promising technique for the commercial manufacturing of modules, the foremost criteria are: (i) low cost, (ii) scalability, (iii) reproducibility and (iv) manufacturability. Several vacuum techniques including co-evaporation, sputtering, molecular beam epitaxy, pulsed laser deposition, etc., have been investigated for the formation of CIGS thin-films. The most successful technique for deposition of CIGS absorber layers for highest efficiency small area cells is the co-evaporation of elements from multiple sources where Se is offered in excess during the process [19]. Although thin-film CIGS solar cells with power conversion efficiencies over 23% have been demonstrated, the vacuum-based processes used therein pose cost and technological barriers in the production of PV modules. With this requirement in mind, development of low-cost methods to fabricate CIGS has become an intensively pursued goal and a variety of solution-based approaches have been demonstrated. Non-vacuum approaches, include electro or electroless-deposition, chemical bath deposition, particulate processes and coating involving molecular precursors are being explored. The efficiency gap between vacuum and non-vacuum deposited CIGS has been reduced in past years and processes from the above categories have now reported cells with efficiencies of 15–17%, thereby, showing promise for commercialization. Among these methods, electrochemical deposition is the most extensively explored technique for the deposition of CIGS absorber layers and has witnessed high efficiency devices [20–23], and will be discussed in detail in the

The process of electro-reduction of precursor ions onto an electrode, substrate of interest, from an appropriate electrolyte by the utilization of electric current or potential between two electrodes is termed as electrodeposition. The properties of the deposit depends on several parameters including ionic concentrations, the electrodes used, the pH of the electrolyte, the temperature, the type of substrate, the stirring rate, the deposition voltage and the time of deposition. Electrodeposition is a promising approach to fabricate the absorber layers in thin

• It is a low cost solution based approach and avoids expensive vacuum technology

• Use of low cost starting materials, e.g., low purity salts and solvents can be purified with

• Deposition of films on a variety of shapes and forms of wires, tapes, coils, and cylinders

(In + Ga) ratio between 25 and 30% [14–18].

following section.

**3. Electrodeposition**

films solar cells due to following advantages:

• Uses low temperature for deposition

• High deposition rate on larger areas with lateral uniformity

application of small voltage before the electrodeposition

## **2. Copper indium selenide (CIS) and copper indium gallium selenide (CIGS)**

CIS is a ternary semiconductor belonging to the I-III-VI class, crystallizes to a chalcopyrite structure, possesses a direct bandgap of 1.04 eV and an absorption coefficient of ≈ 10<sup>5</sup> cm−1 [2, 3]. The CIS-based solar cells exhibit excellent chemical stability, stability with time and doping versatility. Absorber layer is the key element of solar cells, which is produced mainly from the p-type semiconductor in thin films solar cells. The electrical properties of Cu ternary semiconductors are determined by native defects [4]. There are three possible electrically actives defects namely, vacancies, interstitials and antisite defects [5–7]. It is these defects which determine the nature of the conductivity of CIS films whether n type or p type. Intrinsic copper vacancies (VCu) and copper on indium antisite defects (CuIn) are the electrically active defects for a typical p-type CIS film. On the other hand, intrinsic selenium vacancies (VSe) and indium on copper antisite defects (InCu) makes the CIS n-type. The material with Cu rich composition is not preferred mainly because of the formation of copper selenide (Cu<sup>2</sup> Se). Cu<sup>2</sup> Se being highly conductive, shorts out the junction. Adding more Indium than copper reduces the formation of Cu2 Se but it causes other defects like VCu and InCu which are compensating in nature [8]. Hence, the copper to indium ratio (Cu/In) is always maintained around unity. Samples with p-type conductivity are grown if the material is Cu-poor and is annealed under high Se vapor pressure, whereas Cu-rich material with Se deficiency tends to be n-type. CIS films when suitably manufactured tend to be p-type because of the low energy of formation of copper vacancies which give the material its conductivity [4, 7, 9]. CIS solar cells yielded relatively lower open circuit potentials due to its small bandgap. This limitation is overcome by adding controlled amounts of gallium to replace indium in the CIS structure [10]. The band gap of CuIn1-xGa<sup>x</sup> Se2 varies according to the equation [11].

$$E\_{\rm g} = 1.011 + 0.664 \text{x} + 0.249 \text{x} (1 - \text{x}) \tag{1}$$

Depending on the [Ga]/[In + Ga] ratio, the bandgap of CIGS can be varied continuously between 1.02 and 1.68 eV [10, 12]. The addition of about 30% Ga in CIS increases the bandgap to 1.2 eV which has a closer match with the AM 1.5 solar spectrum [12, 13]. Addition of Ga not only increases the band gap but also has other beneficial effects such as improving the adhesion of the film to the Mo substrate, increased carrier concentration, etc. Though it is beneficial to add Ga to improve the properties of CIS, there is a limit to which it serves favorable. Higher Ga content of 40% has a detrimental effect on the device performance, because it negatively impacts the transport properties of the CIGS absorber film. The current, high-efficiency devices are prepared with bandgaps in the range 1.20–1.25 eV, which corresponds to a Ga/ (In + Ga) ratio between 25 and 30% [14–18].

The fabrication methods often are the big difference in industrializing a technology based on a materials system. To determine the most promising technique for the commercial manufacturing of modules, the foremost criteria are: (i) low cost, (ii) scalability, (iii) reproducibility and (iv) manufacturability. Several vacuum techniques including co-evaporation, sputtering, molecular beam epitaxy, pulsed laser deposition, etc., have been investigated for the formation of CIGS thin-films. The most successful technique for deposition of CIGS absorber layers for highest efficiency small area cells is the co-evaporation of elements from multiple sources where Se is offered in excess during the process [19]. Although thin-film CIGS solar cells with power conversion efficiencies over 23% have been demonstrated, the vacuum-based processes used therein pose cost and technological barriers in the production of PV modules. With this requirement in mind, development of low-cost methods to fabricate CIGS has become an intensively pursued goal and a variety of solution-based approaches have been demonstrated. Non-vacuum approaches, include electro or electroless-deposition, chemical bath deposition, particulate processes and coating involving molecular precursors are being explored. The efficiency gap between vacuum and non-vacuum deposited CIGS has been reduced in past years and processes from the above categories have now reported cells with efficiencies of 15–17%, thereby, showing promise for commercialization. Among these methods, electrochemical deposition is the most extensively explored technique for the deposition of CIGS absorber layers and has witnessed high efficiency devices [20–23], and will be discussed in detail in the following section.

## **3. Electrodeposition**

of absorber layer, which is the heart of the solar cell. Silicon has been the foundation of the photovoltaics industry due to its well-known properties, high abundance and well-studied technological aspects of the growth. However, it requires expensive manufacturing technologies such as growing and sawing of ingots. Hence, current trend in photovoltaics requires the development of high performance inexpensive solar absorber materials that can serve in the long term as viable alternatives to the single crystal silicon technology. Among various

CIS is a ternary semiconductor belonging to the I-III-VI class, crystallizes to a chalcopyrite structure, possesses a direct bandgap of 1.04 eV and an absorption coefficient of ≈ 10<sup>5</sup> cm−1 [2, 3]. The CIS-based solar cells exhibit excellent chemical stability, stability with time and doping versatility. Absorber layer is the key element of solar cells, which is produced mainly from the p-type semiconductor in thin films solar cells. The electrical properties of Cu ternary semiconductors are determined by native defects [4]. There are three possible electrically actives defects namely, vacancies, interstitials and antisite defects [5–7]. It is these defects which determine the nature of the conductivity of CIS films whether n type or p type. Intrinsic copper vacancies (VCu) and copper on indium antisite defects (CuIn) are the electrically active defects for a typical p-type CIS film. On the other hand, intrinsic selenium vacancies (VSe) and indium on copper antisite defects (InCu) makes the CIS n-type. The material with Cu rich com-

**2. Copper indium selenide (CIS) and copper indium gallium** 

position is not preferred mainly because of the formation of copper selenide (Cu<sup>2</sup>

varies according to the equation [11].

being highly conductive, shorts out the junction. Adding more Indium than copper reduces

in nature [8]. Hence, the copper to indium ratio (Cu/In) is always maintained around unity. Samples with p-type conductivity are grown if the material is Cu-poor and is annealed under high Se vapor pressure, whereas Cu-rich material with Se deficiency tends to be n-type. CIS films when suitably manufactured tend to be p-type because of the low energy of formation of copper vacancies which give the material its conductivity [4, 7, 9]. CIS solar cells yielded relatively lower open circuit potentials due to its small bandgap. This limitation is overcome by adding controlled amounts of gallium to replace indium in the CIS structure [10]. The band

*Eg* = 1.011 + 0.664*x* + 0.249*x*(1 − *x*) (1)

Depending on the [Ga]/[In + Ga] ratio, the bandgap of CIGS can be varied continuously between 1.02 and 1.68 eV [10, 12]. The addition of about 30% Ga in CIS increases the bandgap to 1.2 eV which has a closer match with the AM 1.5 solar spectrum [12, 13]. Addition of Ga not only increases the band gap but also has other beneficial effects such as improving the adhesion of the film to the Mo substrate, increased carrier concentration, etc. Though it is beneficial to add Ga to improve the properties of CIS, there is a limit to which it serves favorable.

Se but it causes other defects like VCu and InCu which are compensating

based solar cells are the most studied and leading candidates to realize

Se). Cu<sup>2</sup>

Se

technologies, CuInSe<sup>2</sup>

110 Semiconductors - Growth and Characterization

**selenide (CIGS)**

the formation of Cu2

gap of CuIn1-xGa<sup>x</sup>

Se2

commercialization with efficiencies close to 23% [1].

The process of electro-reduction of precursor ions onto an electrode, substrate of interest, from an appropriate electrolyte by the utilization of electric current or potential between two electrodes is termed as electrodeposition. The properties of the deposit depends on several parameters including ionic concentrations, the electrodes used, the pH of the electrolyte, the temperature, the type of substrate, the stirring rate, the deposition voltage and the time of deposition. Electrodeposition is a promising approach to fabricate the absorber layers in thin films solar cells due to following advantages:


**3.2. Pulse electrodeposition**

**3.3. Pulse reverse electrodeposition**

In pulse electrodeposition (PED), current/potential is applied in the form of modulated waves as shown in **Figure 2**. Compared to DC, PED offers additional process control variables like pulse on-time (ton) and off-time (toff). Often, the variation in ton and toff is expressed using a

The precise variation in duty cycle provides the control over electrochemical processes by affecting the diffusion layer, grain size and nucleation. Usually in electroplating a negatively charged layer is formed around the cathode which gets charged to a known thickness and prevents the ions from the bulk. In DC, this charged double layer prevents the passage of ions toward the cathode thereby affecting the features of the deposit and also causes the inhomogeneity in the concentration of ions in the electrolyte. However, in PED, the charged diffusion layer gets discharged and helps easier passage of the ions onto the cathode as the output is periodically turned off. Migration of ions to depleted areas in the bath during off-time makes the even distribution of ions for their easy availability during on-time. The presence of off time aids in the relaxation and the rearrangement of deposited atoms leading to the possibility of new nucleation sites during the subsequent deposition. This, in turn, not only improves the uniformity of deposition but also reduces the porosity and roughness of the deposit (see **Figure 2**). In addition, the entrapped hydrogen and impurities during the deposition diffuse out during the relaxation time. The additional process variables in PED ease the process of optimization for complex stoichiometric ternary/quaternary systems as the control over individual elemental composition is far better with the variation in duty cycle compared to DC technique.

Pulse reverse electrodeposition contains short anodic pulses alongside the cathodic as seen in PED (See **Figure 3**). This small anodic pulse is advantageous since it contributes to the electro-oxidation of the top layer from the deposited film thereby aids in smoothening of the deposit, removing the impurities as well as the entrapped hydrogen. The anodic pulse plays a major role in the deposition of systems like CIS and CIGS, wherein an undesired

**Figure 2.** Schematic of pulse electrodeposition and expected growth process of the deposit.

*t on* + *t off*

Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films

× 100 (2)

http://dx.doi.org/10.5772/intechopen.71857

113

common parameter known as duty cycle, which is defined by the equation:

*Duty Cycle* (%) <sup>=</sup> *<sup>t</sup>* \_\_\_\_\_ *on*

• Bandgap engineering either by varying the contents of individual elements in step-by-step approach or by using combinatorial approach

Owing to these advantages, electrodeposition satisfies all the necessary criteria for the research and development of PV solar cells. In addition, electrodeposition involves two fundamental approaches such as direct current and pulse plating wherein pulse plating further comprises pulse and pulse-reverse features. These features are unique in themselves that they provide additional process control variables and make electrodeposition an attractive tool for the deposition of semiconductors [24].

#### **3.1. Direct current (DC) Electrodeposition**

Direct current electrodeposition is the conventional method wherein a constant current or potential is applied continuously during the deposition to coat the desired materials. **Figure 1** shows the schematic of the direct current applied and the typical growth process corresponding to it. DC technique effectively has two variables, namely, applied potential/current and time of deposition while the precursor concentration and electrolyte pH are the common variables. The morphology, composition and thickness of the deposit can be altered by varying these parameters. In DC, the continuous use of constant potential/current leads to deposition of films without any relaxation leading to the growth of existing nuclei instead of generating new nucleation sites thereby resulting in a rough and porous deposit. In addition, hydrogen evolution reaction (HER) competes along with the deposition of desired materials as often aqueous electrolytes are used. The HER not only influences the current efficiency of the technique but also the deposited film properties. Use of additives in conventional DC deposition has improved the morphology of the deposited films. Despite having several disadvantages, DC electrodeposition is still a leading technique for the production of single element deposits and binary alloys.

**Figure 1.** Schematic of direct current electrodeposition and expected growth process of the deposit.

#### **3.2. Pulse electrodeposition**

• Effective material use as high as 98% with minimum waste generation

trations of dopant precursor in electrodeposition bath

approach or by using combinatorial approach

the deposition of semiconductors [24].

112 Semiconductors - Growth and Characterization

**3.1. Direct current (DC) Electrodeposition**

• Control over the composition of individual elements by varying deposition parameters

• Extrinsic doping of semiconductors with appropriate elements by utilizing small concen-

• Bandgap engineering either by varying the contents of individual elements in step-by-step

Owing to these advantages, electrodeposition satisfies all the necessary criteria for the research and development of PV solar cells. In addition, electrodeposition involves two fundamental approaches such as direct current and pulse plating wherein pulse plating further comprises pulse and pulse-reverse features. These features are unique in themselves that they provide additional process control variables and make electrodeposition an attractive tool for

Direct current electrodeposition is the conventional method wherein a constant current or potential is applied continuously during the deposition to coat the desired materials. **Figure 1** shows the schematic of the direct current applied and the typical growth process corresponding to it. DC technique effectively has two variables, namely, applied potential/current and time of deposition while the precursor concentration and electrolyte pH are the common variables. The morphology, composition and thickness of the deposit can be altered by varying these parameters. In DC, the continuous use of constant potential/current leads to deposition of films without any relaxation leading to the growth of existing nuclei instead of generating new nucleation sites thereby resulting in a rough and porous deposit. In addition, hydrogen evolution reaction (HER) competes along with the deposition of desired materials as often aqueous electrolytes are used. The HER not only influences the current efficiency of the technique but also the deposited film properties. Use of additives in conventional DC deposition has improved the morphology of the deposited films. Despite having several disadvantages, DC electrodeposition is still a leading technique for the production of single element deposits and binary alloys.

**Figure 1.** Schematic of direct current electrodeposition and expected growth process of the deposit.

In pulse electrodeposition (PED), current/potential is applied in the form of modulated waves as shown in **Figure 2**. Compared to DC, PED offers additional process control variables like pulse on-time (ton) and off-time (toff). Often, the variation in ton and toff is expressed using a common parameter known as duty cycle, which is defined by the equation:

$$\text{Duty Cycle} \left(\%\right) = \frac{t\_m}{t\_m + t\_{of}} \times 100\tag{2}$$

The precise variation in duty cycle provides the control over electrochemical processes by affecting the diffusion layer, grain size and nucleation. Usually in electroplating a negatively charged layer is formed around the cathode which gets charged to a known thickness and prevents the ions from the bulk. In DC, this charged double layer prevents the passage of ions toward the cathode thereby affecting the features of the deposit and also causes the inhomogeneity in the concentration of ions in the electrolyte. However, in PED, the charged diffusion layer gets discharged and helps easier passage of the ions onto the cathode as the output is periodically turned off. Migration of ions to depleted areas in the bath during off-time makes the even distribution of ions for their easy availability during on-time. The presence of off time aids in the relaxation and the rearrangement of deposited atoms leading to the possibility of new nucleation sites during the subsequent deposition. This, in turn, not only improves the uniformity of deposition but also reduces the porosity and roughness of the deposit (see **Figure 2**). In addition, the entrapped hydrogen and impurities during the deposition diffuse out during the relaxation time. The additional process variables in PED ease the process of optimization for complex stoichiometric ternary/quaternary systems as the control over individual elemental composition is far better with the variation in duty cycle compared to DC technique.

#### **3.3. Pulse reverse electrodeposition**

Pulse reverse electrodeposition contains short anodic pulses alongside the cathodic as seen in PED (See **Figure 3**). This small anodic pulse is advantageous since it contributes to the electro-oxidation of the top layer from the deposited film thereby aids in smoothening of the deposit, removing the impurities as well as the entrapped hydrogen. The anodic pulse plays a major role in the deposition of systems like CIS and CIGS, wherein an undesired

**Figure 2.** Schematic of pulse electrodeposition and expected growth process of the deposit.

**Figure 3.** Schematic of pulse reverse electrodeposition and expected growth process of the deposit.

secondary Cu-Se phase exists on the surface. With the appropriate control of the anodic pulse, excess copper and the undesired phases can be easily eliminated from the deposited films thereby, forming a phase-pure CIS/CIGS. Similar to PED deposition, the variation in pulse parameters is expressed using a common parameter known as duty cycle, which is defined by the equation:

$$\text{Duty Cycle} \left(\% \right) = \frac{t\_{\text{\\_}}}{t\_{\text{\\_}} + t\_{\text{\\_}}} \times 100 \tag{3}$$

One-step electrodeposition of CIS is usually carried out in an aqueous solution often contain-

often contains a complexing agent in order to shift the reduction potentials of Cu and In closer together to improve the film quality. Complexing agents such as citric acid/citrate [35], ammonia [28], triethanolamine [28, 36], thiocyanate [37], etc., are used during the one-step electrodeposition of CIS thin-films. In addition, a supporting electrolyte such as chloride (LiCl

conductivity of the electrolyte leading to easier mobility of the precursor ions. Often amorphous or poorly crystalline CIS films were observed from electrodeposition which contains frequently degenerate Cu2-xSe phases that are detrimental to the device performance [30, 37]. Also, Cu-rich films have generally larger grain sizes than stoichiometric or In-rich films. Due to these reasons, the electrodeposited CIS films often required to be annealed in a selenium

Incorporation of Ga into the CIS thin-films, to improve the desired properties, was a challenging task for long time for the formation of quaternary CIGS thin-films. However, this bottleneck has been overcome in the recent past. Several researchers have reported the successful incorporation of Ga in the films up to desired range of amounts (6–8 at. %) for the preparation of high efficiency cells [22, 34]. Bhattacharya et al. were the first to report the insertion of Ga from a chloride bath, but to a very low content Ga/In ≈ 0.1, wherein a superimposed alternating voltages has been used at 20 kHz [44]. But the breakthrough for the incorporation of Ga has been realized a little later when the group had used a pH buffer in the chloride bath, also known as Hydrion buffer (pH = 3) consisting of sulfamic acid and potassium hydrogen phthalate. The deposition potential was kept constant while the solution composition has been varied to realize the real possibility of incorporating Ga with the ratio Ga/In from 0.3 to 0.7. This process demonstrated the formation of CIGS films over a wide range of compositions suitable for efficient solar cells by one-step electrodeposition technique. CIGS layers generally deposited from the above mentioned electrodeposition technique often used an additional PVD step to achieve the required composition to form stoichiometric films [45, 46]. Bhattacharya et al. demonstrated a cell efficiency of 9.4% by using a similar PVD step to improve the composition of In and Ga in the as-deposited CIGS thin-films [39]. Valderrama et al. explored a similar electrodeposition technique followed by PVD step to achieve stoichiometric chalcopyrite CIGS films and additionally demonstrated the use of CIGS films to produce hydrogen by the use of

SO4

phases was essentially a key point to achieve high quality CIGS thin-films and hence, higher efficiency. The formation of secondary phases was successfully prevented by the pretreatment of Mo substrate wherein a 1 min pre-deposition of CIGS was performed and a multi-potential

In the similar context, addition of Ga seems to cause morphology related problems, often a concern during the electrodeposition of CIGS films. Fernandez et al. varied the concentration of precursor solution systematically and achieved a better control over composition and morphology of the CIGS films [39]. The presence of cracks in Ga-containing layers is often a serious problem [39, 47], though it can be reduced through the use of alcohol–aqueous solutions [48] or support-

deposition regime was employed to obtain crack-free CIGS layers [47].

SO4

SO4

atmosphere to correct for the stoichiometry and improve crystallinity.

, In3+, and SeO<sup>2</sup>

/H<sup>2</sup>

Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films

[41–43]) is added which results in an improved

[46]. Avoiding the formation of secondary

with gelatin as brightening additive [49]. Complexing

SeO3. The deposition solution

115

http://dx.doi.org/10.5772/intechopen.71857

ing chloride/sulfate precursors of Cu2+ or Cu+

[22, 34, 38, 39] or NaCl [40]) or sulfate (K2

photoelectrochemical testing of the films in H<sup>2</sup>

ing electrolytes such as LiCl or Li<sup>2</sup>

#### **3.4. Electrodeposition of ternary/quaternary chalcopyrites**

The goal of the electrodeposition (ED) is to assure an adherent, compact and a laterally uniform film with the desired stoichiometry. Lateral compositional uniformity is essential over large areas for commercialization of devices. The properties of the deposit solely depend on the control of individual parameters during ED such as electrode material, precursor concentration in electrolyte, applied potential/current and temperature. For single metallic systems like Cu or Zn these properties are well understood. However, deposition of CIGS contains multiple elements including a chalcogen making the ED process substantially complex. Despite that electrochemical deposition appears to be a promising technique for the low-cost solution preparation of CIGS semiconductors [25–27]. Indeed, electrochemical deposition has been widely investigated for CIGS deposition since the pioneering work by Bhattacharya et al. in 1983 [28]. Since then, more than 350 publications have been devoted to the electrochemical preparation of CIGS and several review papers have appeared [25, 29, 30], making electrodeposition the most intensely studied non-vacuum deposition method for CIGS. In light of this, this section reviews and summarizes the previously published work.

The first electrochemical approach to deposit polycrystalline CIS was reported by Bhattacharya in 1983, wherein Cu, In and Se were simultaneously deposited from an acidic solution [28]. Quickly after the first report, several approaches have been proposed for the synthesis of CIS thin films [31–33]. It is of general interest to perform a simultaneous codeposition of all three elements in achieving stoichiometric chalcopyrite CIS films. Such an approach was first initiated by Bhattacharya [28], which is by far the most investigated case as it involves only one electrochemical process, often termed as one-step electrodeposition. Typical electrochemical reactions involving the reduction of elements and formation of CIS/CIGS phase are detailed in [34].

One-step electrodeposition of CIS is usually carried out in an aqueous solution often containing chloride/sulfate precursors of Cu2+ or Cu+ , In3+, and SeO<sup>2</sup> /H<sup>2</sup> SeO3. The deposition solution often contains a complexing agent in order to shift the reduction potentials of Cu and In closer together to improve the film quality. Complexing agents such as citric acid/citrate [35], ammonia [28], triethanolamine [28, 36], thiocyanate [37], etc., are used during the one-step electrodeposition of CIS thin-films. In addition, a supporting electrolyte such as chloride (LiCl [22, 34, 38, 39] or NaCl [40]) or sulfate (K2 SO4 [41–43]) is added which results in an improved conductivity of the electrolyte leading to easier mobility of the precursor ions. Often amorphous or poorly crystalline CIS films were observed from electrodeposition which contains frequently degenerate Cu2-xSe phases that are detrimental to the device performance [30, 37]. Also, Cu-rich films have generally larger grain sizes than stoichiometric or In-rich films. Due to these reasons, the electrodeposited CIS films often required to be annealed in a selenium atmosphere to correct for the stoichiometry and improve crystallinity.

secondary Cu-Se phase exists on the surface. With the appropriate control of the anodic pulse, excess copper and the undesired phases can be easily eliminated from the deposited films thereby, forming a phase-pure CIS/CIGS. Similar to PED deposition, the variation in pulse parameters is expressed using a common parameter known as duty cycle, which is

**Figure 3.** Schematic of pulse reverse electrodeposition and expected growth process of the deposit.

The goal of the electrodeposition (ED) is to assure an adherent, compact and a laterally uniform film with the desired stoichiometry. Lateral compositional uniformity is essential over large areas for commercialization of devices. The properties of the deposit solely depend on the control of individual parameters during ED such as electrode material, precursor concentration in electrolyte, applied potential/current and temperature. For single metallic systems like Cu or Zn these properties are well understood. However, deposition of CIGS contains multiple elements including a chalcogen making the ED process substantially complex. Despite that electrochemical deposition appears to be a promising technique for the low-cost solution preparation of CIGS semiconductors [25–27]. Indeed, electrochemical deposition has been widely investigated for CIGS deposition since the pioneering work by Bhattacharya et al. in 1983 [28]. Since then, more than 350 publications have been devoted to the electrochemical preparation of CIGS and several review papers have appeared [25, 29, 30], making electrodeposition the most intensely studied non-vacuum deposition method for CIGS. In light of this,

The first electrochemical approach to deposit polycrystalline CIS was reported by Bhattacharya in 1983, wherein Cu, In and Se were simultaneously deposited from an acidic solution [28]. Quickly after the first report, several approaches have been proposed for the synthesis of CIS thin films [31–33]. It is of general interest to perform a simultaneous codeposition of all three elements in achieving stoichiometric chalcopyrite CIS films. Such an approach was first initiated by Bhattacharya [28], which is by far the most investigated case as it involves only one electrochemical process, often termed as one-step electrodeposition. Typical electrochemical reactions involving the reduction of elements and formation of CIS/CIGS phase are detailed in [34].

*t <sup>a</sup>* + *t c* × 100 (3)

defined by the equation:

114 Semiconductors - Growth and Characterization

*Duty Cycle* (%) <sup>=</sup> *<sup>t</sup>* \_\_\_\_*<sup>c</sup>*

**3.4. Electrodeposition of ternary/quaternary chalcopyrites**

this section reviews and summarizes the previously published work.

Incorporation of Ga into the CIS thin-films, to improve the desired properties, was a challenging task for long time for the formation of quaternary CIGS thin-films. However, this bottleneck has been overcome in the recent past. Several researchers have reported the successful incorporation of Ga in the films up to desired range of amounts (6–8 at. %) for the preparation of high efficiency cells [22, 34]. Bhattacharya et al. were the first to report the insertion of Ga from a chloride bath, but to a very low content Ga/In ≈ 0.1, wherein a superimposed alternating voltages has been used at 20 kHz [44]. But the breakthrough for the incorporation of Ga has been realized a little later when the group had used a pH buffer in the chloride bath, also known as Hydrion buffer (pH = 3) consisting of sulfamic acid and potassium hydrogen phthalate. The deposition potential was kept constant while the solution composition has been varied to realize the real possibility of incorporating Ga with the ratio Ga/In from 0.3 to 0.7. This process demonstrated the formation of CIGS films over a wide range of compositions suitable for efficient solar cells by one-step electrodeposition technique. CIGS layers generally deposited from the above mentioned electrodeposition technique often used an additional PVD step to achieve the required composition to form stoichiometric films [45, 46]. Bhattacharya et al. demonstrated a cell efficiency of 9.4% by using a similar PVD step to improve the composition of In and Ga in the as-deposited CIGS thin-films [39]. Valderrama et al. explored a similar electrodeposition technique followed by PVD step to achieve stoichiometric chalcopyrite CIGS films and additionally demonstrated the use of CIGS films to produce hydrogen by the use of photoelectrochemical testing of the films in H<sup>2</sup> SO4 [46]. Avoiding the formation of secondary phases was essentially a key point to achieve high quality CIGS thin-films and hence, higher efficiency. The formation of secondary phases was successfully prevented by the pretreatment of Mo substrate wherein a 1 min pre-deposition of CIGS was performed and a multi-potential deposition regime was employed to obtain crack-free CIGS layers [47].

In the similar context, addition of Ga seems to cause morphology related problems, often a concern during the electrodeposition of CIGS films. Fernandez et al. varied the concentration of precursor solution systematically and achieved a better control over composition and morphology of the CIGS films [39]. The presence of cracks in Ga-containing layers is often a serious problem [39, 47], though it can be reduced through the use of alcohol–aqueous solutions [48] or supporting electrolytes such as LiCl or Li<sup>2</sup> SO4 with gelatin as brightening additive [49]. Complexing agents such as citric acid/citrate [50, 51], thiocyanate [52], sodium sulfamate [53], sulfosalicylic acid [54], etc., were often used to improve the composition and morphology of the CIGS films. These additives form complexes with the metal ions in the solution such as Cu, thereby resulting in controlled deposition rates and hence the morphology [55, 56]. Good quality CIGS thin films were also prepared from sulfate–citrate and chloride–citrate solutions [53, 57, 58] and control of the optical band gap by increasing the Ga content in the films was demonstrated [59].

**Absorber material**

CIGSe and CIGSSe

CIGSe and CIGSSe

CIGSe Mo/glass Cu-In-Ga oxide

CIGSe Mo/glass Stacked Cu/In/

CIGSe Mo/glass Stacked Cu/In/Ga

Mo/stainless steel foil

**Substrate Preparation method Reported efficiency/**

precursor followed by thermochemical reduction and selenization

Ga followed by selenization

layers followed by selenization

Four step electrodeposition of CIG layers followed by selenization

CuInGa layers annealed using Rapid thermal Processing under S and Se vapors

precursors followed by rapid thermal processing in S and Se

stacked Cu/In/Ga layers are annealed by RTP under S and Se atmosphere

CIGSe layers followed by physical vapor deposition of In, Ga, Se and selenized

atmosphere

Mo/glass Electrodeposited

CIGSSe Mo/glass Electrodeposited

CIGSSe Mo/glass Electrodeposited

CIGSe Mo/glass Electrodeposited

**photocurrent (cell area** 

15.36% for CIGSe (5.4) 13.4% for CIGSSe (3824.6,

14.1% for CIGSe (0.5) 15.8% for CIGSSe (0.5)

10.6 (0.1) for CISSe 9.9 (0.1) for CISe

15.4% for ED CIGSe (0.4) 12.4% for EL CIGSe (0.4)

submodule)

12.4% (0.1) A new approach

Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films

11.7 (0.4) Step-by-step

12.6 (0.1) Impact of thickness

by one-step electrochemical deposition of metal oxides followed by thermochemical reduction and selenization

electrodeposition of Cu/In/Ga followed by selenization

is studied. 8.7% efficiency is achieved for an ultrathin 370 nm film of CIGS

Development of roll-to-roll electrodeposited CIGS solar cells on flexible stainless steel

substrates

Rapid thermal processing of electrodeposited CIG layers under S and Se atmosphere

Sequentially electrodeposited copper indium precursors are annealed by RTP under S and Se atmosphere

CIGSSe solar cells are made wherein i-ZnO/AZO window layer is deposited using solution based

approach.

Physical vapor deposition of In, Ga and Se was performed on electrodeposited and electroless deposited CIGS layers for compositional adjustment

13.8 (0.09) All solution processed

**Remarks Reference**

http://dx.doi.org/10.5772/intechopen.71857

[82]

117

[22]

[83]

[21]

[84]

[26]

[85]

[23]

**in cm2 )**

To overcome the difficulty of In and Ga incorporation, alternative strategies to co-electrodeposition have been developed, often involving the deposition of stacked elemental layers or else deposition of alloys, followed by a selenization or sulphurization treatment to provide all of the chalcogen [21, 34]. Deposition of Cu-Ga [60] and Cu-In-Ga [52] alloys has been demonstrated, with the latter approach leading to 4% efficiency cells. Solopower explored the layer by layer electrodeposition to form stoichiometric CIGS thin-films [21]. The process also used several varieties of complexing agents, organic additives, etc., to correct the composition and improve the morphology of the films. In addition, current densities were varied during deposition and annealing temperatures were optimized to achieve high quality compact large grained stoichiometric chalcopyrite CIGS films. Devices and modules were fabricated using these films which had shown an efficiency of 15.36% on an area of 5.34 cm<sup>2</sup> and 13.4% on an area of 3.8 m<sup>2</sup> . Most recently, electrodeposited CIGSSe devices have yielded conversion efficiencies upto 17.3% wherein electrodeposited Cu-In-Ga stack is rapid thermal treated in elemental selenium and sulfur atmosphere. Some of the recent notable reports with direct current electrodeposited CIGS solar cells are summarized in **Table 1**.

In addition to the synthesis of CIS and CIGS thin films by DC electrodeposition, PED technique has also been studied. Owing to the advantages of PED discussed in Section 3.2, it results in superior quality thin-films and enables one to avoid additional steps of deposition, post deposition treatments and the use of complexing agents. Kang et al. reported the preparation of CIS thin-films by pulse-reverse electrodeposition followed by selenization wherein stoichiometric CIS films with rough surface morphology exhibited an efficiency of 1.42% [61]. Li et al. employed a square wave modulated by a bell-like wave during the pulse plating for the fabrication of CIS thin-films. The study reported well adherent chalcopyrite CIS films with a uniform morphology [35]. A threestep pulse electrodeposition method was used for the fabrication of CIS thin-films reporting a mixture of phases such as Cu-Se, In-Se, CIS, etc., as confirmed from Raman and optical studies of the samples [62]. Valdes et al. employed different potentials during the pulses which resulted in chalcopyrite p-type CIS thin-films with different morphologies and compositions [63]. Murali et al. prepared CIS thin-films using pulse electrodeposition with varied duty cycle from 6 to 50% and reported the p-type phase pure CIS films with resistivities in the range of 1–10 ohm-cm [64]. Hu et al. employed the deposition of CIS films by pulse electrodeposition wherein multi potentials were used to control the composition of the films which resulted in improved deposition uniformity without any secondary phases [65]. Similar reports on the compositional control of CIS thinfilms by the variation of pulse parameters have appeared in the recent past [66–69].

Fu et al., have explored different plating techniques including DC, pulse and pulse reverse electrodeposition for the fabrication of CIGS thin films and reported the elimination of undesired secondary phases like Cu2−xSe to obtain single phase pure chalcopyrite CIGS thin films [70]. Liu et al., have employed the PED with the variation of duty cycle to remove the excess In and to avoid In-Se compounds during the deposition for the preparation of single phase CIGS


agents such as citric acid/citrate [50, 51], thiocyanate [52], sodium sulfamate [53], sulfosalicylic acid [54], etc., were often used to improve the composition and morphology of the CIGS films. These additives form complexes with the metal ions in the solution such as Cu, thereby resulting in controlled deposition rates and hence the morphology [55, 56]. Good quality CIGS thin films were also prepared from sulfate–citrate and chloride–citrate solutions [53, 57, 58] and control of

To overcome the difficulty of In and Ga incorporation, alternative strategies to co-electrodeposition have been developed, often involving the deposition of stacked elemental layers or else deposition of alloys, followed by a selenization or sulphurization treatment to provide all of the chalcogen [21, 34]. Deposition of Cu-Ga [60] and Cu-In-Ga [52] alloys has been demonstrated, with the latter approach leading to 4% efficiency cells. Solopower explored the layer by layer electrodeposition to form stoichiometric CIGS thin-films [21]. The process also used several varieties of complexing agents, organic additives, etc., to correct the composition and improve the morphology of the films. In addition, current densities were varied during deposition and annealing temperatures were optimized to achieve high quality compact large grained stoichiometric chalcopyrite CIGS films. Devices and modules were fabricated

the optical band gap by increasing the Ga content in the films was demonstrated [59].

using these films which had shown an efficiency of 15.36% on an area of 5.34 cm<sup>2</sup>

films by the variation of pulse parameters have appeared in the recent past [66–69].

Fu et al., have explored different plating techniques including DC, pulse and pulse reverse electrodeposition for the fabrication of CIGS thin films and reported the elimination of undesired secondary phases like Cu2−xSe to obtain single phase pure chalcopyrite CIGS thin films [70]. Liu et al., have employed the PED with the variation of duty cycle to remove the excess In and to avoid In-Se compounds during the deposition for the preparation of single phase CIGS

current electrodeposited CIGS solar cells are summarized in **Table 1**.

efficiencies upto 17.3% wherein electrodeposited Cu-In-Ga stack is rapid thermal treated in elemental selenium and sulfur atmosphere. Some of the recent notable reports with direct

In addition to the synthesis of CIS and CIGS thin films by DC electrodeposition, PED technique has also been studied. Owing to the advantages of PED discussed in Section 3.2, it results in superior quality thin-films and enables one to avoid additional steps of deposition, post deposition treatments and the use of complexing agents. Kang et al. reported the preparation of CIS thin-films by pulse-reverse electrodeposition followed by selenization wherein stoichiometric CIS films with rough surface morphology exhibited an efficiency of 1.42% [61]. Li et al. employed a square wave modulated by a bell-like wave during the pulse plating for the fabrication of CIS thin-films. The study reported well adherent chalcopyrite CIS films with a uniform morphology [35]. A threestep pulse electrodeposition method was used for the fabrication of CIS thin-films reporting a mixture of phases such as Cu-Se, In-Se, CIS, etc., as confirmed from Raman and optical studies of the samples [62]. Valdes et al. employed different potentials during the pulses which resulted in chalcopyrite p-type CIS thin-films with different morphologies and compositions [63]. Murali et al. prepared CIS thin-films using pulse electrodeposition with varied duty cycle from 6 to 50% and reported the p-type phase pure CIS films with resistivities in the range of 1–10 ohm-cm [64]. Hu et al. employed the deposition of CIS films by pulse electrodeposition wherein multi potentials were used to control the composition of the films which resulted in improved deposition uniformity without any secondary phases [65]. Similar reports on the compositional control of CIS thin-

. Most recently, electrodeposited CIGSSe devices have yielded conversion

on an area of 3.8 m<sup>2</sup>

116 Semiconductors - Growth and Characterization

and 13.4%


**Table 1.** Notable reports on direct current electrodeposited CIS/CIGS with reported efficiencies.

thin films [58]. In one of the most successful efforts of fabrication of CIGS by pulse electrodeposition, Bi et al. have effectively utilized the parameters in electrodeposition to demonstrate CIGS solar cells with conversion efficiencies upto 10.39% [71] and 11.04% [72]. Cu-In-Ga metal precursors were electrodeposited by pulse current method wherein the charge density was chosen such that the desirable thickness of each layer is achieved. Electrodeposited metal stack was annealed in a three step process to achieve dense large grained CIGS layer which yielded high conversion efficiencies. **Table 2** summarizes the recent literature reported for CIGS thin films fabrication using pulse electrodeposition.


Despite having several advantages, pulse and pulse reverse electrodeposition techniques are underutilized for the fabrication of CIGS thin films. Considering the additional process variables these techniques offer, thin films photovoltaic community might consider exploring these processes for fabricating high quality CIGS films. Our group has explored these features while developing a simplified pulsed electrodeposition technique for the fabrication of CIGS

**Table 2.** Summary of recent reports on pulse electrodeposited CIS/CIGS for application in thin films solar cells.

**Materials Substrate Preparation method Notable inference Remarks Reference**

Photocurrents in μA range are observed

PEC performance reported

PEC performance reported

PEC performance reported

Variation in duty cycle in pulse reverse plating yielded device quality

http://dx.doi.org/10.5772/intechopen.71857

conditions for each layer of Cu, In and Ga during pulse electrodeposition

Duty cycle has been optimized for device quality CIGS films

A novel nanoflake morphology has been reported while obtaining stoichiometric CISe with optimization in duty

DC and pulse plated CIGS films are compared. Optimization in pulse plating improves the morphology and crystallinity of CIGSe

A sequential pulse electrodeposition approach is proposed so as to minimize the utilization of In concentration during electrodeposition of CIGSe films

cycle

[87]

119

[71]

[88]

[76]

[75]

[73]

CIGS films

10.39% (0.34) Optimization of pulse

Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films

A new simplified pulse electrodeposition approach for the fabrication of CIS and CIGS films with compact and nanostructured morphologies was developed [73–76]. The approach utilizes a two electrode system by avoiding the conventional third reference electrode for the deposition of thin films. The commonly used expensive platinum counter electrode is substituted by

**4. Pulsed electrodeposition of CIS and CIGS films: case studies**

thin films and is discussed below.

CIGSe Mo/glass DC and pulse reverse

CIGSe Mo/glass Step by step pulse

CIGS ITO Pulse electrodeposition

CISe Mo foil Pulse electrodeposition

CIGSe Mo foil DC and pulse plating of

CIGSe Mo/glass Two stage pulse

plating of CIGS followed by selenization

electrodeposition of Cu/ In/Ga films followed by

of CIGS followed by

of CIS followed by annealing

CIGSe films followed by

electrodeposition of CuGaSe and In films followed by annealing

selenization

annealing

annealing


thin films [58]. In one of the most successful efforts of fabrication of CIGS by pulse electrodeposition, Bi et al. have effectively utilized the parameters in electrodeposition to demonstrate CIGS solar cells with conversion efficiencies upto 10.39% [71] and 11.04% [72]. Cu-In-Ga metal precursors were electrodeposited by pulse current method wherein the charge density was chosen such that the desirable thickness of each layer is achieved. Electrodeposited metal stack was annealed in a three step process to achieve dense large grained CIGS layer which yielded high conversion efficiencies. **Table 2** summarizes the recent literature reported for

**Materials Substrate Preparation method Notable inference Remarks Reference**

11.04% efficiency for pulse plated CIGSe (0.34) 8.18% efficiency for direct current deposited CIGSe (0.34)

CIGS exhibited p-type and dark I-V curves are recorded

Efficiency of 1.42% on 0.4 cm<sup>2</sup>

PEC curves are recorded indicating photoactivity of CIGSe films

cell

electrolyte

**photocurrent (cell area** 

Photocurrent of 1 mA/cm<sup>2</sup> for CISe and for CIGSe in ethyl viologen perchlorate

10.9% (0.4) Electrodeposited

**Remarks Reference**

[38]

[86]

[72]

[70]

[61]

[58]

precursor layers are prepared in a three stage sequential route wherein Cu and In second and third layers are deposited on

CIGSe layer

Direct current and pulsed current plating are explored for step by step Cu/In/Ga layers and pulse plated devices exhibited higher

Use of pulse reverse potential eliminated the undesired CuSe phase

Pulse reverse plating is employed to make CISe with device quality

Variation in duty cycle is adopted and growth mechanism is studied

efficiency

features

CISe films were deposited from spray coating of nanoparticles and selenized while CIGSe were electrodeposited and selenized

**in cm2 )**

CIGS thin films fabrication using pulse electrodeposition.

DC and PC methods followed by selenization

reverse plating of CIGS films followed by annealing in Ar

electrodeposition of CISe followed by selenization

of CIGSe followed by

annealing

CIGSe Mo/glass Stacked Cu/In/Ga by

CIGSe Mo foil DC, Pulse and pulse

CIGSe Mo foil Pulse electrodeposition

CISe Mo/glass Pulse reverse

**Substrate Preparation method Reported efficiency/**

electrodeposition (CIGSe/Cu/In) process followed by annealing under Se atmosphere

CIGSe is annealed under selenium atmosphere and spray coated CIS nanoparticles followed by selenization

**Table 1.** Notable reports on direct current electrodeposited CIS/CIGS with reported efficiencies.

**Absorber material**

CIGSe Mo/glass Three stage

118 Semiconductors - Growth and Characterization

CIGSe CISe Mo/glass Electrodeposited

**Table 2.** Summary of recent reports on pulse electrodeposited CIS/CIGS for application in thin films solar cells.

Despite having several advantages, pulse and pulse reverse electrodeposition techniques are underutilized for the fabrication of CIGS thin films. Considering the additional process variables these techniques offer, thin films photovoltaic community might consider exploring these processes for fabricating high quality CIGS films. Our group has explored these features while developing a simplified pulsed electrodeposition technique for the fabrication of CIGS thin films and is discussed below.

#### **4. Pulsed electrodeposition of CIS and CIGS films: case studies**

A new simplified pulse electrodeposition approach for the fabrication of CIS and CIGS films with compact and nanostructured morphologies was developed [73–76]. The approach utilizes a two electrode system by avoiding the conventional third reference electrode for the deposition of thin films. The commonly used expensive platinum counter electrode is substituted by high purity graphite electrode thereby cutting down the cost. The technique by virtue of utilization of pulse plating avoids the use of any additives and complexing agents during electrodeposition while obtaining the stoichiometric films in a single step thereby eliminating multi-step deposition. The films are annealed under Ar atmosphere while avoiding the conventional selenium atmosphere making it an environmental friendly approach. In addition, a novel strategy has been developed to address the scarcity of Indium wherein In precursor has been effectively minimized compared to existing reports. Typical pulse electrodeposition set-up, applied pulse voltage wave form and the corresponding current density are depicted in **Figure 4**.

Since the approach avoids the use of reference electrode and employs a two-electrode system, prior to electrodeposition, cyclic voltammetry of the CIGS system using a two and three-electrode system are studied as shown in **Figure 5**. The reduction peak of CIGS is shifted from −0.8 to −1.1 V from three to two-electrode systems. Apparently, the potential shift is approximately same as the standard electrode potential of SCE vs. normal hydrogen electrode (NHE). However, experiments with different systems might be required to verify the same. The inference is utilized to adopt higher deposition potentials in the pulsed electrodeposition of CIS and CIGS films.

#### **4.1. Study I: pulse electrodeposition of CuInSe2 films**

The pulse electrodeposition of CuInSe<sup>2</sup> (CIS) films is performed from a bath containing metallic chlorides of Cu, In and Ga while selenous acid is used as selenium precursor dissolved in pH 3 Hydrion buffer. The final pH of the bath is maintained around 2.0–2.5. Additives and complexing agents are completely avoided in this study. CIS films were electrodeposited using a deposition potential of −1.5 V while the variation in duty cycle has been studied. The applied pulse-voltage and the corresponding current density curves from **Figure 6** note that there is a small positive current density during the pulse off-time though no voltage is applied. This could be due to the presence of an electric double layer at cathode-electrolyte interface forming a capacitor of molecular dimension [58, 77]. The duty cycle for the application of pulses is varied in the range of 17–67% (varied off-time with fixed on-time). The PED technique employed in the present study did not result in any disruption and dissolution of the deposited film into the electrolyte, which generally happens with higher deposition voltages as previously reported [78]. The PED deposited CIS films are annealed at 550°C for 30 min under Ar atmosphere.

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**Figure 6.** (a) Nanoflake morphology, (b) Raman spectra, (c) Tauc's plot and (d) photoelectrochemical I-V characteristics

of pulse electrodeposited and annealed CIS films.

**Figure 5.** Cyclic voltammogram of Cu-In-Ga-Se system using a three-electrode and two-electrode configuration.

**Figure 4.** (a) Pulse electrodeposition set-up, (b) applied deposition voltage waveform and (c) corresponding current density employed for pulse electrodeposition of CIS and CIGS thin films.

Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films http://dx.doi.org/10.5772/intechopen.71857 121

high purity graphite electrode thereby cutting down the cost. The technique by virtue of utilization of pulse plating avoids the use of any additives and complexing agents during electrodeposition while obtaining the stoichiometric films in a single step thereby eliminating multi-step deposition. The films are annealed under Ar atmosphere while avoiding the conventional selenium atmosphere making it an environmental friendly approach. In addition, a novel strategy has been developed to address the scarcity of Indium wherein In precursor has been effectively minimized compared to existing reports. Typical pulse electrodeposition set-up, applied pulse

Since the approach avoids the use of reference electrode and employs a two-electrode system, prior to electrodeposition, cyclic voltammetry of the CIGS system using a two and three-electrode system are studied as shown in **Figure 5**. The reduction peak of CIGS is shifted from −0.8 to −1.1 V from three to two-electrode systems. Apparently, the potential shift is approximately same as the standard electrode potential of SCE vs. normal hydrogen electrode (NHE). However, experiments with different systems might be required to verify the same. The inference is utilized to

 **films**

lic chlorides of Cu, In and Ga while selenous acid is used as selenium precursor dissolved in pH 3 Hydrion buffer. The final pH of the bath is maintained around 2.0–2.5. Additives and complexing agents are completely avoided in this study. CIS films were electrodeposited using a deposition potential of −1.5 V while the variation in duty cycle has been studied. The applied pulse-voltage and the corresponding current density curves from **Figure 6** note that there is a small positive current density during the pulse off-time though no voltage is applied. This could be due to the presence of an electric double layer at cathode-electrolyte interface forming a capacitor of molecular dimension [58, 77]. The duty cycle for the application of pulses is varied in the range of 17–67% (varied off-time with fixed on-time). The PED technique employed in

**Figure 4.** (a) Pulse electrodeposition set-up, (b) applied deposition voltage waveform and (c) corresponding current

density employed for pulse electrodeposition of CIS and CIGS thin films.

(CIS) films is performed from a bath containing metal-

voltage wave form and the corresponding current density are depicted in **Figure 4**.

adopt higher deposition potentials in the pulsed electrodeposition of CIS and CIGS films.

**4.1. Study I: pulse electrodeposition of CuInSe2**

The pulse electrodeposition of CuInSe<sup>2</sup>

120 Semiconductors - Growth and Characterization

**Figure 5.** Cyclic voltammogram of Cu-In-Ga-Se system using a three-electrode and two-electrode configuration.

the present study did not result in any disruption and dissolution of the deposited film into the electrolyte, which generally happens with higher deposition voltages as previously reported [78]. The PED deposited CIS films are annealed at 550°C for 30 min under Ar atmosphere.

**Figure 6.** (a) Nanoflake morphology, (b) Raman spectra, (c) Tauc's plot and (d) photoelectrochemical I-V characteristics of pulse electrodeposited and annealed CIS films.

The compositional analysis indicated that with an increase in the pulse off time from about 5 to 50 ms during electrodeposition, the relative content of In in the film decreases from almost 40 to 2 at. %. This is due to the unintended positive current density observed in **Figure 4c**, which oxidizes the elements with least electronegativity from the deposited film leading to the dissolution of corresponding element into the electrolyte. A duty cycle of 50% has been considered optimal for obtaining stoichiometric CIS films. Morphological analysis reveal that the CIS films deposited with optimal conditions possess novel nanoflake-like architectures (see **Figure 6a**). Such morphology is expected to be advantageous since they possess high surface area of the film thereby causing improvement in absorption of light and photoresponse. This is also expected to be advantageous at the device stage since it facilitates increase in the p-n junction interface area of the solar cell which directly influences its performance. Additionally, flake like crystallite structure is favorable to increase current carrier concentration, electron transmission and thus induce the generation of photocurrent [79]. Another interesting aspect from this study is the effective reduction in the secondary phases during the electrodeposition of CIS/CIGS films. The variation in duty cycle explored in the present study yielded the films with reduced copper selenide phase for optimal conditions. Raman spectra of the electrodeposited and annealed CIS films are shown in **Figure 6b**. The spectra contain well-defined peaks at wavenumbers of 172, 215, 234 and 260 cm−1 and are attributed to the A<sup>1</sup> , B2 and E modes of CIS films and the peak at 260 cm−1 is attributed to the A<sup>1</sup> mode of Cu2-xSe. As the spectra reveal the optimized condition with relative Cu/ In ratio close to 1, shows least intense peak for copper selenide thereby affirming its reduction. Bandgap of CIS films as obtained from optical absorption spectroscopy is 1.02 eV (from **Figure 6c**), which matches well with the theoretical value of 1.04 eV for stoichiometric chalcopyrite CIS films [80]. Photoelectrochemical performance of CIS films as shown in **Figure 6d** confirms the photoactivity and p-type conductivity.

#### **4.2. Study II: comparison of direct current and pulse electrodeposited CIGS films**

and (424) for CIGS are observed thereby confirming the presence of chalcopyrite CIGS phase. Additionally, Mo substrate and MoSe<sup>2</sup> peaks are observed [81]. XRD pattern also reveals the presence of copper selenide in DC electroplated CIGS while it is absent in PC plated films indicating the phase-pure CIGS formation by PC approach. **Figure 7d** shows the

**Figure 7.** (a) and (b) surface morphologies, (c) XRD patterns and (d) Raman spectra of direct current and pulse

modes of the CIGS at 176, 205 and 232 cm−1, respectively. In addition, a less intense peak cor-

thin film. The copper rich composition in DC plated CIGS films facilitated the formation of the Cu2−xSe phase, which is generally dispersed on the surface [84, 87]. PC electrodeposition with suitable optimization of parameters aided control over the composition of elements and eliminated Cu2−xSe phase. The bandgap of direct current and pulse plated CIGS films is obtained from Tauc's plots as shown in **Figure 8a**, which are determined to be 1.21 and 1.29 eV, respectively. The photoelectrochemical J-V characteristics of annealed DC and PC

cathodic current with potential confirms the p-type conductivity. Also, the PC plated CIGS films show lower dark current and higher photocurrent compared to DC plated films, which could be attributed to the dense morphology with stoichiometric chalcopyrite CIGS without secondary phases. **Figure 8b** shows the amperometric J-t curve confirming the photoactivity

mode of Cu2−xSe at 260 cm−1 is found in case of DC electrodeposited CIGS

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123

SO4

, B2

are shown in **Figure 8b**. Increase in

and E

Raman spectra of the DC and PC electrodeposited CIGS thin films and contain A<sup>1</sup>

electrodeposited CIGS thin films studied in 0.5 M Na<sup>2</sup>

of CIGS under chopped illumination.

responding to A<sup>1</sup>

electrodeposited CIGS films.

CIGS thin films are deposited using direct current (DC) and pulsed current (PC) methods from a bath of metal chlorides and selenous acid in pH 3 Hydrion buffer devoid of additives and complexing agents. **Figure 7a** and **b** shows the surface morphologies of DC and PC plated CIGS films. DC deposited CIGS films exhibited porous and rough morphology with finer spherical particles while dense uniform films with coarser spherical particles are observed for PC deposited CIGS. PC method can produce relatively more homogeneous surface with good adhesion to the substrate because the rate-determining step of the deposition process is controlled by a mass-transfer process. Relaxation during the pulse off-time in the PC electrodeposition not only allows the diffusion of ad atoms but also facilitates the formation of new nucleation sites thereby leading to the homogeneous and compact morphology unlike the DC deposition wherein building-up of material takes place at same nucleation sites leading to the roughness of the film. The stoichiometry of the DC and PC electrodeposited CIGS films is determined, from EDS, to be Cu1.10In0.54Ga0.23Se2.13 and Cu0.98In0.73Ga0.25Se2.03, respectively, indicating copper rich composition in DC CIGS films.

**Figure 7c** shows the XRD patterns of annealed DC and PC electroplated CIGS films, which show the preferred orientation corresponding to (112) and other peaks to (211), (220), (312)

The compositional analysis indicated that with an increase in the pulse off time from about 5 to 50 ms during electrodeposition, the relative content of In in the film decreases from almost 40 to 2 at. %. This is due to the unintended positive current density observed in **Figure 4c**, which oxidizes the elements with least electronegativity from the deposited film leading to the dissolution of corresponding element into the electrolyte. A duty cycle of 50% has been considered optimal for obtaining stoichiometric CIS films. Morphological analysis reveal that the CIS films deposited with optimal conditions possess novel nanoflake-like architectures (see **Figure 6a**). Such morphology is expected to be advantageous since they possess high surface area of the film thereby causing improvement in absorption of light and photoresponse. This is also expected to be advantageous at the device stage since it facilitates increase in the p-n junction interface area of the solar cell which directly influences its performance. Additionally, flake like crystallite structure is favorable to increase current carrier concentration, electron transmission and thus induce the generation of photocurrent [79]. Another interesting aspect from this study is the effective reduction in the secondary phases during the electrodeposition of CIS/CIGS films. The variation in duty cycle explored in the present study yielded the films with reduced copper selenide phase for optimal conditions. Raman spectra of the electrodeposited and annealed CIS films are shown in **Figure 6b**. The spectra contain well-defined peaks at wavenumbers of 172, 215, 234 and 260 cm−1 and

and E modes of CIS films and the peak at 260 cm−1 is attributed

mode of Cu2-xSe. As the spectra reveal the optimized condition with relative Cu/

In ratio close to 1, shows least intense peak for copper selenide thereby affirming its reduction. Bandgap of CIS films as obtained from optical absorption spectroscopy is 1.02 eV (from **Figure 6c**), which matches well with the theoretical value of 1.04 eV for stoichiometric chalcopyrite CIS films [80]. Photoelectrochemical performance of CIS films as shown in **Figure** 

CIGS thin films are deposited using direct current (DC) and pulsed current (PC) methods from a bath of metal chlorides and selenous acid in pH 3 Hydrion buffer devoid of additives and complexing agents. **Figure 7a** and **b** shows the surface morphologies of DC and PC plated CIGS films. DC deposited CIGS films exhibited porous and rough morphology with finer spherical particles while dense uniform films with coarser spherical particles are observed for PC deposited CIGS. PC method can produce relatively more homogeneous surface with good adhesion to the substrate because the rate-determining step of the deposition process is controlled by a mass-transfer process. Relaxation during the pulse off-time in the PC electrodeposition not only allows the diffusion of ad atoms but also facilitates the formation of new nucleation sites thereby leading to the homogeneous and compact morphology unlike the DC deposition wherein building-up of material takes place at same nucleation sites leading to the roughness of the film. The stoichiometry of the DC and PC electrodeposited CIGS films is determined, from EDS, to be Cu1.10In0.54Ga0.23Se2.13 and Cu0.98In0.73Ga0.25Se2.03, respectively, indi-

**Figure 7c** shows the XRD patterns of annealed DC and PC electroplated CIGS films, which show the preferred orientation corresponding to (112) and other peaks to (211), (220), (312)

**4.2. Study II: comparison of direct current and pulse electrodeposited CIGS films**

are attributed to the A<sup>1</sup>

122 Semiconductors - Growth and Characterization

to the A<sup>1</sup>

, B2

**6d** confirms the photoactivity and p-type conductivity.

cating copper rich composition in DC CIGS films.

**Figure 7.** (a) and (b) surface morphologies, (c) XRD patterns and (d) Raman spectra of direct current and pulse electrodeposited CIGS films.

and (424) for CIGS are observed thereby confirming the presence of chalcopyrite CIGS phase. Additionally, Mo substrate and MoSe<sup>2</sup> peaks are observed [81]. XRD pattern also reveals the presence of copper selenide in DC electroplated CIGS while it is absent in PC plated films indicating the phase-pure CIGS formation by PC approach. **Figure 7d** shows the Raman spectra of the DC and PC electrodeposited CIGS thin films and contain A<sup>1</sup> , B2 and E modes of the CIGS at 176, 205 and 232 cm−1, respectively. In addition, a less intense peak corresponding to A<sup>1</sup> mode of Cu2−xSe at 260 cm−1 is found in case of DC electrodeposited CIGS thin film. The copper rich composition in DC plated CIGS films facilitated the formation of the Cu2−xSe phase, which is generally dispersed on the surface [84, 87]. PC electrodeposition with suitable optimization of parameters aided control over the composition of elements and eliminated Cu2−xSe phase. The bandgap of direct current and pulse plated CIGS films is obtained from Tauc's plots as shown in **Figure 8a**, which are determined to be 1.21 and 1.29 eV, respectively. The photoelectrochemical J-V characteristics of annealed DC and PC electrodeposited CIGS thin films studied in 0.5 M Na<sup>2</sup> SO4 are shown in **Figure 8b**. Increase in cathodic current with potential confirms the p-type conductivity. Also, the PC plated CIGS films show lower dark current and higher photocurrent compared to DC plated films, which could be attributed to the dense morphology with stoichiometric chalcopyrite CIGS without secondary phases. **Figure 8b** shows the amperometric J-t curve confirming the photoactivity of CIGS under chopped illumination.

**Figure 8.** (a) Tauc's plots and (b) photoelectrochemical J-V characteristics of direct current and pulse electrodeposited CIGS films.

#### **4.3. Study III: sequential pulsed electrodeposition of CIGS thin-films**

The sequential technique has been proposed to essentially minimize In precursor to address the scarcity of In. The usage of In in the growing electronic and optoelectronic industries is very high in the form of materials such as indium doped tin oxide (ITO), CIS, CIGS, InP, InN, InGaAs, InAlAs, etc., making it one of the most scarce elements in the near future. In this context, a novel sequential PC approach is explored for the fabrication of CIGS thin-films. Deposition of Cu-Ga-Se films is carried out by optimizing the deposition voltage in the first stage followed by subsequent deposition of In in the second stage. The sequentially deposited Cu-Ga-Se/In thin-films are annealed in Ar atmosphere and characterized.

The optimized CuGaSe/In annealed films are noted to have a compact morphology (**Figure 9a**), which is well-suited for application as solar absorber layers, since it facilitates easier diffusion of minority charge carriers and reduces recombination. XRD and Raman analyses confirm the presence of chalcopyrite CIGS films without any undesired phases as shown in **Figure 9b** and **c**. In addition to the micro-Raman analysis of CIGS films, Raman mapping of the optimized films is performed to further verify their phase-purity. **Figure 9d** shows the Raman mapping wherein red refers to the dominant CIGS, green to Cu2-xSe and blue to In2 Se3 phase. The map of annealed films contains only CIGS phase, thereby indicating the absence of any secondary phases. It also affirms the fact that the elements undergo interdiffusion during the annealing and form the desired chalcopyrite CIGS phase.

the linear section to x-axis. In addition, N<sup>a</sup> ≈ 2.6 × 1016 cm−3, calculated using the slope of the curve. The flat-band potential and acceptor density determined herein are close to the values reported previously [46]. **Figure 10c** shows the amperometric current-time (I-t) curve of CIGS films obtained at −0.4 V by chopped light, which demonstrates the nature of photoactivity of

**Figure 10.** (a) TEM image (inset: TEM-EDS, SAED pattern), (b) Mott-Schottky analysis and (c) amperometric J-t curve of

**Figure 9.** (a) Surface morphology, (b) XRD pattern, (c) Raman spectrum and (d) surface Raman mapping of sequentially

Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films

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125

pulse electrodeposited CIGS films.

sequentially pulse electrodeposited CIGS films.

TEM analysis is used to confirm the inter-diffusion of In and to infer the CIGS phase formation. The TEM image of CIGS particles is shown in **Figure 10a**. The selected area diffraction (SAED) pattern and TEM-EDS analyses are included as an inset in **Figure 10a**. SAED shows the orientations corresponding to (112), (220), (312) and (400) of CIGS, which corroborate the observation from XRD. TEM-EDS analysis shown in **Figure 10a** confirm the presence of Cu, In Ga and Se affirming the interdiffusion of In (deposited in stage II) during annealing. **Figure 10b** shows the Mott-Schottky (1/C<sup>2</sup> vs. V) plot of CIGS thin-films in 0.5 M H<sup>2</sup> SO4 at a frequency of 10 kHz in the dark. The slope of the MS plots is negative, thereby confirming that the CIGS film is p-type. The flat band potential was found to be −0.15 V (vs. SCE) by extrapolating Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films http://dx.doi.org/10.5772/intechopen.71857 125

**4.3. Study III: sequential pulsed electrodeposition of CIGS thin-films**

Cu-Ga-Se/In thin-films are annealed in Ar atmosphere and characterized.

wherein red refers to the dominant CIGS, green to Cu2-xSe and blue to In2

and form the desired chalcopyrite CIGS phase.

**10b** shows the Mott-Schottky (1/C<sup>2</sup>

CIGS films.

124 Semiconductors - Growth and Characterization

The sequential technique has been proposed to essentially minimize In precursor to address the scarcity of In. The usage of In in the growing electronic and optoelectronic industries is very high in the form of materials such as indium doped tin oxide (ITO), CIS, CIGS, InP, InN, InGaAs, InAlAs, etc., making it one of the most scarce elements in the near future. In this context, a novel sequential PC approach is explored for the fabrication of CIGS thin-films. Deposition of Cu-Ga-Se films is carried out by optimizing the deposition voltage in the first stage followed by subsequent deposition of In in the second stage. The sequentially deposited

**Figure 8.** (a) Tauc's plots and (b) photoelectrochemical J-V characteristics of direct current and pulse electrodeposited

The optimized CuGaSe/In annealed films are noted to have a compact morphology (**Figure 9a**), which is well-suited for application as solar absorber layers, since it facilitates easier diffusion of minority charge carriers and reduces recombination. XRD and Raman analyses confirm the presence of chalcopyrite CIGS films without any undesired phases as shown in **Figure 9b** and **c**. In addition to the micro-Raman analysis of CIGS films, Raman mapping of the optimized films is performed to further verify their phase-purity. **Figure 9d** shows the Raman mapping

of annealed films contains only CIGS phase, thereby indicating the absence of any secondary phases. It also affirms the fact that the elements undergo interdiffusion during the annealing

TEM analysis is used to confirm the inter-diffusion of In and to infer the CIGS phase formation. The TEM image of CIGS particles is shown in **Figure 10a**. The selected area diffraction (SAED) pattern and TEM-EDS analyses are included as an inset in **Figure 10a**. SAED shows the orientations corresponding to (112), (220), (312) and (400) of CIGS, which corroborate the observation from XRD. TEM-EDS analysis shown in **Figure 10a** confirm the presence of Cu, In Ga and Se affirming the interdiffusion of In (deposited in stage II) during annealing. **Figure** 

of 10 kHz in the dark. The slope of the MS plots is negative, thereby confirming that the CIGS film is p-type. The flat band potential was found to be −0.15 V (vs. SCE) by extrapolating

vs. V) plot of CIGS thin-films in 0.5 M H<sup>2</sup>

Se3

SO4

phase. The map

at a frequency

**Figure 9.** (a) Surface morphology, (b) XRD pattern, (c) Raman spectrum and (d) surface Raman mapping of sequentially pulse electrodeposited CIGS films.

the linear section to x-axis. In addition, N<sup>a</sup> ≈ 2.6 × 1016 cm−3, calculated using the slope of the curve. The flat-band potential and acceptor density determined herein are close to the values reported previously [46]. **Figure 10c** shows the amperometric current-time (I-t) curve of CIGS films obtained at −0.4 V by chopped light, which demonstrates the nature of photoactivity of

**Figure 10.** (a) TEM image (inset: TEM-EDS, SAED pattern), (b) Mott-Schottky analysis and (c) amperometric J-t curve of sequentially pulse electrodeposited CIGS films.

CIGS films with a photocurrent density of ≈ 0.8 mA/cm<sup>2</sup> . The improved photoresponse of CIGS films indicate their potential for application in thin-film solar cells and photoelectrochemical hydrogen generation.

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Pulsed Electrochemical Deposition of CuInSe2 and Cu(In,Ga)Se2 Semiconductor Thin Films

Se2

## **5. Conclusions**

Electrodeposition is a versatile technique for the growth of semiconductor thin films on large areas with a very low capital investment. Considering the state of energy requirement, it is pertinent to explore this low cost technique for production of copper indium selenide (CIS) and copper indium gallium selenide (CIGS) semiconductor films which are the potential candidates for application in solar photovoltaics. The features of electrodeposition with its advanced techniques are discussed in detail while also reviewing the utilization of these for the fabrication of CIS and CIGS thin films. A state-of-the-art summary has been presented on the direct current and pulse electrodeposition of CIS and CIGS thin films detailing various approaches explored while obtaining high efficient CIGS devices. In addition, a new low cost environmental friendly pulse electrodeposition technique has been proposed for the fabrication of CIS and CIGS thin films. In conclusion, the chapter puts forward the idea to photovoltaic community to explore the economic pulse electrodeposition technique for the fabrication of high quality CIGS semiconductor thin films for application in thin films solar cells.

## **Author details**

Sreekanth Mandati1,2\*, Bulusu V. Sarada<sup>1</sup> , Suhash R. Dey<sup>2</sup> and Shrikant V. Joshi<sup>3</sup>

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

1 Center for Solar Energy Materials, International Advanced Research Center for Powder Metallurgy and New Materials (ARCI), Hyderabad, Telangana, India

2 Department of Materials Science and Metallurgical Engineering, Indian Institute of Technology Hyderabad, Sangareddy, Telangana, India

3 Department of Subtractive and Additive Manufacturing, University West, Trollhatten, Sweden

## **References**


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CIGS films with a photocurrent density of ≈ 0.8 mA/cm<sup>2</sup>

hydrogen generation.

126 Semiconductors - Growth and Characterization

**5. Conclusions**

**Author details**

Sweden

**References**

CuInSe2

Sreekanth Mandati1,2\*, Bulusu V. Sarada<sup>1</sup>

(version 50). 2017;**25**:668-676

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

Technology Hyderabad, Sangareddy, Telangana, India

Metallurgy and New Materials (ARCI), Hyderabad, Telangana, India

thin films. Thin Solid Films. 1978;**48**:67-72

films indicate their potential for application in thin-film solar cells and photoelectrochemical

Electrodeposition is a versatile technique for the growth of semiconductor thin films on large areas with a very low capital investment. Considering the state of energy requirement, it is pertinent to explore this low cost technique for production of copper indium selenide (CIS) and copper indium gallium selenide (CIGS) semiconductor films which are the potential candidates for application in solar photovoltaics. The features of electrodeposition with its advanced techniques are discussed in detail while also reviewing the utilization of these for the fabrication of CIS and CIGS thin films. A state-of-the-art summary has been presented on the direct current and pulse electrodeposition of CIS and CIGS thin films detailing various approaches explored while obtaining high efficient CIGS devices. In addition, a new low cost environmental friendly pulse electrodeposition technique has been proposed for the fabrication of CIS and CIGS thin films. In conclusion, the chapter puts forward the idea to photovoltaic community to explore the economic pulse electrodeposition technique for the fabrication

of high quality CIGS semiconductor thin films for application in thin films solar cells.

, Suhash R. Dey<sup>2</sup>

1 Center for Solar Energy Materials, International Advanced Research Center for Powder

2 Department of Materials Science and Metallurgical Engineering, Indian Institute of

3 Department of Subtractive and Additive Manufacturing, University West, Trollhatten,

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**Chapter 7**

Provisional chapter

**The Electrochemical Performance of Deposited**

The Electrochemical Performance of Deposited

**Electrochemical Capacitor Application**

surface adsorption of electrolyte ions on the electrode surface.

Electrochemical Capacitor Application

Chan Pei Yi and Siti Rohana Majid

Chan Pei Yi and Siti Rohana Majid

http://dx.doi.org/10.5772/intechopen.71957

Abstract

materials

1. Introduction

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

**Manganese Oxide-Based Film as Electrode Material for**

DOI: 10.5772/intechopen.71957

Manganese Oxide-Based Film as Electrode Material for

The transition metal oxide has been recognized as one of the promising electrode materials for electrochemical capacitor application. Due to the participation of charge transfer reactions, the capacitance offered by transition metal oxide can be higher compared to double layer capacitance. The investigation on hydrous ruthenium oxide has revealed the surface redox reactions that contributed to the wide potential window shown on cyclic voltammetry curve. Although the performance of ruthenium oxide is impressive, its toxicity has limited itself from commercial application. Manganese oxide is a pseudocapacitive material behaves similar to ruthenium oxide. It consists of various oxidation states which allow the occurrence of redox reactions. It is also environmental friendly, low cost, and natural abundant. The charge storage of manganese oxide film takes into account of the redox reactions between Mn3+ and Mn4+ and can be accounted to two mechanisms. The first one involves the intercalation/deintercalation of electrolyte ions and/or protons upon reduction/oxidation processes. The second contributor for the charge storage is due to the

Keywords: pseudocapacitors, energy storage devices, metal oxides, thin films, electrode

The electronics technologies we are granted nowadays are the results of many years' researches. From Benjamin Franklin, Alessandro Volta, Michael Faraday, and Nikola Tesla, the continuous efforts have developed the electrical knowledge which has been practicalized for the sake of living standards. The better quality of life and the transformation from agricultural-based economy into information-based economy indicating the globalization has taken place. Although the fruits of globalization are attractive, the globalization also causes unavoidable negative

> © 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 eproduction in any medium, provided the original work is properly cited.

© 2018 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 Electrochemical Performance of Deposited Manganese Oxide-Based Film as Electrode Material for Electrochemical Capacitor Application** The Electrochemical Performance of Deposited Manganese Oxide-Based Film as Electrode Material for Electrochemical Capacitor Application

DOI: 10.5772/intechopen.71957

Chan Pei Yi and Siti Rohana Majid Chan Pei Yi and Siti Rohana Majid

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/intechopen.71957

#### Abstract

[78] Dharmadasa IM, Burton RP, Simmonds M. Electrodeposition of CuInSe<sup>2</sup>

Energy Materials and Solar Cells. 2006;**90:**2191-2200

Chemistry. 2011;**C115**:19632-19639

132 Semiconductors - Growth and Characterization

CuInSe2

Cu(In,Ga)Se2

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interface in CIGS solar cells. Thin Solid Films. 2001;**387**:118-122

sor layer. Solar Energy Materials and Solar Cells. 2013;**119**:241-245

absorbers. Solar Energy Materials and Solar Cells. 2015;**143**:212-217

efficient CIGS device. Advanced Functional Materials. 2015;**25**:12-27

Photoelectrochemical characterization of CuInSe<sup>2</sup>

Surface Science. 2013;**268**:391-396

S2

plated CuIn1−xGa<sup>x</sup>

2015;**126**:558-563

cells. The Journal of Physical Chemistry C. 2011;**115**:234-240

two-electrode system for applications in multi-layer graded bandgap solar cells. Solar

low-cost solar cell: Morphology control and growth mechanism. The Journal of Physical

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[82] Duchatelet A, Sidali T, Loones N, Savidand G, Chassaing E, Lincot D. 12.4% efficient

[83] Duchatelet A, Letty E, Jaime-Ferrer S, Grand PP, Mollica F, Naghavi N. The impact of reducing the thickness of electrodeposited stacked cu/in/Ga layers on the performance

[84] Insignares-Cuello C, Oliva F, Neuschitzer M, Fontané X, Broussillou C, de Monsabert TG, Saucedo E, Ruiz C, Pérez-Rodríguez A, Izquierdo-Roca V. Advanced characterization of electrodeposition-based high efficiency solar cells: Non-destructive Raman scattering quantitative assessment of the anion chemical composition in cu (in, Ga)(S, se)<sup>2</sup>

[85] Romanyuk YE, Hagendorfer H, Stücheli P, Fuchs P, Uhl AR, Sutter-Fella CM, Werner M, Haass S, Stückelberger J, Broussillou C, Grand P-P, Bermudez V, Tiwari AN. All solution-processed chalcogenide solar cells—From single functional layers towards a 13.8%

[86] Ye H, Park HS, Akhavan VA, Goodfellow BW, Panthani MG, Korgel BA, Bard AJ.

[87] Jadhav HS, Kalubarme RS, Ahn S, Yun JH, Park C-J. Effects of duty cycle on properties of CIGS thin films fabricated by pulse-reverse electrodeposition technique. Applied

[88] Vadivel S, Srinivasan K, Murali KR. Structural and optoelectronic properties of pulse

and cu(In1−xGa<sup>x</sup>

films. Optik: International Journal for Light and Electron Optics.

)Se2

thin films for solar

of CIGS solar cells. Solar Energy Materials and Solar Cells. 2017;**162**:114-119

solar cell prepared from one step electrodeposited Cu–In–Ga oxide precur-

single crystals. Physica Status Solidi A. 1979;**56**:K137-K140

[81] Wada T, Kohara N, Nishiwaki S, Negami T. Characterization of the cu( in,Ga)Se<sup>2</sup>

layers using a

particles for a

Mo

ZnSnS4

The transition metal oxide has been recognized as one of the promising electrode materials for electrochemical capacitor application. Due to the participation of charge transfer reactions, the capacitance offered by transition metal oxide can be higher compared to double layer capacitance. The investigation on hydrous ruthenium oxide has revealed the surface redox reactions that contributed to the wide potential window shown on cyclic voltammetry curve. Although the performance of ruthenium oxide is impressive, its toxicity has limited itself from commercial application. Manganese oxide is a pseudocapacitive material behaves similar to ruthenium oxide. It consists of various oxidation states which allow the occurrence of redox reactions. It is also environmental friendly, low cost, and natural abundant. The charge storage of manganese oxide film takes into account of the redox reactions between Mn3+ and Mn4+ and can be accounted to two mechanisms. The first one involves the intercalation/deintercalation of electrolyte ions and/or protons upon reduction/oxidation processes. The second contributor for the charge storage is due to the surface adsorption of electrolyte ions on the electrode surface.

Keywords: pseudocapacitors, energy storage devices, metal oxides, thin films, electrode materials

## 1. Introduction

The electronics technologies we are granted nowadays are the results of many years' researches. From Benjamin Franklin, Alessandro Volta, Michael Faraday, and Nikola Tesla, the continuous efforts have developed the electrical knowledge which has been practicalized for the sake of living standards. The better quality of life and the transformation from agricultural-based economy into information-based economy indicating the globalization has taken place. Although the fruits of globalization are attractive, the globalization also causes unavoidable negative

© The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons © 2018 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.

Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and eproduction in any medium, provided the original work is properly cited.

consequence such as global warming. The increased concentration of greenhouse gases in the atmosphere is referred as the impact of human activities such as the burning of fossil fuel for transports and power electricity plants. It brings into our concern that an alternative energy source is required, at the same time solving the energy demand issue that caused by globalization is needed. It is therefore, necessary to improve the energy management. Energy storage system, as part of the energy management, receives tremendous attentions for this purpose.

Energy storage system is a big family comprising various energy storage devices, for example, solar cell, battery, electrochemical capacitor, and fuel cell, which can be used for specific applications, for instance, the employment of fuel cell-based system over gasoline engine in transportation [1]. By reducing the rated fuel cell power and controlling the energy loss, the fuel efficiency can be enhanced [2]. Another technique of energy management is by distributing the power onto different energy sources. For example, a hybrid system based on a polymer electrolyte membrane fuel cell and nickel-metal hydride battery has been evaluated for tramway in Spain [3]. Tsukahara and Kondo have inspected the prospective hybridizations of fuel cell with Li-ion battery and electric double layer capacitor (EDLC) to power the railway vehicles [4].

Apart from transportation, many portable electronic devices also supported by the battery, that is also an energy storage device. From nickel-cadmium battery, nickel-metal hydride battery, to Li-ion battery, the studies on batteries have never stopped. To date, there are more than 35,000 published papers regarding to lithium battery according to Web of Science. Nowadays, Li-ion battery is the most promising battery. A diagram describing charging-discharging mechanism of Li-ion battery is displayed in Figure 1.

Nowadays, the Ragone plot, which is an indication of possible development for certain energy storage devices, has included some storage devices other than batteries. For example, fuel cell, capacitor, and supercapacitor. However, it does not show all the other important properties for instance cycle stability, temperature range of operation, and energy efficiency. Therefore, the Ragone plot cannot be the only reference for evaluating the performance of an energy storage device. Nevertheless, it can be used as a source of information which is continually updated

**103 104 105**

**1 s**

http://dx.doi.org/10.5772/intechopen.71957

135

**0.01 s**

**Electrolytic capacitor**

**100 s**

The Electrochemical Performance of Deposited Manganese Oxide-Based Film as Electrode Material…

**Electrochemical capacitor**

**Power density, W kg-1**

**Flywheels**

**SMES**

An electrochemical capacitor, also known as supercapacitor, is an energy storage device whereby its electrochemical performance lies in between a conventional capacitor and a battery. It can store larger amount of energy compared to a capacitor, but lesser when compared to a battery. Nonetheless, an electrochemical capacitor has higher power density than a battery. The electrochemical capacitor is categorized into two types: electric double layer capacitor (EDLC) and pseudocapacitor. The working principle behind an EDLC is based on the electrostatic interaction. The EDLC charges/discharges are according to the adsorption/de-adsorption processes. The schematic diagram of an EDLC is depicted in

As shown, the positive and negative charges are attracted to the electrode surfaces with opposite charges respectively. The charges are thus stored by means of the electric double layer formed at the electrode surface. Therefore, material with high surface area and conductivity such as carbon materials are suitable to be employed as electrode for EDLC. Conversely, a pseudocapacitor makes use of the charge transfer reactions for charge storage purpose. The

Although the charge storage mechanisms are different, the electrochemical signatures of EDLC and pseudocapacitor are similar when they are examined using cyclic voltammetry test [9]. As

corresponding capacitance is named as pseudocapacitance.

**101**

**101 102**

**Batteries**

**102**

**103**

**Energy density, J kg-1**

**104**

**105**

and improved [8].

Figure 2. Ragone plot.

Figure 3.

As a rechargeable battery, Li-ion battery encourages a rapid growth in the mobile devices technology. Abraham has outlined the limitation and future outlook of Li-ion battery [5]. Li-O2 or Li-air battery will be the battery that can achieve the utmost energy density. However, a downto-earth Li-air battery is still well on the way to be practical. Recently, an all-solid-state rechargeable battery based on a solid glass electrolyte promises a safe, low cost, and longer cycle life properties for battery [6]. The performances of batteries are summarized in Ragone plot [7], Figure 2.

Figure 1. Charging-discharging schematic diagram of li-ion rechargeable battery.

The Electrochemical Performance of Deposited Manganese Oxide-Based Film as Electrode Material… http://dx.doi.org/10.5772/intechopen.71957 135

Figure 2. Ragone plot.

consequence such as global warming. The increased concentration of greenhouse gases in the atmosphere is referred as the impact of human activities such as the burning of fossil fuel for transports and power electricity plants. It brings into our concern that an alternative energy source is required, at the same time solving the energy demand issue that caused by globalization is needed. It is therefore, necessary to improve the energy management. Energy storage system,

Energy storage system is a big family comprising various energy storage devices, for example, solar cell, battery, electrochemical capacitor, and fuel cell, which can be used for specific applications, for instance, the employment of fuel cell-based system over gasoline engine in transportation [1]. By reducing the rated fuel cell power and controlling the energy loss, the fuel efficiency can be enhanced [2]. Another technique of energy management is by distributing the power onto different energy sources. For example, a hybrid system based on a polymer electrolyte membrane fuel cell and nickel-metal hydride battery has been evaluated for tramway in Spain [3]. Tsukahara and Kondo have inspected the prospective hybridizations of fuel cell with Li-ion

Apart from transportation, many portable electronic devices also supported by the battery, that is also an energy storage device. From nickel-cadmium battery, nickel-metal hydride battery, to Li-ion battery, the studies on batteries have never stopped. To date, there are more than 35,000 published papers regarding to lithium battery according to Web of Science. Nowadays, Li-ion battery is the most promising battery. A diagram describing charging-discharging mechanism of

As a rechargeable battery, Li-ion battery encourages a rapid growth in the mobile devices technology. Abraham has outlined the limitation and future outlook of Li-ion battery [5]. Li-O2 or Li-air battery will be the battery that can achieve the utmost energy density. However, a downto-earth Li-air battery is still well on the way to be practical. Recently, an all-solid-state rechargeable battery based on a solid glass electrolyte promises a safe, low cost, and longer cycle life properties for battery [6]. The performances of batteries are summarized in Ragone plot [7],

**Load**

**Electrons Electrons**

**Li+ Charging ions Discharging**

**Electrolyte**

**Cathode Separator Anode**

as part of the energy management, receives tremendous attentions for this purpose.

battery and electric double layer capacitor (EDLC) to power the railway vehicles [4].

Li-ion battery is displayed in Figure 1.

134 Semiconductors - Growth and Characterization

**Charger**

**Cathode Separator Anode**

Figure 1. Charging-discharging schematic diagram of li-ion rechargeable battery.

Figure 2.

Nowadays, the Ragone plot, which is an indication of possible development for certain energy storage devices, has included some storage devices other than batteries. For example, fuel cell, capacitor, and supercapacitor. However, it does not show all the other important properties for instance cycle stability, temperature range of operation, and energy efficiency. Therefore, the Ragone plot cannot be the only reference for evaluating the performance of an energy storage device. Nevertheless, it can be used as a source of information which is continually updated and improved [8].

An electrochemical capacitor, also known as supercapacitor, is an energy storage device whereby its electrochemical performance lies in between a conventional capacitor and a battery. It can store larger amount of energy compared to a capacitor, but lesser when compared to a battery. Nonetheless, an electrochemical capacitor has higher power density than a battery. The electrochemical capacitor is categorized into two types: electric double layer capacitor (EDLC) and pseudocapacitor. The working principle behind an EDLC is based on the electrostatic interaction. The EDLC charges/discharges are according to the adsorption/de-adsorption processes. The schematic diagram of an EDLC is depicted in Figure 3.

As shown, the positive and negative charges are attracted to the electrode surfaces with opposite charges respectively. The charges are thus stored by means of the electric double layer formed at the electrode surface. Therefore, material with high surface area and conductivity such as carbon materials are suitable to be employed as electrode for EDLC. Conversely, a pseudocapacitor makes use of the charge transfer reactions for charge storage purpose. The corresponding capacitance is named as pseudocapacitance.

Although the charge storage mechanisms are different, the electrochemical signatures of EDLC and pseudocapacitor are similar when they are examined using cyclic voltammetry test [9]. As

Figure 3. Schematic diagram of EDLC.

the charge stored is linearly dependent with the potential, the resultant voltammetry curve exhibits a rectangular and symmetry shape. The redox materials such as conducting polymers and transition metal oxides are usually studied as the potential electrode material for pseudocapacitor. The hydrous ruthenium oxide is found to exhibit a high specific capacitance of 720 F g<sup>1</sup> , making it a promising electrode material for electrochemical capacitor [10]. Nevertheless, ruthenium oxide is expensive and toxic. This has prevented hydrous ruthenium oxide from commercial use. Besides, the usage of strong acidic electrolyte also forbids it from application. Other materials such as SnO2, MnO2, TiO2, VO2, and MoO3 have been studied as the alternative electrode material and they are showing different potentials for the practical applications [11–15]. Among the materials, manganese oxide has been widely studied due to the fact that the manganese oxide can form a variety of composites with different materials using various synthesis routes [16–18]. In the following section, we will discuss about the properties of manganese oxide.

The octahedral structure can be arranged through edge- and/or corner-sharing. In common, there are two types of structure: (a) tunnel or chain structure and (b) layer structure. Chain structure is made up of corner-sharing arrangement while tunnel structure is resulted from the combination of single, double, or triple chains of MnO6 octahedral. Layer structure is constructed by the sheets or layers stacking of MnO6 octahedral. Both kinds of structure can

**Mn O**

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As a transition metal, manganese exists in various valence states which in turns form a variety of mineral with distinct physical and chemical properties. When different phases of manganese oxide are mixed, they can intergrow to form a new structure. The manganese oxides mineral is inclusive of MnO2, Mn2O3, Mn3O4, and MnOOH. The different forms of minerals are named as the polymorphs. For instance, MnO2 has three polymorphs: pyrolusite, ramsdellite, and nautile while MnOOH has polymorphs of manganite, groutite, and feiknechtite. The polymorphic form is dependent on the linkage between MnO6 octahedrals, which is determined by the preparation process of manganese oxide [27]. Over the years, the variation in structural forms and properties has made manganese oxide a suitable candidate for different applications such as energy storage system, biosensor, coating, environmental, and nuclear science [28–34]. According to Pourbaix diagram, Figure 5, the electrochemical performances of manganese oxides in aqueous media are arisen from its various oxidation states (Mn2+, Mn3+, and Mn4+) that are emerged from different oxide phases (Mn3O4, Mn2O3, and MnO2) which are thermodynamically stable in

In general, the charge storage of manganese oxide relies on two mechanisms: (a) surface adsorption of electrolyte ions and (b) intercalation/deintercalation of electrolyte ions and/ or proton upon reduction/oxidation [36]. Although higher surface area can lead to higher value of specific capacitance, further increasing in surface area is found to not contribute

hold proton and electrolyte cation.

Figure 4. Schematic diagram of MnO6 octahedral.

alkaline region [35].

#### 2. Manganese oxide

Manganese is one of the most abundant elements on Earth which is widely distributed across the crust. In general, manganese deposit forms from seawater. Thus, its redox sensitivity toward oxidation state of ocean enables us to explore the ancient environmental conditions [19]. In human body, the manganese participates in carbohydrate metabolism as well as formation of bone and connective tissues. Manganese oxide has been used since ancient time. It was employed for elucidation of glass and acted as black pigment. Throughout the years, the researches on manganese oxide have deepen the understanding on the chemistry of this compound [20–26]. At the same time, the utilization of manganese oxide has been exploited and diversified.

The advantage of manganese oxide lies in the feasible formation of various structures based on different arrangements of basic building structure of manganese oxide, which is MnO6 octahedral. This is a structure where O2 ions are octahedrically coordinated to the central of Mn4+ ion, as shown in Figure 4.

Figure 4. Schematic diagram of MnO6 octahedral.

the charge stored is linearly dependent with the potential, the resultant voltammetry curve exhibits a rectangular and symmetry shape. The redox materials such as conducting polymers and transition metal oxides are usually studied as the potential electrode material for pseudocapacitor. The hydrous ruthenium oxide is found to exhibit a high specific capacitance

Nevertheless, ruthenium oxide is expensive and toxic. This has prevented hydrous ruthenium oxide from commercial use. Besides, the usage of strong acidic electrolyte also forbids it from application. Other materials such as SnO2, MnO2, TiO2, VO2, and MoO3 have been studied as the alternative electrode material and they are showing different potentials for the practical applications [11–15]. Among the materials, manganese oxide has been widely studied due to the fact that the manganese oxide can form a variety of composites with different materials using various synthesis routes [16–18]. In the following section, we will discuss about the

Manganese is one of the most abundant elements on Earth which is widely distributed across the crust. In general, manganese deposit forms from seawater. Thus, its redox sensitivity toward oxidation state of ocean enables us to explore the ancient environmental conditions [19]. In human body, the manganese participates in carbohydrate metabolism as well as formation of bone and connective tissues. Manganese oxide has been used since ancient time. It was employed for elucidation of glass and acted as black pigment. Throughout the years, the researches on manganese oxide have deepen the understanding on the chemistry of this compound [20–26]. At the same time, the utilization of manganese oxide has been exploited

The advantage of manganese oxide lies in the feasible formation of various structures based on different arrangements of basic building structure of manganese oxide, which is MnO6 octahedral. This is a structure where O2 ions are octahedrically coordinated to the central of Mn4+

, making it a promising electrode material for electrochemical capacitor [10].

of 720 F g<sup>1</sup>

properties of manganese oxide.

Figure 3. Schematic diagram of EDLC.

136 Semiconductors - Growth and Characterization

2. Manganese oxide

and diversified.

ion, as shown in Figure 4.

The octahedral structure can be arranged through edge- and/or corner-sharing. In common, there are two types of structure: (a) tunnel or chain structure and (b) layer structure. Chain structure is made up of corner-sharing arrangement while tunnel structure is resulted from the combination of single, double, or triple chains of MnO6 octahedral. Layer structure is constructed by the sheets or layers stacking of MnO6 octahedral. Both kinds of structure can hold proton and electrolyte cation.

As a transition metal, manganese exists in various valence states which in turns form a variety of mineral with distinct physical and chemical properties. When different phases of manganese oxide are mixed, they can intergrow to form a new structure. The manganese oxides mineral is inclusive of MnO2, Mn2O3, Mn3O4, and MnOOH. The different forms of minerals are named as the polymorphs. For instance, MnO2 has three polymorphs: pyrolusite, ramsdellite, and nautile while MnOOH has polymorphs of manganite, groutite, and feiknechtite. The polymorphic form is dependent on the linkage between MnO6 octahedrals, which is determined by the preparation process of manganese oxide [27]. Over the years, the variation in structural forms and properties has made manganese oxide a suitable candidate for different applications such as energy storage system, biosensor, coating, environmental, and nuclear science [28–34]. According to Pourbaix diagram, Figure 5, the electrochemical performances of manganese oxides in aqueous media are arisen from its various oxidation states (Mn2+, Mn3+, and Mn4+) that are emerged from different oxide phases (Mn3O4, Mn2O3, and MnO2) which are thermodynamically stable in alkaline region [35].

In general, the charge storage of manganese oxide relies on two mechanisms: (a) surface adsorption of electrolyte ions and (b) intercalation/deintercalation of electrolyte ions and/ or proton upon reduction/oxidation [36]. Although higher surface area can lead to higher value of specific capacitance, further increasing in surface area is found to not contribute

3. General principle of electrodeposition

A common configuration of electrodeposition is displayed in Figure 6.

referred to SHE and the corresponding values are recorded in Figure 7.

**Working electrode** 

**(WE)**

**Deposition electrolyte**

Figure 6. General electrodeposition setup.

Electrodeposition refers to an electrical process such as electrolytic and electrophoretic deposition, which allows the accumulated mass of a metal ions, or deposit, coated onto an electrode.

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It is usually made up of a working electrode, counter electrode, and reference electrode. The working electrode is the substrate where the deposition reaction takes place. The reference electrode is used to maintain the voltage stability for the working electrode while the counter electrode or auxiliary electrode is utilized to complete the current flow. There is various reference electrodes served for different electrolyte solutions. A standard hydrogen electrode (SHE) consists of 1.0 M H+ (aq.) solution and is used to compare with other reference electrodes since the standard electrode potential of hydrogen is 0 V. Saturated calomel electrode (SCE) is a reference electrode composes of KCl solution and establishes based on the reaction between elemental mercury and mercury chloride. However, the dangerous nature of mercury content has prohibited the SCE reference electrode from widely use. Instead, the silver chloride (Ag/AgCl) electrode is employed. Ag/AgCl reference electrode is mostly utilized for electrochemical tests and industrial application due its simplicity for fabrication, stability, and non-toxicity. Nonetheless, the usage of Ag/AgCl in basic solution must be heedful because a long immersion of electrode allows the diffusion of OH ions into the internal filling solution, which can subsequently increase the pH and shift the reaction potential. The measurement of potential using various reference electrodes can be

> **Reference electrode**

> > **Counter electrode**

**(WE)**

Figure 5. Pourbaix diagram of manganese.

to the specific capacitance [12, 37]. The structure of the manganese oxide can determine its electrochemical performance. A structure with more rooms for the insertion of electrolyte ions will offer higher charge storage capacity, and subsequently higher specific capacitance. To date, the researchers are still working best to figure out the charge storage mechanism of manganese oxide in order to better explain its electrochemical behavior [38–40]. This knowledge will inspire us and provide us a way on how to fully utilize the potential capacitance of manganese oxide.

Manganese oxide-based film can be prepared via a variety of methods. A material will exhibit different properties such as particle sizes and types of defects, depending on the fabrication routes. In addition to this, the change in experimental parameters also produces materials with different electrochemical properties. The common preparation methods are the hydrothermal, chemical bath deposition, polyol synthesis, sol–gel, electrodeposition, solvothermal, and co-precipitation [20, 36, 39, 41–45]. The specific capacitances obtained ranges from 121.5 to 698 F g<sup>1</sup> , which is still lower than the theoretical specific capacitance of manganese oxide (1380 F g<sup>1</sup> ). To better utilize the electrochemical active sites, one should understand the relationship between the experimental method and the corresponding structure formed. This chapter will focus on the green and relatively simple method of electrodeposition. Thereafter, we will discuss about the electrochemical performance of manganese oxide-based film fabricated using electrodeposition technique.

## 3. General principle of electrodeposition

Electrodeposition refers to an electrical process such as electrolytic and electrophoretic deposition, which allows the accumulated mass of a metal ions, or deposit, coated onto an electrode. A common configuration of electrodeposition is displayed in Figure 6.

It is usually made up of a working electrode, counter electrode, and reference electrode. The working electrode is the substrate where the deposition reaction takes place. The reference electrode is used to maintain the voltage stability for the working electrode while the counter electrode or auxiliary electrode is utilized to complete the current flow. There is various reference electrodes served for different electrolyte solutions. A standard hydrogen electrode (SHE) consists of 1.0 M H+ (aq.) solution and is used to compare with other reference electrodes since the standard electrode potential of hydrogen is 0 V. Saturated calomel electrode (SCE) is a reference electrode composes of KCl solution and establishes based on the reaction between elemental mercury and mercury chloride. However, the dangerous nature of mercury content has prohibited the SCE reference electrode from widely use. Instead, the silver chloride (Ag/AgCl) electrode is employed. Ag/AgCl reference electrode is mostly utilized for electrochemical tests and industrial application due its simplicity for fabrication, stability, and non-toxicity. Nonetheless, the usage of Ag/AgCl in basic solution must be heedful because a long immersion of electrode allows the diffusion of OH ions into the internal filling solution, which can subsequently increase the pH and shift the reaction potential. The measurement of potential using various reference electrodes can be referred to SHE and the corresponding values are recorded in Figure 7.

Figure 6. General electrodeposition setup.

to the specific capacitance [12, 37]. The structure of the manganese oxide can determine its electrochemical performance. A structure with more rooms for the insertion of electrolyte ions will offer higher charge storage capacity, and subsequently higher specific capacitance. To date, the researchers are still working best to figure out the charge storage mechanism of manganese oxide in order to better explain its electrochemical behavior [38–40]. This knowledge will inspire us and provide us a way on how to fully utilize the

**Mn2+**

**0 2 4 6 8 10 pH**

**Mn (s)**

**MnO2 (s)**

**Mn(OH)2 (s)**

**Mn3O4 (s)**

**MnO4** -

**Mn2O3 (s)**

Manganese oxide-based film can be prepared via a variety of methods. A material will exhibit different properties such as particle sizes and types of defects, depending on the fabrication routes. In addition to this, the change in experimental parameters also produces materials with different electrochemical properties. The common preparation methods are the hydrothermal, chemical bath deposition, polyol synthesis, sol–gel, electrodeposition, solvothermal, and co-precipitation [20, 36, 39, 41–45]. The specific capaci-

active sites, one should understand the relationship between the experimental method and the corresponding structure formed. This chapter will focus on the green and relatively simple method of electrodeposition. Thereafter, we will discuss about the electrochemical performance of manganese oxide-based film fabricated using electrodeposition

, which is still lower than the theoretical

). To better utilize the electrochemical

potential capacitance of manganese oxide.

Figure 5. Pourbaix diagram of manganese.

**0**

**E / V**

138 Semiconductors - Growth and Characterization

**-0.5**

**-1.0**

**-1.5**

**0.5**

**1.0**

**1.5**

tances obtained ranges from 121.5 to 698 F g<sup>1</sup>

technique.

specific capacitance of manganese oxide (1380 F g<sup>1</sup>

Due to this reason, the mass loading plays an important role in determining the capacity of manganese oxide formed. Among these precursors, Mn(CH3COO)2 decomposes at lower potential and offers higher deposition rate which make it a favorable precursor. In addition, the Mn3+ is more stable with acetate compared to sulfate which makes it a thermodynamic favor for oxidation kinetics [53]. On the other hand, KMnO4 is usually the Mn7+ precursor for

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As the oxidation process can lead to the dissolution of metal substrate, the cathodic reduction is usually preferred over anodic oxidation as various metals can be employed as the substrates [51]. The anodic oxidation of Mn2+ involves other oxidation state of manganese ions. Initially the Mn2+ will diffuse and adsorb onto the electrode surface to form Mnads2+. The adsorbed ions oxidize to Mn3+ which subsequently forms intermediate with water. Under appropriate heating temperature, the MnOOH intermediate can be transformed to MnO2. The detailed oxidation reaction of Mn2+ to MnO2 in the medium with different acidity is described in Section 4 [54, 55]. According to Pourbaix diagram, the reduction process of Mn7+ occurs in a neutral aqueous solution is as

The electrodeposition can be carried out either by a constant potential or current technique. For constant potential technique, the resulted current-time transient can be determined by following factors: (1) the potential and time of the oxide formation, (2) the potential and time of oxide reduction, and (3) the maintenance time of electrode at reversible potential before reduction takes place [56]. The chronopotentiometry is an electrodeposition technique performed with a constant current. The constant current is applied between working and auxiliary electrodes while the potential of working electrode is measured against a reference electrode. The potential reveals the reaction takes place on the electrode during the electrodeposition process. Before the electrodeposition starts, there is no concentration gradient of oxidants in the solution regardless of the distance from electrode surface. This initial potential is also called as open circuit potential. However, a concentration gradient of oxidants is formed when the reduction initiates by the application of current. The reduction is a resultant process of oxidants responds to the applied current. As a result, the oxidants near to the electrode surface are consumed, causing the oxidants diffuse from bulk solution toward electrode surface in order to accommodate the reduction process. The potential is comparable to the redox potential of certain electron transfer reaction. Since the concentration changes with time, this potential alters correspondingly. Once the current cannot be further sustained by this redox potential, which is due to the concentration of oxidants turns to zero, this potential will adjust to another redox potential in order to maintain the fixed current. The changing potential during galvanostatic deposition can lead to the formation of various morphologies. Different from chronoamperometry, which is a deposition method employing a constant voltage, a desired morphology can be formed accordingly. Knowledge of the interactions between electrode and electrodeposition parameters allows us to construct and improve the electrode film performance. Herein, the manganese oxide deposited using various deposition techniques are

<sup>4</sup> þ 2H2O þ 3e� ! MnO2 þ 4OH� (1)

cathodic deposition.

follow, Reaction 1 [35]:

discussed and evaluated.

MnO�

Figure 7. Voltage conversion between reference electrodes.

As electrolytic deposition makes use of solution containing metal ions for deposition purpose, electrophoretic deposition utilizes a mixture consists of suspended colloidal particles. Electrophoretic deposition was revealed when a Russian scientist observed the movement of clay particles in water induced by electric field. It occurs when the non-conductive electrically charged particles migrate to the electrode surface under an applied electric field. The charged particles suspend in the electrolyte due to the mutual electrostatic repulsion [46]. Inside the suspension, the surface charges attract the electrolyte ions with opposite charge. As a result, the counter ions adsorb onto the surface charges, forming a diffuse cloud of counter ions. The process is governed by electrostatic interaction. At the same time, the adsorbed ions repulse from each other. When they are close enough to overcome the electrostatic force, van der Waals attraction will be predominant and adhesion can occur. To achieve electrophoretic deposition, a stable suspension containing well-dispersed particles with desired electrophoretic mobility must be prepared firstly. The usage of additives such as dopamine and triethanolamine can help to stabilize the suspension [47–48]. In addition, it offers a more uniform and adherent deposit. The parameters that play the role on electrophoretic deposition are composition of dispersion medium, pH of electrolyte, and concentrations of particles and electrolyte. The principles behind electrolytic deposition and electrophoretic deposition have been studied [49–51]. An understanding about these two processes is required in order to fabricate desired electrode material. The oxidation kinetics of manganese oxide from Mn2+ is reliant on the deposition methods. Depends on the particle charge, electrodeposition can be divided into two types: anodic deposition and cathodic deposition. The anodic deposition is resulted from the oxidation of negative ions on anode (positively charged electrode). In contrast, the reduction of positive ions on cathode results in a cathodic deposition. For manganese oxide film, the anodic deposition involves the oxidation of cationic Mn2+ precursors while cathodic deposition is achieved by the reduction of anionic Mn7+ from MnO4 . Manganese (II) sulfate (MnSO4), manganese (II) nitrate (Mn (NO3)2), manganese (II) acetate (Mn(CH3COO)2), and manganese (II) chloride (MnCl2) can be used as precursors for Mn2+. As the deposition mechanism is independent of precursors, it does not affect significantly on the capacitive behaviors of manganese oxides formed [52].

Due to this reason, the mass loading plays an important role in determining the capacity of manganese oxide formed. Among these precursors, Mn(CH3COO)2 decomposes at lower potential and offers higher deposition rate which make it a favorable precursor. In addition, the Mn3+ is more stable with acetate compared to sulfate which makes it a thermodynamic favor for oxidation kinetics [53]. On the other hand, KMnO4 is usually the Mn7+ precursor for cathodic deposition.

As the oxidation process can lead to the dissolution of metal substrate, the cathodic reduction is usually preferred over anodic oxidation as various metals can be employed as the substrates [51]. The anodic oxidation of Mn2+ involves other oxidation state of manganese ions. Initially the Mn2+ will diffuse and adsorb onto the electrode surface to form Mnads2+. The adsorbed ions oxidize to Mn3+ which subsequently forms intermediate with water. Under appropriate heating temperature, the MnOOH intermediate can be transformed to MnO2. The detailed oxidation reaction of Mn2+ to MnO2 in the medium with different acidity is described in Section 4 [54, 55]. According to Pourbaix diagram, the reduction process of Mn7+ occurs in a neutral aqueous solution is as follow, Reaction 1 [35]:

As electrolytic deposition makes use of solution containing metal ions for deposition purpose, electrophoretic deposition utilizes a mixture consists of suspended colloidal particles. Electrophoretic deposition was revealed when a Russian scientist observed the movement of clay particles in water induced by electric field. It occurs when the non-conductive electrically charged particles migrate to the electrode surface under an applied electric field. The charged particles suspend in the electrolyte due to the mutual electrostatic repulsion [46]. Inside the suspension, the surface charges attract the electrolyte ions with opposite charge. As a result, the counter ions adsorb onto the surface charges, forming a diffuse cloud of counter ions. The process is governed by electrostatic interaction. At the same time, the adsorbed ions repulse from each other. When they are close enough to overcome the electrostatic force, van der Waals attraction will be predominant and adhesion can occur. To achieve electrophoretic deposition, a stable suspension containing well-dispersed particles with desired electrophoretic mobility must be prepared firstly. The usage of additives such as dopamine and triethanolamine can help to stabilize the suspension [47–48]. In addition, it offers a more uniform and adherent deposit. The parameters that play the role on electrophoretic deposition are composition of dispersion medium, pH of electrolyte, and concentrations of particles and electrolyte. The principles behind electrolytic deposition and electrophoretic deposition have been studied [49–51]. An understanding about these two processes is required in order to fabricate desired electrode material. The oxidation kinetics of manganese oxide from Mn2+ is reliant on the deposition methods. Depends on the particle charge, electrodeposition can be divided into two types: anodic deposition and cathodic deposition. The anodic deposition is resulted from the oxidation of negative ions on anode (positively charged electrode). In contrast, the reduction of positive ions on cathode results in a cathodic deposition. For manganese oxide film, the anodic deposition involves the oxidation of cationic Mn2+ precursors while cathodic deposition is achieved by the reduction

**Potential / V vs. SHE** 

**Silver Chloride Electrode (Ag/AgCl / Sat. KCl) + 0.197 V**

> **Saturated Calomel Electrode (SCE / Sat. KCl) + 0.241 V**

**Silver Chloride Electrode (Ag/AgCl / 3M KCl) + 0.210 V**

**0.0 V 0.1 V 0.2 V**

**Standard Hydrogen Electrode (SHE)** 

140 Semiconductors - Growth and Characterization

Figure 7. Voltage conversion between reference electrodes.

. Manganese (II) sulfate (MnSO4), manganese (II) nitrate (Mn

(NO3)2), manganese (II) acetate (Mn(CH3COO)2), and manganese (II) chloride (MnCl2) can be used as precursors for Mn2+. As the deposition mechanism is independent of precursors, it does not affect significantly on the capacitive behaviors of manganese oxides formed [52].

of anionic Mn7+ from MnO4

$$\rm{MnO\_4^- + 2H\_2O + 3e^- \to MnO\_2 + 4OH^-} \tag{1}$$

The electrodeposition can be carried out either by a constant potential or current technique. For constant potential technique, the resulted current-time transient can be determined by following factors: (1) the potential and time of the oxide formation, (2) the potential and time of oxide reduction, and (3) the maintenance time of electrode at reversible potential before reduction takes place [56]. The chronopotentiometry is an electrodeposition technique performed with a constant current. The constant current is applied between working and auxiliary electrodes while the potential of working electrode is measured against a reference electrode. The potential reveals the reaction takes place on the electrode during the electrodeposition process. Before the electrodeposition starts, there is no concentration gradient of oxidants in the solution regardless of the distance from electrode surface. This initial potential is also called as open circuit potential. However, a concentration gradient of oxidants is formed when the reduction initiates by the application of current. The reduction is a resultant process of oxidants responds to the applied current. As a result, the oxidants near to the electrode surface are consumed, causing the oxidants diffuse from bulk solution toward electrode surface in order to accommodate the reduction process. The potential is comparable to the redox potential of certain electron transfer reaction. Since the concentration changes with time, this potential alters correspondingly. Once the current cannot be further sustained by this redox potential, which is due to the concentration of oxidants turns to zero, this potential will adjust to another redox potential in order to maintain the fixed current. The changing potential during galvanostatic deposition can lead to the formation of various morphologies. Different from chronoamperometry, which is a deposition method employing a constant voltage, a desired morphology can be formed accordingly. Knowledge of the interactions between electrode and electrodeposition parameters allows us to construct and improve the electrode film performance. Herein, the manganese oxide deposited using various deposition techniques are discussed and evaluated.

### 4. Electrochemical performance of deposited manganese oxide-based film

#### 4.1. Effect of electrolyte composition

For anodic deposition, manganese acetate (Mn(CH3COO)2) and manganese sulfate (MnSO4) are always chosen as the Mn2+ precursors. They offer different kinds of morphologies although the electrodeposition is carried out under the same conditions. For example, the morphology produced from 0.01 M Mn(CH3COO)2 at constant current density of 30 mA cm�<sup>2</sup> exhibits interconnected but non-continuous nanorods structure, Figure 8(a) [57]. On the other hand, 0.01 M MnSO4 leads to a continuous and homogenous nanorods structure (Figure 8(b)).

With the addition of H2SO4, a discrete crystallite of manganese oxide is formed ([58]). H2SO4 acts as a supporting electrolyte and enhances the stability of the soluble Mn3+ intermediate before further reactions [59]. This allows more manganese ions to be deposited and form a film on the substrate after certain electrodeposition time. The electrodeposition mechanism of manganese oxide with the presence of H2SO4 is shown as following [60]:

$$\text{Mn}^{2+} \rightarrow \text{Mn}^{3+} + \text{e}^- \tag{2}$$

concentration, Mn3+ achieves relatively higher stability allowing it to experience disproportionation which subsequently forms Mn2+ and Mn4+ (Reaction 3). Mn4+ is then hydrolyzed to form MnO2 on the surface of substrate (Reaction 4). On the other hand, at lower acidity condition, Mn3+ is less stable and thus can hydrolyze easily to form MnOOH (Reaction 5). MnOOH is then converted to MnO2 under suitable annealing temperature (Reaction 6). The manganese oxide electrodeposited from the electrolyte containing H2SO4 can obtain specific capacitance as high as 5600 F g<sup>1</sup> [59]. A less acidic medium also can be achieved by adding sodium sulfate (Na2SO4). It

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143

In an electrolyte of Mn(CH3COO)2, 0.5 V (vs. SCE) of anodic potential forms a manganese oxide consists of two oxidation states: trivalent (Mn3+) and tetravalent states (Mn4+), while divalent ions Mn2+ and Mn3+ are observed for manganese oxide formed at anodic potential lower than 0.5 V [63, 64]. The manganese oxide resulted from deposition potential of 0.2 V (vs. SCE) for 1 hour consists of γ-Mn2O3 and/or Mn3O4 while the phase transforms from tetragonal to hexagonal system when the deposition potential increased to 0.4 V and higher potentials. At anodic potential higher than 0.5 V, Mn4+ state becomes dominant and hydrous MnO2 can be

The anodic potential is found to be correlated to the adsorbed water content in the manganese oxide structure. The increase in anodic deposition potential leads to the formation of higher oxidation state of manganese ions, at the same time results in the reduction in adsorbed water content [64, 66]. As the rate of nucleus formation increases more rapidly than the rate of crystal growth with the increase of anodic potential, the morphology preferably grows horizontally, which subsequently forms a compact and layered structure. The formation of this structure is resulted from the local fluctuation of reactants. As the reactant consumption rate increases with higher deposition potential, the reactions take place at the vicinity of electrode also occur rapidly. The release of adsorbed water molecules further resulted in the formation of an

uneven surface. The highest specific capacitance obtained is 240 F g<sup>1</sup> at 5 mV s<sup>1</sup>

(1.8 V) can result in a porous structure with nanowall architecture [68].

potential range of 0.10.4 V whereby the MnO4

specific capacitance value is achieved when the technique applied is cathodic potentiostatic electrodeposition. The MnO2 prepared through cathodic reduction process at a constant potential of 0.55 V (vs. SCE) has achieved around 233 F g<sup>1</sup> [67]. However, instead of a smooth surface which is produced at 0.5 V (vs. SCE) of anodic potential, a rough surface is formed for this cathodic potentiostatic electrodeposition. A further decrease in deposition potential

For cathodic deposition, the increase in deposition potential leads to the uniform distribution of manganese ions onto the substrate. At relatively low cathodic deposition potential, the morphology formed is a cluster-like structure. The increase of cathodic deposition potential from 0.4 V to 0.1 V motivates the formation of flaky-like structure [62]. Further increase in cathodic deposition potential forms a more homogenous and flat structure. The formation of various structures with the cathodic deposition potential arises from the different deposition mechanisms with different potential values. The first mechanism takes place in the cathodic

<sup>2</sup> produced will form H2MnO4

. A similar

that

can act as the supporting electrolyte for Mn2+ and Mn7+ precursors [61, 62].

4.2. Effect of deposition potential

formed [65].

Disproportionation pathway:

$$\text{M}\mathbf{M}\mathbf{n}^{3+} \rightarrow \mathbf{M}\mathbf{n}^{2+} + \mathbf{M}\mathbf{n}^{4+} \tag{3}$$

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

Hydrolysis pathway:

$$\text{Mn}^{3+} + 2\text{H}\_2\text{O} \rightarrow \text{MnOOH} + 3\text{H}^+ \tag{5}$$

$$\text{MnOOH} \rightarrow \text{MnO}\_2 + \text{H}^+ + \text{e}^- \tag{6}$$

Depending on the concentration of acid, there are two proposed pathways for the formation of manganese oxide: disproportionation and hydrolysis. Reaction 2 shows the oxidation of Mn2+ become soluble Mn3+ intermediate. If the supporting acid, which is H2SO4 in this case, has higher

Figure 8. Formation of nanorods structure in (a) 0.01 M Mn(CH3COO)2 and (b) 0.01 M MnSO4.

concentration, Mn3+ achieves relatively higher stability allowing it to experience disproportionation which subsequently forms Mn2+ and Mn4+ (Reaction 3). Mn4+ is then hydrolyzed to form MnO2 on the surface of substrate (Reaction 4). On the other hand, at lower acidity condition, Mn3+ is less stable and thus can hydrolyze easily to form MnOOH (Reaction 5). MnOOH is then converted to MnO2 under suitable annealing temperature (Reaction 6). The manganese oxide electrodeposited from the electrolyte containing H2SO4 can obtain specific capacitance as high as 5600 F g<sup>1</sup> [59]. A less acidic medium also can be achieved by adding sodium sulfate (Na2SO4). It can act as the supporting electrolyte for Mn2+ and Mn7+ precursors [61, 62].

#### 4.2. Effect of deposition potential

4. Electrochemical performance of deposited manganese oxide-based film

For anodic deposition, manganese acetate (Mn(CH3COO)2) and manganese sulfate (MnSO4) are always chosen as the Mn2+ precursors. They offer different kinds of morphologies although the electrodeposition is carried out under the same conditions. For example, the morphology produced from 0.01 M Mn(CH3COO)2 at constant current density of 30 mA cm�<sup>2</sup> exhibits interconnected but non-continuous nanorods structure, Figure 8(a) [57]. On the other hand, 0.01 M MnSO4 leads to a continuous and homogenous nanorods structure (Figure 8(b)).

With the addition of H2SO4, a discrete crystallite of manganese oxide is formed ([58]). H2SO4 acts as a supporting electrolyte and enhances the stability of the soluble Mn3+ intermediate before further reactions [59]. This allows more manganese ions to be deposited and form a film on the substrate after certain electrodeposition time. The electrodeposition mechanism of

Depending on the concentration of acid, there are two proposed pathways for the formation of manganese oxide: disproportionation and hydrolysis. Reaction 2 shows the oxidation of Mn2+ become soluble Mn3+ intermediate. If the supporting acid, which is H2SO4 in this case, has higher

**(a) (b)**

Figure 8. Formation of nanorods structure in (a) 0.01 M Mn(CH3COO)2 and (b) 0.01 M MnSO4.

Mn<sup>2</sup><sup>þ</sup> ! Mn<sup>3</sup><sup>þ</sup> <sup>þ</sup> <sup>e</sup>� (2)

2Mn<sup>3</sup><sup>þ</sup> ! Mn<sup>2</sup><sup>þ</sup> <sup>þ</sup> Mn<sup>4</sup><sup>þ</sup> (3)

Mn<sup>4</sup><sup>þ</sup> <sup>þ</sup> 2H2O ! MnO2 <sup>þ</sup> 4H<sup>þ</sup> (4)

Mn<sup>3</sup><sup>þ</sup> <sup>þ</sup> 2H2O ! MnOOH <sup>þ</sup> 3H<sup>þ</sup> (5)

MnOOH ! MnO2 þ H<sup>þ</sup> þ e� (6)

manganese oxide with the presence of H2SO4 is shown as following [60]:

4.1. Effect of electrolyte composition

142 Semiconductors - Growth and Characterization

Disproportionation pathway:

Hydrolysis pathway:

In an electrolyte of Mn(CH3COO)2, 0.5 V (vs. SCE) of anodic potential forms a manganese oxide consists of two oxidation states: trivalent (Mn3+) and tetravalent states (Mn4+), while divalent ions Mn2+ and Mn3+ are observed for manganese oxide formed at anodic potential lower than 0.5 V [63, 64]. The manganese oxide resulted from deposition potential of 0.2 V (vs. SCE) for 1 hour consists of γ-Mn2O3 and/or Mn3O4 while the phase transforms from tetragonal to hexagonal system when the deposition potential increased to 0.4 V and higher potentials. At anodic potential higher than 0.5 V, Mn4+ state becomes dominant and hydrous MnO2 can be formed [65].

The anodic potential is found to be correlated to the adsorbed water content in the manganese oxide structure. The increase in anodic deposition potential leads to the formation of higher oxidation state of manganese ions, at the same time results in the reduction in adsorbed water content [64, 66]. As the rate of nucleus formation increases more rapidly than the rate of crystal growth with the increase of anodic potential, the morphology preferably grows horizontally, which subsequently forms a compact and layered structure. The formation of this structure is resulted from the local fluctuation of reactants. As the reactant consumption rate increases with higher deposition potential, the reactions take place at the vicinity of electrode also occur rapidly. The release of adsorbed water molecules further resulted in the formation of an uneven surface. The highest specific capacitance obtained is 240 F g<sup>1</sup> at 5 mV s<sup>1</sup> . A similar specific capacitance value is achieved when the technique applied is cathodic potentiostatic electrodeposition. The MnO2 prepared through cathodic reduction process at a constant potential of 0.55 V (vs. SCE) has achieved around 233 F g<sup>1</sup> [67]. However, instead of a smooth surface which is produced at 0.5 V (vs. SCE) of anodic potential, a rough surface is formed for this cathodic potentiostatic electrodeposition. A further decrease in deposition potential (1.8 V) can result in a porous structure with nanowall architecture [68].

For cathodic deposition, the increase in deposition potential leads to the uniform distribution of manganese ions onto the substrate. At relatively low cathodic deposition potential, the morphology formed is a cluster-like structure. The increase of cathodic deposition potential from 0.4 V to 0.1 V motivates the formation of flaky-like structure [62]. Further increase in cathodic deposition potential forms a more homogenous and flat structure. The formation of various structures with the cathodic deposition potential arises from the different deposition mechanisms with different potential values. The first mechanism takes place in the cathodic potential range of 0.10.4 V whereby the MnO4 <sup>2</sup> produced will form H2MnO4 that subsequently reduce to form manganese oxide, as shown in Reaction 7. Another mechanism occurs at cathodic potential less than 0.2 V. It takes into account of the dissolution of manganese oxide that has formed initially. This brings about the formation of Mn2+ which is then reacted with MnO4 � to produce manganese oxide, as described in Reactions 8 and 9.

$$\text{3Mn}^{2+} + \text{2MnO}\_4^{-} + \text{2H}\_2\text{O} \rightarrow \text{5MnO}\_2 + \text{4H}^+ \tag{7}$$

Although the integration with other materials shown enhancement in charge storage, the incorporation of carbon nanotubes (CNT) does not always work in this way. As the sp2 carbon basal plane has low chemical and electrochemical reactivity, the nucleation sites on the surface of CNT-manganese oxide composite film are limited. In addition, the manganese oxide particles tend to grow at the CNT's junction [71]. A flower-like morphology built up by nanosheets that are originated radially from a central is thus formed. Nevertheless, this composite indeed acquires better cycle stability due to CNT acting as a conductive backbone that reduces the dissolution of particles. There is a variety of substrates can be used for deposition. Tantalum (Ta) foil has high melting point and provides good corrosion resistance, strength, and ductility. Cotton sheet supplies flexibility and textile structure that eases a uniform coating of material. Nickel (Ni) foil contributes a good electrical conductivity while the stainless steel is inert and owns a stable passivity. As a metal substrate, Ta foil and Ni foil usually offer a good cycle stability and specific capacitance of around 413 F g<sup>1</sup> in average for CNT-manganese oxide film [72, 73]. In spite of the fact that stainless steel has poorer electrical conductivity compared to other metals, the CNT-manganese oxide film composite film formed on stainless steel was found to achieve higher specific capacitance, that is 869 F g<sup>1</sup> [74]. Meanwhile, it also has good cycle stability. Given a similar morphology, which is nanowires-structured manganese oxide coated on CNT, formed on these three substrates, the distinct specific capacitance value achieved suggests the hidden advantages of stainless steel as a substrate. Direct deposition of manganese oxide on CNT paper using potentiodynamic method gives rise to around 168 F g<sup>1</sup> [70]. It is therefore important to choose an appropriate substrate for optimal electrochemical performance. Without the CNT, the manganese oxide tends to appear in nanorod-like structure and sphere-like structure at oxidation condition and reduction condition, respectively [75]. For the manganese oxide deposited within the same potential range (0.10.4 V), a desired structure can be determined by applying different annealing temperatures. For example, 300C leads to the formation of nanotubes structure while 100C forms nanorods structure [76]. The manganese oxide-based films prepared at different scan rates during potentiodynamic depo-

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145

sition possess distinct morphologies. At scan rate lower than 100 mV s<sup>1</sup>

tance of 410 F g<sup>1</sup>

achieved around 150 F g<sup>1</sup>

.

4.3. Effect of deposition current density

forms are irregular and dense. It starts to evolve and becomes more porous when the scan rate of deposition is increased. A nanoflake structure can be formed at scan rate of 200 mV s<sup>1</sup> [77]. This structure is beneficial for charge storage purpose at which it has offered specific capaci-

Potentiostatic deposition tends to form a more compact structure compared to the galvanostatic deposition. This is due to the consistent deposition rate during potentiostatic deposition. As the potential is maintained throughout the deposition process, the deposition rate is sustained. In contrast, the potential varies during galvanostatic deposition in order to sustain the current supply. The structure formed is thus less compact and higher in surface area [78]. The difference in compactness contributes to 50% increment in specific capacitance value and 15% higher capacitance retention upon 5000 cycles for galvanostatic deposited manganese oxide film. During the galvanostatic deposition, the structure changes from irregular to regular

. The compact and non-porous structure produced at relatively scan rate

, the morphology

$$2\text{MnO}\_4^{3-} + 8\text{H}^+ \leftrightarrow \text{MnO}\_4^- + 2\text{MnO}\_2 + 4\text{H}\_2\text{O} \tag{8}$$

$$\text{H}\_2\text{MnO}\_4^- + \text{e}^- \rightarrow 2\text{MnO}\_2 + 4\text{OH}^- \tag{9}$$

Depending on the potential range, a potentiodynamic deposition can produce a nanorod-like structure with various diameters, given that a same scan rate is employed, as shown in Figure 9. Smaller potential range (0.1�0.4 V vs. SCE) leads to the formation of nanorods structure with higher diameter while wider potential range (0.1�0.6 V vs. SCE) tends to form a nanowire structure, which has relatively smaller diameter compared to nanorods [69]. Other studies that applying a small potential range (0.3�0.6 V vs. SCE) also led to the formation of nanorod structure with similar diameter, that is around 30 nm in average [70]. This suggests the tendency to form nanorod structured manganese oxide at 0.3 V of potential range in electrolyte consisting manganese acetate and sodium sulfate. The specific capacitances estimated from charge–discharge at 0.1 mA cm�<sup>2</sup> for nanorods formed in the potential ranges of (0.1�0.4) V and (0.3�0.6) V are in the range of 200 to 250 F g�<sup>1</sup> . The extension of deposition potential range from 0.3 to 0.5 V not only improves the surface area, at the same time it also enhances the specific capacitance for 44%.

Figure 9. Morphologies formed at different potential range.

Although the integration with other materials shown enhancement in charge storage, the incorporation of carbon nanotubes (CNT) does not always work in this way. As the sp2 carbon basal plane has low chemical and electrochemical reactivity, the nucleation sites on the surface of CNT-manganese oxide composite film are limited. In addition, the manganese oxide particles tend to grow at the CNT's junction [71]. A flower-like morphology built up by nanosheets that are originated radially from a central is thus formed. Nevertheless, this composite indeed acquires better cycle stability due to CNT acting as a conductive backbone that reduces the dissolution of particles. There is a variety of substrates can be used for deposition. Tantalum (Ta) foil has high melting point and provides good corrosion resistance, strength, and ductility. Cotton sheet supplies flexibility and textile structure that eases a uniform coating of material. Nickel (Ni) foil contributes a good electrical conductivity while the stainless steel is inert and owns a stable passivity. As a metal substrate, Ta foil and Ni foil usually offer a good cycle stability and specific capacitance of around 413 F g<sup>1</sup> in average for CNT-manganese oxide film [72, 73]. In spite of the fact that stainless steel has poorer electrical conductivity compared to other metals, the CNT-manganese oxide film composite film formed on stainless steel was found to achieve higher specific capacitance, that is 869 F g<sup>1</sup> [74]. Meanwhile, it also has good cycle stability. Given a similar morphology, which is nanowires-structured manganese oxide coated on CNT, formed on these three substrates, the distinct specific capacitance value achieved suggests the hidden advantages of stainless steel as a substrate. Direct deposition of manganese oxide on CNT paper using potentiodynamic method gives rise to around 168 F g<sup>1</sup> [70]. It is therefore important to choose an appropriate substrate for optimal electrochemical performance. Without the CNT, the manganese oxide tends to appear in nanorod-like structure and sphere-like structure at oxidation condition and reduction condition, respectively [75]. For the manganese oxide deposited within the same potential range (0.10.4 V), a desired structure can be determined by applying different annealing temperatures. For example, 300C leads to the formation of nanotubes structure while 100C forms nanorods structure [76]. The manganese oxide-based films prepared at different scan rates during potentiodynamic deposition possess distinct morphologies. At scan rate lower than 100 mV s<sup>1</sup> , the morphology forms are irregular and dense. It starts to evolve and becomes more porous when the scan rate of deposition is increased. A nanoflake structure can be formed at scan rate of 200 mV s<sup>1</sup> [77]. This structure is beneficial for charge storage purpose at which it has offered specific capacitance of 410 F g<sup>1</sup> . The compact and non-porous structure produced at relatively scan rate achieved around 150 F g<sup>1</sup> .

#### 4.3. Effect of deposition current density

subsequently reduce to form manganese oxide, as shown in Reaction 7. Another mechanism occurs at cathodic potential less than 0.2 V. It takes into account of the dissolution of manganese oxide that has formed initially. This brings about the formation of Mn2+ which is then

Depending on the potential range, a potentiodynamic deposition can produce a nanorod-like structure with various diameters, given that a same scan rate is employed, as shown in Figure 9. Smaller potential range (0.1�0.4 V vs. SCE) leads to the formation of nanorods structure with higher diameter while wider potential range (0.1�0.6 V vs. SCE) tends to form a nanowire structure, which has relatively smaller diameter compared to nanorods [69]. Other studies that applying a small potential range (0.3�0.6 V vs. SCE) also led to the formation of nanorod structure with similar diameter, that is around 30 nm in average [70]. This suggests the tendency to form nanorod structured manganese oxide at 0.3 V of potential range in electrolyte consisting manganese acetate and sodium sulfate. The specific capacitances estimated from charge–discharge at 0.1 mA cm�<sup>2</sup> for nanorods formed in the potential ranges of

potential range from 0.3 to 0.5 V not only improves the surface area, at the same time it also

**0.1 V 0.2 V 0.3 V 0.4 V 0.5 V 0.6 V**

**(c)**

Figure 9. Morphologies formed at different potential range.

**nanowires**

**nanorods**

**(a) (b)**

3Mn<sup>2</sup><sup>þ</sup> <sup>þ</sup> 2MnO4

H2MnO4

(0.1�0.4) V and (0.3�0.6) V are in the range of 200 to 250 F g�<sup>1</sup>

enhances the specific capacitance for 44%.

<sup>3</sup>� <sup>þ</sup> 8H<sup>þ</sup> \$ MnO4

3MnO4

� to produce manganese oxide, as described in Reactions 8 and 9.

� þ 2H2O ! 5MnO2 þ 4H<sup>þ</sup> (7)

� þ e� ! 2MnO2 þ 4OH� (9)

� þ 2MnO2 þ 4H2O (8)

. The extension of deposition

*vs***. SCE**

reacted with MnO4

144 Semiconductors - Growth and Characterization

Potentiostatic deposition tends to form a more compact structure compared to the galvanostatic deposition. This is due to the consistent deposition rate during potentiostatic deposition. As the potential is maintained throughout the deposition process, the deposition rate is sustained. In contrast, the potential varies during galvanostatic deposition in order to sustain the current supply. The structure formed is thus less compact and higher in surface area [78]. The difference in compactness contributes to 50% increment in specific capacitance value and 15% higher capacitance retention upon 5000 cycles for galvanostatic deposited manganese oxide film. During the galvanostatic deposition, the structure changes from irregular to regular and uniform structure with the increased in current density [79, 80]. When the current density exceeds the optimal value to deposit Mn4+, a clustered structure which consists of soluble Mn6+ and/or Mn7+ may forms. Other than oxidation states, the morphology also changes with current density. The evolution of morphology is easily observed using field emission scanning electron microscopy (FESEM). As the nucleation rate is directly related to the current density, a lower current density produces lower nucleation rate. As a result, there is not many nuclei formed on the substrate surface and a continuous deposit layer is hard to be constructed [81]. In addition, the deposit usually possesses rough surface. Higher current density can lead to higher nucleation and growth rate. The deposit accumulates on the structure formed ahead and filling up the space or crack, which leads to the formation of a uniform coating, Figure 10.

At optimal current density, various structures such as nanowires, nanoflakes, nanosheets, and nanorods can be formed [57, 82, 83]. For example, manganese oxide formed from MnSO4 precursor at 4 mA cm<sup>2</sup> presents as agglomerated clusters. A small decreased in current density, 3.7 mA cm<sup>2</sup> , brings about a grain-like structure constructed by nanowires [84]. 2 mA cm<sup>2</sup> of galvanostatic deposition in Mn(NO3)2 produces a flower-like structure made up of nanowires [83]. This manganese oxide film is capable to maintain 84% of specific capacitance after 1000 cycles.

It can be seen that the stability of manganese oxide suspension is important to ensure the success of electrophoretic deposition. The ethanol can be used as the liquid medium for suspension. However, the deposit resulted from the manganese oxide suspension in ethanol is shown to be irregular and tends to form agglomerates. The addition of phosphate ester has been shown to enhance the stability of the manganese oxide suspension and increase the mass load [85]. This improves the electrochemical performance of the manganese oxide film. Without the phosphate ester, the manganese oxide film only exhibits 236 F g<sup>1</sup> of specific capacitance and drops 15% of the initial specific capacitance after 25 cycles [86]. Meanwhile, the phosphate ester has offered around 60% of increment in specific capacitance [85]. Sodium alginate is a good dispersant as well. It has been proposed that the sodium alginate provides the electrostatic and steric stabilization for the manganese oxide suspension. Additionally, it supplies electric charge for the suspension particles which is beneficial for the deposition process [87]. The manganese oxide film deposited from the dispersant electrolyte of sodium

**Substrate**

**(a) (b)**

Figure 11. Schematic diagram of electrophoretic deposited manganese oxide film.

**Charged manganese oxide nanofibers**

The Electrochemical Performance of Deposited Manganese Oxide-Based Film as Electrode Material…

nanotubes or reduced graphene oxide with the manganese oxide does not alter the nanostructure of the oxide itself [88]. The nanostructured manganese oxide particles attach on the carbon

We have discussed about the impact of few deposition parameters on the deposited manganese oxide-based film. The as-deposited manganese oxide-based film is dominantly amorphous with inherent cation deficiency [90]. The defect is most likely to form at relatively low deposition temperatures (80200C). Except the one prepared using electrophoretic deposition, at which the manganese oxide particles are firstly fabricated before deposition, the as-deposited manganese oxide-based film requires certain post-treatments to improve the crystalline structure. Annealing is one of the common post-treatment. When the manganese oxide-based film undergoes post-heating process, different crystalline structure forms based on the annealing temperature. Most of the water content at the surface layer of structure desorbs gradually at 120350C [91]. The manganese oxide presents as γ-manganese oxide at annealing temperature lower than 350C. The transformation of crystalline structure is initiated at around 300C, which allows the γ-phase changes to β-phase of MnO2. When the manganese oxide undergoes further heating, α-Mn2O3 phase starts to form [92]. Any desired crystalline structure is thus can be prepared. Other than crystalline structure, the morphology

) a little higher than the one prepared from

**Manganese oxide film with nanofiber structure**

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147

). The integration of carbon

alginate has obtained specific capacitance (412 F g<sup>1</sup>

**Substrate**

the ethanol with phosphate ester dispersant electrolyte (377 F g<sup>1</sup>

nanotube surface and deposit together onto the substrate [89].

#### 4.4. Electrophoretic deposition

As one kind of the electrodeposition techniques, the oxidation and dissolution of the metal substrate during the anodic electrophoretic deposition is a concern. Thus, cathodic electrophoretic deposition is more favorable compared to its counterpart. Different from the electrodeposition techniques mentioned above, the electrophoretic deposition does not form the oxide during or after the process. Instead, the oxide is fabricated beforehand. The desired nanostructure can be manipulated using various preparation methods such as hydrothermal, chemical reduction, chemical precipitation, spray pyrolysis, and wet-chemical processes. The nanostructured oxide powder formed is then suspended in a dispersant electrolyte for further action, which is the electrophoretic deposition in this case. The structure of manganese oxide particles will be retained even after deposition, as depicted in Figure 11.

Figure 10. Morphological evolution with the increase in current density.

Figure 11. Schematic diagram of electrophoretic deposited manganese oxide film.

and uniform structure with the increased in current density [79, 80]. When the current density exceeds the optimal value to deposit Mn4+, a clustered structure which consists of soluble Mn6+ and/or Mn7+ may forms. Other than oxidation states, the morphology also changes with current density. The evolution of morphology is easily observed using field emission scanning electron microscopy (FESEM). As the nucleation rate is directly related to the current density, a lower current density produces lower nucleation rate. As a result, there is not many nuclei formed on the substrate surface and a continuous deposit layer is hard to be constructed [81]. In addition, the deposit usually possesses rough surface. Higher current density can lead to higher nucleation and growth rate. The deposit accumulates on the structure formed ahead and filling up the space or crack, which leads to the formation of a uniform coating,

At optimal current density, various structures such as nanowires, nanoflakes, nanosheets, and nanorods can be formed [57, 82, 83]. For example, manganese oxide formed from MnSO4 precursor at 4 mA cm<sup>2</sup> presents as agglomerated clusters. A small decreased in current

2 mA cm<sup>2</sup> of galvanostatic deposition in Mn(NO3)2 produces a flower-like structure made up of nanowires [83]. This manganese oxide film is capable to maintain 84% of specific

As one kind of the electrodeposition techniques, the oxidation and dissolution of the metal substrate during the anodic electrophoretic deposition is a concern. Thus, cathodic electrophoretic deposition is more favorable compared to its counterpart. Different from the electrodeposition techniques mentioned above, the electrophoretic deposition does not form the oxide during or after the process. Instead, the oxide is fabricated beforehand. The desired nanostructure can be manipulated using various preparation methods such as hydrothermal, chemical reduction, chemical precipitation, spray pyrolysis, and wet-chemical processes. The nanostructured oxide powder formed is then suspended in a dispersant electrolyte for further action, which is the electrophoretic deposition in this case. The structure of manganese oxide particles

will be retained even after deposition, as depicted in Figure 11.

**Mn2+ ion**

Figure 10. Morphological evolution with the increase in current density.

**Substrate**

**(a) (b) (c) Increase in current density**

, brings about a grain-like structure constructed by nanowires [84].

**Substrate**

**MnOx film**

Figure 10.

density, 3.7 mA cm<sup>2</sup>

**Substrate**

capacitance after 1000 cycles.

146 Semiconductors - Growth and Characterization

4.4. Electrophoretic deposition

It can be seen that the stability of manganese oxide suspension is important to ensure the success of electrophoretic deposition. The ethanol can be used as the liquid medium for suspension. However, the deposit resulted from the manganese oxide suspension in ethanol is shown to be irregular and tends to form agglomerates. The addition of phosphate ester has been shown to enhance the stability of the manganese oxide suspension and increase the mass load [85]. This improves the electrochemical performance of the manganese oxide film. Without the phosphate ester, the manganese oxide film only exhibits 236 F g<sup>1</sup> of specific capacitance and drops 15% of the initial specific capacitance after 25 cycles [86]. Meanwhile, the phosphate ester has offered around 60% of increment in specific capacitance [85]. Sodium alginate is a good dispersant as well. It has been proposed that the sodium alginate provides the electrostatic and steric stabilization for the manganese oxide suspension. Additionally, it supplies electric charge for the suspension particles which is beneficial for the deposition process [87]. The manganese oxide film deposited from the dispersant electrolyte of sodium alginate has obtained specific capacitance (412 F g<sup>1</sup> ) a little higher than the one prepared from the ethanol with phosphate ester dispersant electrolyte (377 F g<sup>1</sup> ). The integration of carbon nanotubes or reduced graphene oxide with the manganese oxide does not alter the nanostructure of the oxide itself [88]. The nanostructured manganese oxide particles attach on the carbon nanotube surface and deposit together onto the substrate [89].

We have discussed about the impact of few deposition parameters on the deposited manganese oxide-based film. The as-deposited manganese oxide-based film is dominantly amorphous with inherent cation deficiency [90]. The defect is most likely to form at relatively low deposition temperatures (80200C). Except the one prepared using electrophoretic deposition, at which the manganese oxide particles are firstly fabricated before deposition, the as-deposited manganese oxide-based film requires certain post-treatments to improve the crystalline structure. Annealing is one of the common post-treatment. When the manganese oxide-based film undergoes post-heating process, different crystalline structure forms based on the annealing temperature. Most of the water content at the surface layer of structure desorbs gradually at 120350C [91]. The manganese oxide presents as γ-manganese oxide at annealing temperature lower than 350C. The transformation of crystalline structure is initiated at around 300C, which allows the γ-phase changes to β-phase of MnO2. When the manganese oxide undergoes further heating, α-Mn2O3 phase starts to form [92]. Any desired crystalline structure is thus can be prepared. Other than crystalline structure, the morphology

The addition of secondary or ternary materials can enhance the electrochemical performance of manganese oxide film. However, these are not shown in this chapter. The core idea is to gather the advantages of various materials then compensate shortcomings of each other. The common materials combination involves the carbon material, transition metal oxide, and conducting polymer. The carbon material offers conductivity while transition metal oxide and conducting polymer provides more electroactive sites for charge storage purpose. Electrodeposition is a widely used technique to protect and strengthen the function of parts used in various industries. The wide application of electrodeposition technology can be attributed to its simplicity, manufacturability, and scalability. The electrodeposition method also requires a relatively low fabrication cost for energy storage device compared to other methods. This technique allows a direct formation of film on the substrate desired and the film properties are governable by varying the deposition parameters. The traditional method of electrode fabrication for electrochemical capacitor involves the pressing of electrode material onto the substrate. This process can increase the contact resistance and reduce the porous surface area which brings about damage to the electrode materials formed. In overall, manganese oxidebased film prepared using electrodeposition is prospective and practical. There are numerous related studies carrying out every year, although the reported specific capacitance of manganese oxide-based film are yet far from the expected performances (Table 1), we are convinced

The Electrochemical Performance of Deposited Manganese Oxide-Based Film as Electrode Material…

Deposition electrolyte Specific

2 mM MnSO4 + 50 mM KCl 163.40 [44]

0.1 M MnSO4 + 0.01 M TTAB 343.00 [100]

0.5 M KMnO4 128.00 [78]

2 mM MnSO4 + 20 mM Na2MoO4 190.90 [104]

MnOx Potentiostatic 0.2 V vs. SCE 0.001 M KMnO4 + 1 M Na2SO4 368.00 [62] MnOx Potentiostatic 0.50 V vs. SCE 0.25 M Mn(CH3COO)2 240.00 [63] MnO2 Potentiostatic 0.55 V vs. SCE KMnO4 232.94 [67]

MnOx 0.1 M MnSO4 294.00

SLS

Mn3O4 Potentiostatic 1.3 V vs. SCE 0.25 M MnNO3 416.00 [101] MnO2 Potentiostatic 0.60 vs. SCE 0.1 M Mn(CH3COO)2 + 0.1 M Na2SO4 240.00 [102]

Potentiostatic 0.75 V vs. SCE 0.25 M MnSO45H2O 285.00 [65]

0.05 M Mn(CH3COO)2 + 0.1 M Ni(CH3COO)2 + 0.2 M CH3COONa

0.5 M MnSO4 + 0.5 M Na2SO4 + 100 mM

capacitance, F g<sup>1</sup>

http://dx.doi.org/10.5772/intechopen.71957

167.50 [70]

169.00 [103]

310.00 [105]

Reference

149

that there is still a big room for improvement.

Ag/AgCl

Ag/AgCl

Ag/AgCl

0.30–0.60 V vs. SCE

0.4 1.2 V vs. Ag/AgCl

0.00 1.00 V vs. Ag/AgCl

1.10 1.50 V vs.

SCE

Deposition mode

MnOx Potentiostatic 1.0 V vs.

MnOx Potentiostatic 1.1 V vs.

MnO2 Potentiostatic 10 V vs.

voltammetry

voltammetry

Cyclic voltammetry

Cyclic voltammetry

MnO2-CNT Cyclic

MnOx Cyclic

Mn-Ni mixed oxide

Mn-Mo mixed oxide

Electrode materials

a-

MnO2nH2O

Figure 12. Morphologies of manganese oxide film formed at annealing temperature of (a) 100 200C, (b) 300C, (c) 500C, and (d) 600C.

of the manganese oxide-based film can also be modified by varying the annealing temperatures. At annealing temperatures lower than 200C, a fibrous and granular structure is formed. The entanglement of fibers causes the morphology to evolve to a cluster-like structure when the temperature reaches 300C. Further increasing in temperature can lead to the formation of flaky-like (500C) and rod-like structures (600C) [93], Figure 12.

There are still many studies carried out to investigate other possible factors affecting the deposited manganese oxide-based film. For example, by studying the porosity of the manganese oxide-based film, one can also inspect more details about the structure and charge transport properties [94]. More and more deposition techniques have been developed and studied to prepare the manganese oxide-based film. For instance, a redox deposition that is took place when a substrate is immersed substrate into the Mn ions precursors [95, 96]. From chemical bath deposition, chemical vapor deposition to spray pyrolysis deposition, all deposition techniques are intended to grow the deposit with a good quality along with good physical and chemical properties in a large scale.

#### 5. Summary

Manganese oxide-based film is shown to be prepared in various electrodeposition conditions by varying electrodeposition potentials, current densities, additives, and electrolytes. The potential application of manganese oxide-based film as the electrode material for electrochemical capacitor is thus discussed. As deposited manganese oxide film is amorphous in nature [97], the amorphousness can be transformed to crystalline phase by employing appropriate annealing temperature. The crystallinity start to arise when the annealing temperature increases to 300C. MnO2 is the first crystal structure detected at 300C, further increasing the temperature leads to the formation of Mn3O4 and Mn2O3 [98]. Since proton participates in the charge storage mechanism of manganese oxide, it is believed that this oxide will perform better in hydrous form. Previous study about the RuO2 has shown the significant improvement on the electrochemical performance with the presence of hydrous phase [99]. Not long after that, the crystalline manganese oxide has exhibited its potential application as electrode material as well [12]. To date, the charge storage mechanism is found to be dependent on the crystalline structure, water content, and surface area. It turns out that, the electrochemical performance of manganese oxide-based film is not totally relied on any of these factors. In contrast, it is resulted from the combination of all the factors that have been found and studied. For this reason, the researchers are still making effort to understand this complication.

The addition of secondary or ternary materials can enhance the electrochemical performance of manganese oxide film. However, these are not shown in this chapter. The core idea is to gather the advantages of various materials then compensate shortcomings of each other. The common materials combination involves the carbon material, transition metal oxide, and conducting polymer. The carbon material offers conductivity while transition metal oxide and conducting polymer provides more electroactive sites for charge storage purpose. Electrodeposition is a widely used technique to protect and strengthen the function of parts used in various industries. The wide application of electrodeposition technology can be attributed to its simplicity, manufacturability, and scalability. The electrodeposition method also requires a relatively low fabrication cost for energy storage device compared to other methods. This technique allows a direct formation of film on the substrate desired and the film properties are governable by varying the deposition parameters. The traditional method of electrode fabrication for electrochemical capacitor involves the pressing of electrode material onto the substrate. This process can increase the contact resistance and reduce the porous surface area which brings about damage to the electrode materials formed. In overall, manganese oxidebased film prepared using electrodeposition is prospective and practical. There are numerous related studies carrying out every year, although the reported specific capacitance of manganese oxide-based film are yet far from the expected performances (Table 1), we are convinced that there is still a big room for improvement.

of the manganese oxide-based film can also be modified by varying the annealing temperatures. At annealing temperatures lower than 200C, a fibrous and granular structure is formed. The entanglement of fibers causes the morphology to evolve to a cluster-like structure when the temperature reaches 300C. Further increasing in temperature can lead to the formation of

Figure 12. Morphologies of manganese oxide film formed at annealing temperature of (a) 100 200C, (b) 300C, (c)

**Substrate Substrate Substrate Substrate (a) (b) (c) (d)**

There are still many studies carried out to investigate other possible factors affecting the deposited manganese oxide-based film. For example, by studying the porosity of the manganese oxide-based film, one can also inspect more details about the structure and charge transport properties [94]. More and more deposition techniques have been developed and studied to prepare the manganese oxide-based film. For instance, a redox deposition that is took place when a substrate is immersed substrate into the Mn ions precursors [95, 96]. From chemical bath deposition, chemical vapor deposition to spray pyrolysis deposition, all deposition techniques are intended to grow the deposit with a good quality along with good physical

Manganese oxide-based film is shown to be prepared in various electrodeposition conditions by varying electrodeposition potentials, current densities, additives, and electrolytes. The potential application of manganese oxide-based film as the electrode material for electrochemical capacitor is thus discussed. As deposited manganese oxide film is amorphous in nature [97], the amorphousness can be transformed to crystalline phase by employing appropriate annealing temperature. The crystallinity start to arise when the annealing temperature increases to 300C. MnO2 is the first crystal structure detected at 300C, further increasing the temperature leads to the formation of Mn3O4 and Mn2O3 [98]. Since proton participates in the charge storage mechanism of manganese oxide, it is believed that this oxide will perform better in hydrous form. Previous study about the RuO2 has shown the significant improvement on the electrochemical performance with the presence of hydrous phase [99]. Not long after that, the crystalline manganese oxide has exhibited its potential application as electrode material as well [12]. To date, the charge storage mechanism is found to be dependent on the crystalline structure, water content, and surface area. It turns out that, the electrochemical performance of manganese oxide-based film is not totally relied on any of these factors. In contrast, it is resulted from the combination of all the factors that have been found and studied. For this reason, the researchers are still making effort to understand

flaky-like (500C) and rod-like structures (600C) [93], Figure 12.

and chemical properties in a large scale.

5. Summary

500C, and (d) 600C.

148 Semiconductors - Growth and Characterization

this complication.



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Table 1. The electrochemical performance of deposited manganese oxide-based film.

#### Acknowledgements

The authors acknowledge the University of Malaya for providing financial support through the projects BKS030-2017 and FG034-17AFR.

#### Author details

Chan Pei Yi and Siti Rohana Majid\*

\*Address all correspondence to: shana@um.edu.my

Centre for Ionics University of Malaya, Department of Physics, Faculty of Science, University of Malaya, Kuala Lumpur, Malaysia

### References

Acknowledgements

Electrode Materials

Fe-doped MnOx

Co-doped MnOx

Fe-doped MnOx

MnO2 Galvanostatic 10.5 mA

150 Semiconductors - Growth and Characterization

MnO2 Galvanostatic 165 mA

MnO2-RGO Electrophoretic deposition

MnOX Electrophoretic deposition

MnOx Electrophoretic deposition

MnOx-CNT Electrophoretic deposition

MnOx Electrophoretic deposition

MnOx Electrophoretic deposition

cm<sup>2</sup>

cm<sup>2</sup>

0.3 V vs. SCE

5 10 V vs. SCE

10 100 V vs. SCE

100 V vs. Ag/AgCl

100 V vs. Ag/AgCl

Table 1. The electrochemical performance of deposited manganese oxide-based film.

Author details

the projects BKS030-2017 and FG034-17AFR.

\*Address all correspondence to: shana@um.edu.my

Chan Pei Yi and Siti Rohana Majid\*

of Malaya, Kuala Lumpur, Malaysia

The authors acknowledge the University of Malaya for providing financial support through

Deposition mode Electrolyte composition Specific

MnO2 Galvanostatic 2 mA cm<sup>2</sup> 5 mM Mn(NO3)2 246.00 [83] MnO2 Galvanostatic 3 mA cm<sup>2</sup> 0.02 M KMnO4 188.00 [106] MnO2-PPy Galvanostatic 4 mA cm<sup>2</sup> 0.2 M MnSO4 + PPy 620.00 [84]

MnOx Galvanostatic 5 mA cm<sup>2</sup> 0.01 M Mn(CH3COO)2 185.00 [57]

Galvanostatic MnSO4, iron sulfate, EDTA 298.40

Galvanostatic 5 mA cm<sup>2</sup> 0.1 M MnSO4 + 0.1 M citric acid 218.00 [107]

Galvanostatic 50 mA cm<sup>2</sup> MnSO4, cobalt sulfate, EDTA 186.20 [109]

Manganese oxide + ethanol +

Manganese oxide + ethylene alcohol +

phosphate ester

H2SO4

15 V vs. SCE Manganese oxide + sodium alginate + carbon nanotubes

0.02 M Mn(CH3COO)2 201.00 [108]

0.5 M KMnO4 196.00 [78]

MnO2 + RGO 392.00 [88]

Manganese oxide + sodium alginate 412.00 [87]

Manganese oxide + ethanol 236.00 [86]

capacitance, F g<sup>1</sup>

377.00 [85]

Around 210.00 [89]

275.00 [110]

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**Chapter 8**

Provisional chapter

**Semiconducting Electrospun Nanofibers for Energy**

DOI: 10.5772/intechopen.72817

Nowadays, semiconducting thin films, thanks to their unique and excellent properties, play a crucial role for the design of devices for energy conversion and storage, such as solar cells, perovskite solar cells, lithium-ion batteries (LIBs), and fuel cells. Since the nanostructured arrangements can improve the behavior of the materials in several application fields, in this chapter we propose the electrospinning process as electrohydrodynamic deposition to obtain semiconducting materials, in the form of nanofiber mats. The nanostructured mats are able to provide high surface-area-to-volume ratio and a microporous structure, which are crucial aspects for energetic application. In this chapter, we deeply describe the electrospinning process and how nanofibers obtained can be used in energy devices, satisfying all the requirements to improve overall final performances.

Keywords: semiconducting materials, nanofiber mats, electrospinning, energy

Nowadays, different deposition techniques, based on the application of an external electric potential, are carried out in order to prepare thin film and coatings. A thin film is defined as a layer of material with a thickness in the range from few nanometers (namely monolayer) to several micrometers. All electrodeposition processes, which can be divided in chemical methods and physical methods as proposed in Figure 1, ensure the deposition of different classes of materials as metals, semiconductors, ceramics, and organo-ceramics in the form of thin films, onto several substrate materials. Semiconducting thin films show a variety of unique and excellent properties that make them particularly attractive in several application areas. Among them, these materials play a preeminent role for the design of devices for energy conversion, such as solar cells and perovskite solar cells [1]. The behavior of semiconducting materials can be improved toward energy-related applications, when their shape and dimension are controlled

> © 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 eproduction in any medium, provided the original work is properly cited.

© 2018 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.

Semiconducting Electrospun Nanofibers

**Conversion**

Abstract

1. Introduction

Giulia Massaglia and Marzia Quaglio

Giulia Massaglia and Marzia Quaglio

for Energy Conversion

http://dx.doi.org/10.5772/intechopen.72817

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

conversion devices, energy storage devices


#### **Semiconducting Electrospun Nanofibers for Energy Conversion** Semiconducting Electrospun Nanofibers for Energy Conversion

DOI: 10.5772/intechopen.72817

#### Giulia Massaglia and Marzia Quaglio Giulia Massaglia and Marzia Quaglio

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/intechopen.72817

#### Abstract

[101] Nguyen T, João Carmezim M, Boudard M, Fátima Montemor M. Cathodic electrodeposition and electrochemical response of manganese oxide pseudocapacitor electrodes.

[102] Chou S, Cheng F, Chen J. Electrodeposition synthesis and electrochemical properties of nanostructured γ-MnO2 films. Journal of Power Sources. 2006;162(1):727-734 [103] Tahmasebi MH, Raeissi K, Golozar MA, Vicenzo A, Hashempour M, Bestetti M. Tailoring the pseudocapacitive behavior of electrochemically deposited manganese-nickel

[104] Masaharu Nakayama AT, Sato Y, Tonosaki T, Ogura K. Electrodeposition of manganese and molybdenum mixed oxide thin films and their charge storage properties. Langmuir.

[105] Devaraj S, Munichandraiah N. High capacitance of electrodeposited MnO2 by the effect of a surface-active agent. Electrochemical and Solid-State Letters. 2005;8(7):A373-A377

[106] Jacob GM, Zhitomirsky I. Microstructure and properties of manganese dioxide films prepared by electrodeposition. Applied Surface Science. 2008;254(20):6671-6676

[107] Dubal DP, Kim WB, Lokhande CD. Galvanostatically deposited Fe: MnO2 electrodes for supercapacitor application. Journal of Physics and Chemistry of Solids. 2012;73(1):18-24

[108] Kundu M, Liu L. Direct growth of mesoporous MnO2 nanosheet arrays on nickel foam current collectors for high-performance pseudocapacitors. Journal of Power Sources.

[109] Trung Dung D, Thi Thu Hang L, Thi Bich Thuy H, Thanh Tung M. Synthesis of nanostructured manganese oxides based materials and application for supercapacitor. Advances in Natural Sciences: Nanoscience and Nanotechnology. 2015;6(2):025011 [110] Chen C-Y, Lyu Y-R, Su C-Y, Lin H-M, Lin C-K. Characterization of spray pyrolyzed manganese oxide powders deposited by electrophoretic deposition technique. Surface

International Journal of Hydrogen Energy. 2015;40(46):16355-16364

oxide films. Electrochimica Acta. 2016;190:636-647

and Coating Technology. 2007;202(4):1277-1281

2005;21(13):5907-5913

158 Semiconductors - Growth and Characterization

2013;243:676-681

Nowadays, semiconducting thin films, thanks to their unique and excellent properties, play a crucial role for the design of devices for energy conversion and storage, such as solar cells, perovskite solar cells, lithium-ion batteries (LIBs), and fuel cells. Since the nanostructured arrangements can improve the behavior of the materials in several application fields, in this chapter we propose the electrospinning process as electrohydrodynamic deposition to obtain semiconducting materials, in the form of nanofiber mats. The nanostructured mats are able to provide high surface-area-to-volume ratio and a microporous structure, which are crucial aspects for energetic application. In this chapter, we deeply describe the electrospinning process and how nanofibers obtained can be used in energy devices, satisfying all the requirements to improve overall final performances.

Keywords: semiconducting materials, nanofiber mats, electrospinning, energy conversion devices, energy storage devices

#### 1. Introduction

Nowadays, different deposition techniques, based on the application of an external electric potential, are carried out in order to prepare thin film and coatings. A thin film is defined as a layer of material with a thickness in the range from few nanometers (namely monolayer) to several micrometers. All electrodeposition processes, which can be divided in chemical methods and physical methods as proposed in Figure 1, ensure the deposition of different classes of materials as metals, semiconductors, ceramics, and organo-ceramics in the form of thin films, onto several substrate materials. Semiconducting thin films show a variety of unique and excellent properties that make them particularly attractive in several application areas. Among them, these materials play a preeminent role for the design of devices for energy conversion, such as solar cells and perovskite solar cells [1]. The behavior of semiconducting materials can be improved toward energy-related applications, when their shape and dimension are controlled

© The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons © 2018 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.

Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and eproduction in any medium, provided the original work is properly cited.

2. Electrospinning technique and its principles

The electrospinning is an electro-hydrodynamic process that provides polymer-based fibers with diameter distribution in the range from few nanometers to several micrometers by involving electrostatic forces [4–7]. The process is based on the concept that electrostatic forces induce columbic interactions between charged elements of the polymeric fluid, leading then to overcome the surface tension in a charged polymeric jet and ensuring the nanofiber formation. An electrospinning system is constituted by three major components, as sketched in Figure 2(a): (i) high-voltage supply; (ii) a spinneret, which represents one of the two electrodes, containing the metallic needle of the syringe, where the polymeric solution is loaded; and (iii) the counter electrode, also named grounded electrode, which is the second electrode, where the nanofibers are collected. It involves a high-voltage supply in order to inject charges with a certain polarity in the polymeric solution and then generates a polymeric charged jet, accelerated toward a counter electrode with opposite polarity. In a typical process, the voltage (0–30 kV) is applied between the first electrode (tip of needle) and the second electrode (counter electrode). This implies the indirectly definition of electric field intensity as the ratio between the voltage value and working distance. The working distance is the distance between the first electrode and the counter electrode. The spinneret is linked with a syringe, in which the polymeric (or melt) solution is loaded and a syringe pump allows to control the solution flows with a constant rate, defined as flow rate. When the voltage is applied, the drop at the tip of the needle becomes highly electrified, and the charges are uniformly distributed on its surface. Therefore, the repulsive forces, acted between all charged elements of polymeric solution, induce an elongation of the spherical drop to form a conical shape, known as Taylor's cone. When the repulsive forces

Semiconducting Electrospun Nanofibers for Energy Conversion

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161

Figure 2. A sketch of electrospinning setup is proposed in (a). In (b) a representation of bending instabilities characterizing the charged polymeric jet during the electrospinning process is proposed. (reprinted with the permission from

(polymer, 2008, 49, 2387–2425) copyright (2008) Elsevier).

Figure 1. Different deposition techniques obtained by applying an external electric field and/or potential difference. These techniques can be divided in electrodeposition methods and electro-hydrodynamic techniques, as explained in the diagram.

up to the nanometer scale, in the so-called nanostructured arrangements. In this chapter, electrohydrodynamic depositions are proposed in order to obtain semiconducting materials in the form of nanofiber mats, able to combine high surface-area-to-volume ratio together with a microporous structure well suited for energetic application. The electro-hydrodynamic techniques involve an external electric field applied to a polymeric solution to provide the final deposition of nanomaterial. As proposed in Figure 1, two different processes, i.e., electrospinning and electrospray, can be classified as electro-hydrodynamic techniques. Indeed, electrospinning ensures the direct assembly of nanofiber mats with different morphologies and properties, as described in this chapter.

Electrospinning process is based on the principle that strong repulsive forces, induced by external applied electric field, can overcome the surface tension in a charged polymeric jet [2, 3]. Therefore, through this technique, the polymer can be arranged in a mat with a high surfacearea-ratio-to-volume, showing a micro/macroporous structure. Moreover, final nanofibers are based not only on polymers but also on metals, ceramics, and metal oxides, obtained by implementing further different chemical and thermal treatments. However, electrospray technique is an electro-hydrodynamic technique, which occurs at low viscosity values of initial polymeric solution. Indeed, in this case, the surface tension overcomes the viscoelastic forces, and consequently, the instauration of charged droplets with different diameters and concentrations occurs during the process.

During the last decade, different works have been presented in the literature, focusing their attention on nanostructured semiconducting metal oxides (as TiO2, ZnO, CuO, and SnO2) in order to design well-performing and green energy systems (such as in dye-sensitized solar cells, lithium-ion batteries (LIBs), fuel cells). In this scenario, nanofibers progressively increased their importance as one of the most important nanostructures to be selected to improve the final performance of the devices.

## 2. Electrospinning technique and its principles

up to the nanometer scale, in the so-called nanostructured arrangements. In this chapter, electrohydrodynamic depositions are proposed in order to obtain semiconducting materials in the form of nanofiber mats, able to combine high surface-area-to-volume ratio together with a microporous structure well suited for energetic application. The electro-hydrodynamic techniques involve an external electric field applied to a polymeric solution to provide the final deposition of nanomaterial. As proposed in Figure 1, two different processes, i.e., electrospinning and electrospray, can be classified as electro-hydrodynamic techniques. Indeed, electrospinning ensures the direct assembly of nanofiber mats with different morphologies and properties, as

Figure 1. Different deposition techniques obtained by applying an external electric field and/or potential difference. These techniques can be divided in electrodeposition methods and electro-hydrodynamic techniques, as explained in the

Electrospinning process is based on the principle that strong repulsive forces, induced by external applied electric field, can overcome the surface tension in a charged polymeric jet [2, 3]. Therefore, through this technique, the polymer can be arranged in a mat with a high surfacearea-ratio-to-volume, showing a micro/macroporous structure. Moreover, final nanofibers are based not only on polymers but also on metals, ceramics, and metal oxides, obtained by implementing further different chemical and thermal treatments. However, electrospray technique is an electro-hydrodynamic technique, which occurs at low viscosity values of initial polymeric solution. Indeed, in this case, the surface tension overcomes the viscoelastic forces, and consequently, the instauration of charged droplets with different diameters and concentrations

During the last decade, different works have been presented in the literature, focusing their attention on nanostructured semiconducting metal oxides (as TiO2, ZnO, CuO, and SnO2) in order to design well-performing and green energy systems (such as in dye-sensitized solar cells, lithium-ion batteries (LIBs), fuel cells). In this scenario, nanofibers progressively increased their importance as one of the most important nanostructures to be selected to improve the final

described in this chapter.

160 Semiconductors - Growth and Characterization

diagram.

occurs during the process.

performance of the devices.

The electrospinning is an electro-hydrodynamic process that provides polymer-based fibers with diameter distribution in the range from few nanometers to several micrometers by involving electrostatic forces [4–7]. The process is based on the concept that electrostatic forces induce columbic interactions between charged elements of the polymeric fluid, leading then to overcome the surface tension in a charged polymeric jet and ensuring the nanofiber formation. An electrospinning system is constituted by three major components, as sketched in Figure 2(a): (i) high-voltage supply; (ii) a spinneret, which represents one of the two electrodes, containing the metallic needle of the syringe, where the polymeric solution is loaded; and (iii) the counter electrode, also named grounded electrode, which is the second electrode, where the nanofibers are collected. It involves a high-voltage supply in order to inject charges with a certain polarity in the polymeric solution and then generates a polymeric charged jet, accelerated toward a counter electrode with opposite polarity. In a typical process, the voltage (0–30 kV) is applied between the first electrode (tip of needle) and the second electrode (counter electrode). This implies the indirectly definition of electric field intensity as the ratio between the voltage value and working distance. The working distance is the distance between the first electrode and the counter electrode. The spinneret is linked with a syringe, in which the polymeric (or melt) solution is loaded and a syringe pump allows to control the solution flows with a constant rate, defined as flow rate. When the voltage is applied, the drop at the tip of the needle becomes highly electrified, and the charges are uniformly distributed on its surface. Therefore, the repulsive forces, acted between all charged elements of polymeric solution, induce an elongation of the spherical drop to form a conical shape, known as Taylor's cone. When the repulsive forces

Figure 2. A sketch of electrospinning setup is proposed in (a). In (b) a representation of bending instabilities characterizing the charged polymeric jet during the electrospinning process is proposed. (reprinted with the permission from (polymer, 2008, 49, 2387–2425) copyright (2008) Elsevier).

overcame the surface tension of the droplet, the charged polymeric jet is ejected from the tip of Taylor's cone. During the flight, the solvent evaporation together with the instauration of several instabilities (defined whipping or bending instabilities) [8–10] occurred, leading then to the deposition of nanofiber mat, characterized by a small-size diameter distribution and by a high surface-area-to-volume ratio. In particular, the electrified jet proceeded with a straight path directly toward the counter electrode until the formation of successive instabilities, as sketched in Figure 2(b).

The theoretical principle that explains the correlation between the formation of bending instabilities with the columbic interactions, the external electric field, and the surface tension is not widely investigated in the literature. However, during the process, the free end of the jet shows different envelope loops, which repeats itself in a smaller and smaller scale as the jet diameter is reduced [10]. During the bending instability, the charged jet is divided in sub jets, achieving a progressive diameter reduction, determined by Eq. (1) as explained in the literature [2]:

$$
\sigma\_0^{\,3} = \frac{4\varepsilon m\dot{m}\_0}{k\pi\sigma\rho} \tag{1}
$$

3.1. Polymer solution parameters

tip of the needle.

Several solution parameters, such as viscosity solution, conductivity, dielectric constant, and surface tension, influence the formation of polymeric charged jet and consequently morphological properties of nanofiber mats. The solution viscosity represents the resistance offered by a fluid to its progressive deformation, induced by shear stress or tensile stress. In particular, the viscosity is due to the collisions between all particles that move in a fluid at different velocities. Therefore, solution viscosity can be defined as the measure of force/stress needed to keep the fluid moving in a certain space. The concentration of polymer, dissolved in the solution, directly influences its viscosity the higher the polymeric concentration, the higher the solution viscosity. As determined by several works in the literature, in order to guarantee the instauration of charged polymeric jet during the process, leading to the collection of

The formation of nanofibers with or without defects depends on both viscosity and surface tension of the solution. The surface tension is a polymeric solution property due to all cohesive forces between fluidic molecules, ensuring then/leading then to the distribution of a fluid into the minimum surface area condition. Indeed, in one liquid all inner molecules interact with each neighboring molecule, inducing a resulting force equal to zero and a lower state of energy. On the contrary, since the same number of neighboring does not surround the molecules on the surface, an internal pressure is occurred, which induces the liquid surface to occupy the minimal area, reducing its energy state. According to Laplace's law, the spherical shape can satisfy conditions of minimal area for a liquid [3], minimizing the "wall tension" of the drop surface, as sketched in Figure 3. Related to the electrospinning process, the surface tension of polymeric solution ensures the generation of a spherical droplet, suspended at the

Figure 3. Representation of theoretical concept of Laplace's law (a) represents the cylindrical vessel (T = PR, where R is

the radius of tube) and (b) represents the spherical vessel (T = PR/2, where R is the radius of tube).

ð Þ 0:02 ≤ η ≤ 300 Pa∗s (2)

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163

nanofibers, the viscosity must be in the following range (Eq. (2)) [13–15]:

where ε is the fluidic permittivity (C V�<sup>1</sup> cm�<sup>1</sup> ), m\_<sup>0</sup> is the mass flow rate (g s�<sup>1</sup> ) when r0 (cm) is defined, k is a dimensionless parameter depending on the electric currents, σ is electric conductivity (A V�<sup>1</sup> cm�<sup>1</sup> ), and r is the density (g cm�<sup>3</sup> ) of obtained nanofibers mats. Electrospinning process is applied on the polymer-based materials, including both synthetic and natural polymers. However, metallic carbon nanofibers and ceramic nanofibers can be obtained by electrospinning process, starting from polymeric solutions, and by occurring successive treatments, such as pyrolysis, calcination, and so on. One of the main advantages of this process is represented by the different nanostructures that can be obtained, such as hollow, porous, and dense nanofibers. Therefore, all wide nanostructures are achieved by varying and defining the process parameters, such as electric potential, flow rate, polymer concentration, working distance, and ambient condition.

#### 3. Definition of electrospinning parameters and their correlation with the nanofiber properties

Since the modulation of morphological properties of the nanofiber mats is directly dependent on the process parameters, it is mandatory to define all these process parameters, which can be divided in three main categories [3, 11, 12]:


The first two categories are analyzed in the following paragraphs.

#### 3.1. Polymer solution parameters

overcame the surface tension of the droplet, the charged polymeric jet is ejected from the tip of Taylor's cone. During the flight, the solvent evaporation together with the instauration of several instabilities (defined whipping or bending instabilities) [8–10] occurred, leading then to the deposition of nanofiber mat, characterized by a small-size diameter distribution and by a high surface-area-to-volume ratio. In particular, the electrified jet proceeded with a straight path directly toward the counter electrode until the formation of successive instabilities, as sketched

The theoretical principle that explains the correlation between the formation of bending instabilities with the columbic interactions, the external electric field, and the surface tension is not widely investigated in the literature. However, during the process, the free end of the jet shows different envelope loops, which repeats itself in a smaller and smaller scale as the jet diameter is reduced [10]. During the bending instability, the charged jet is divided in sub jets, achieving a progressive diameter reduction, determined by Eq. (1) as explained in the literature [2]:

> <sup>3</sup> <sup>¼</sup> <sup>4</sup>εm\_<sup>0</sup> kπσr

defined, k is a dimensionless parameter depending on the electric currents, σ is electric conduc-

process is applied on the polymer-based materials, including both synthetic and natural polymers. However, metallic carbon nanofibers and ceramic nanofibers can be obtained by electrospinning process, starting from polymeric solutions, and by occurring successive treatments, such as pyrolysis, calcination, and so on. One of the main advantages of this process is represented by the different nanostructures that can be obtained, such as hollow, porous, and dense nanofibers. Therefore, all wide nanostructures are achieved by varying and defining the process parameters, such as electric potential, flow rate, polymer concentration, working dis-

3. Definition of electrospinning parameters and their correlation with

Since the modulation of morphological properties of the nanofiber mats is directly dependent on the process parameters, it is mandatory to define all these process parameters, which can be

i. Parameters of the polymeric solution (or polymer melt), i.e., viscosity, concentration, and

ii. Parameters of electrospinning process, i.e., voltage, flow rate, and working distance

), m\_<sup>0</sup> is the mass flow rate (g s�<sup>1</sup>

) of obtained nanofibers mats. Electrospinning

(1)

) when r0 (cm) is

r0

), and r is the density (g cm�<sup>3</sup>

where ε is the fluidic permittivity (C V�<sup>1</sup> cm�<sup>1</sup>

in Figure 2(b).

162 Semiconductors - Growth and Characterization

tivity (A V�<sup>1</sup> cm�<sup>1</sup>

tance, and ambient condition.

the nanofiber properties

divided in three main categories [3, 11, 12]:

iii. External parameters, i.e., humidity and temperature

The first two categories are analyzed in the following paragraphs.

polymer molecular weight

between two electrodes

Several solution parameters, such as viscosity solution, conductivity, dielectric constant, and surface tension, influence the formation of polymeric charged jet and consequently morphological properties of nanofiber mats. The solution viscosity represents the resistance offered by a fluid to its progressive deformation, induced by shear stress or tensile stress. In particular, the viscosity is due to the collisions between all particles that move in a fluid at different velocities. Therefore, solution viscosity can be defined as the measure of force/stress needed to keep the fluid moving in a certain space. The concentration of polymer, dissolved in the solution, directly influences its viscosity the higher the polymeric concentration, the higher the solution viscosity. As determined by several works in the literature, in order to guarantee the instauration of charged polymeric jet during the process, leading to the collection of nanofibers, the viscosity must be in the following range (Eq. (2)) [13–15]:

$$(0.02 \le \eta \le 300) \text{ Pa\ast s} \tag{2}$$

The formation of nanofibers with or without defects depends on both viscosity and surface tension of the solution. The surface tension is a polymeric solution property due to all cohesive forces between fluidic molecules, ensuring then/leading then to the distribution of a fluid into the minimum surface area condition. Indeed, in one liquid all inner molecules interact with each neighboring molecule, inducing a resulting force equal to zero and a lower state of energy. On the contrary, since the same number of neighboring does not surround the molecules on the surface, an internal pressure is occurred, which induces the liquid surface to occupy the minimal area, reducing its energy state. According to Laplace's law, the spherical shape can satisfy conditions of minimal area for a liquid [3], minimizing the "wall tension" of the drop surface, as sketched in Figure 3. Related to the electrospinning process, the surface tension of polymeric solution ensures the generation of a spherical droplet, suspended at the tip of the needle.

Figure 3. Representation of theoretical concept of Laplace's law (a) represents the cylindrical vessel (T = PR, where R is the radius of tube) and (b) represents the spherical vessel (T = PR/2, where R is the radius of tube).

In order to explain the correlation between the solution concentration, solution viscosity, and surface tension, it is important to distinguish two different electro-hydrodynamic processes, obtained by using polymeric solutions with different viscosity values. Indeed, if viscosity is lower than 0.1 Pa\*s (η ≤ 0.1 Pa\*s), the surface tension overcomes the viscoelastic forces, a noncontinuous charged polymeric jet is generated, and consequently, droplets with different diameters and with different concentrations are collected. This process is defined as electrospray [13–15]. When viscosity value is higher than 2 Pa\*s (η ≥ 2 Pa\*s), the electrospinning process is ensured, thus providing the formation of nanofibers. Therefore, the charged polymeric jet travels as a continuous jet toward counter electrode, in which dried nanofiber mats were collected on.

charged jet results to be continuous, leading to the formation of nanofiber mats, characterized by a large number of beads together with a nonuniform diameter distribution. Indeed, as the

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This overview on all those parameters, which influence the morphology properties of nanofibers nets, provides important instruments in order to tune/control, during the electrospinning, the formation of some defects, i.e., nano-netting, able to optimize nanofiber mats, involved in several

The nano-nets appear as spider-weblike structure, characterized by secondary ultrathin nanofibers interconnected with the main nanofibers [19]. There are some works in the literature [20–22] that study spider-weblike nanofibrous mat obtained by using an electrospun colloidal solution, containing the polymer and metal oxide nanoparticles. Kim et al. [20] obtained the nano-netting structure starting from a polymeric solution containing solid powder of ZnO mixed with a solution of nylon-6 and acetic acid. The ZnO nanoparticles induce a solution charge density increase, providing the separation of the thinner fibers from the main nanofiber web.

Moreover, Amna et al. [21] proposed the formation of secondary thinner web when ZnO nanoparticles were dissolved in a sol–gel solution of polyurethane in dimethylformamide

One of the main important aspects of the electrospinning process is represented by the possibility to provide different types of nanofiber morphology, obtained by modifying the electrospinning technique. As an example, coaxial electrospinning is applied to the preparation of polymer core-shell nanofibers and hollow nanofibers composed not only of polymers but also of ceramics. Coaxial electrospinning is obtained by using two syringe supports disposed in a concentric configuration, and each syringe contains different spinning solutions, as sketched in Figure 4(A). All parameters, described above, which influence the formation of polymeric jet during the electrospinning process, are the same. Coaxial electrospinning provides further advantages, when the molecular weight of polymer is too low to ensure the fiber formations, avoiding the droplets and consequently the electrospray process. Incorporating these kinds of polymer as the core into a core-shell nanofibers, it is possible to ensure the formation of a continuous jet and consequently the collection of nanofibers on the counter electrode. Moreover, core-shell fibers can offer a solution when it is needed to keep the functional components (proteins, enzymes, bacteria, viruses) maintaining their functionality. Core-shell nanofibers are characterized by a shell, based on solid materials, such as natural or synthetic polymers, and by a core, which is commonly a solvent (like water) with bio-systems. However, the hollow nanofibers are carried out when a wall is based on inorganic polymeric composites or ceramic materials, and the core results to be empty (as sketched in Figure 4(B–D)). There are two different approaches implemented to obtain hollow nanofibers. The first one is based on the concept of sacrificial polymer templates that is then removed. Choi et al. [23] fabricated hollow ZnO nanofibers by using polyvinyl alcohol (PVA) as polymeric template.

applications, like catalysis, sensors, optics, tissue engineering, and energy storage.

flow rate increases, the diameters of nanofibers mat increase [3].

4. Hollow nanostructures and coaxial electrospinning

(DMF).

Different works in the literature, moreover, demonstrate that the increment of solution viscosity guarantees the formation of a uniform mat of nanofibers, without the presence of beads, known as one of the most common defects into the nanofibers mats [13]. It is widely explained how the molecular weight of polymer (Mw) and the polymeric concentration can control the presence of defects and diameter distributions inside the nanofiber mats. As the molecular weight increases, the number of beads and droplet is reduced. Since the increasing of the molecular weight can increase the instabilities distribution, the final nanofiber mats show a nonuniform distribution of diameters. Moreover, a low polymer concentration induces thinner fiber diameters, due to the evaporation of the solvent [14]. The direct correlation between polymeric concentration and viscosity modifies the jet deformations induced by viscoelastic forces during electrospinning. Therefore, when the polymeric concentration is too low, the electrospray process results to be the main electrified process deposition. On the contrary, when the solution viscosity is too high, during the electrospinning, the leak of charged jet from the tip of the needle could be compromised.

#### 3.2. Electrospinning process parameters

All the parameters, related to the electrospinning process, such as voltage, flow rate, and working distance, tune the diameter distribution in the nanofiber mats, thus controlling the porosity distributions and the surface area of nanofibers.

Different works in the literature [3] demonstrated that the correlation between the voltage applied and the nanofiber morphology is not well defined. Nevertheless, this process parameter is quite important in order to establish, for each solution, the threshold value, above which the charged polymeric jet is originated, thus ensuring the nanofiber deposition on the counter electrode. Another fundamental parameter is the working distance, whose value can influence the completely evaporation of the solvent. Indeed, it is needed to define the minimum value of distance, able to provide the fiber's sufficient time to dry before depositing on the collector [15–18].

The flow rate is known as the rate at which the polymer solution is injected to the tip of the needle, defining then the flowing mass of solution and consequently the position of Taylor's cone related to the syringe needle. Moreover, a direct correlation between flow rate and the length of liner path, which precedes the bending instabilities, can be observed during the process [3]. At low values of flow rate, Taylor's cone is formed inside the tip of the needle, thus leading to an intermittent polymeric jet; however, at too high values of flow rates, the charged jet results to be continuous, leading to the formation of nanofiber mats, characterized by a large number of beads together with a nonuniform diameter distribution. Indeed, as the flow rate increases, the diameters of nanofibers mat increase [3].

This overview on all those parameters, which influence the morphology properties of nanofibers nets, provides important instruments in order to tune/control, during the electrospinning, the formation of some defects, i.e., nano-netting, able to optimize nanofiber mats, involved in several applications, like catalysis, sensors, optics, tissue engineering, and energy storage.

The nano-nets appear as spider-weblike structure, characterized by secondary ultrathin nanofibers interconnected with the main nanofibers [19]. There are some works in the literature [20–22] that study spider-weblike nanofibrous mat obtained by using an electrospun colloidal solution, containing the polymer and metal oxide nanoparticles. Kim et al. [20] obtained the nano-netting structure starting from a polymeric solution containing solid powder of ZnO mixed with a solution of nylon-6 and acetic acid. The ZnO nanoparticles induce a solution charge density increase, providing the separation of the thinner fibers from the main nanofiber web.

Moreover, Amna et al. [21] proposed the formation of secondary thinner web when ZnO nanoparticles were dissolved in a sol–gel solution of polyurethane in dimethylformamide (DMF).

#### 4. Hollow nanostructures and coaxial electrospinning

In order to explain the correlation between the solution concentration, solution viscosity, and surface tension, it is important to distinguish two different electro-hydrodynamic processes, obtained by using polymeric solutions with different viscosity values. Indeed, if viscosity is lower than 0.1 Pa\*s (η ≤ 0.1 Pa\*s), the surface tension overcomes the viscoelastic forces, a noncontinuous charged polymeric jet is generated, and consequently, droplets with different diameters and with different concentrations are collected. This process is defined as electrospray [13–15]. When viscosity value is higher than 2 Pa\*s (η ≥ 2 Pa\*s), the electrospinning process is ensured, thus providing the formation of nanofibers. Therefore, the charged polymeric jet travels as a continuous jet toward counter electrode, in which dried nanofiber mats were

Different works in the literature, moreover, demonstrate that the increment of solution viscosity guarantees the formation of a uniform mat of nanofibers, without the presence of beads, known as one of the most common defects into the nanofibers mats [13]. It is widely explained how the molecular weight of polymer (Mw) and the polymeric concentration can control the presence of defects and diameter distributions inside the nanofiber mats. As the molecular weight increases, the number of beads and droplet is reduced. Since the increasing of the molecular weight can increase the instabilities distribution, the final nanofiber mats show a nonuniform distribution of diameters. Moreover, a low polymer concentration induces thinner fiber diameters, due to the evaporation of the solvent [14]. The direct correlation between polymeric concentration and viscosity modifies the jet deformations induced by viscoelastic forces during electrospinning. Therefore, when the polymeric concentration is too low, the electrospray process results to be the main electrified process deposition. On the contrary, when the solution viscosity is too high, during the electrospinning, the leak of charged jet from

All the parameters, related to the electrospinning process, such as voltage, flow rate, and working distance, tune the diameter distribution in the nanofiber mats, thus controlling the

Different works in the literature [3] demonstrated that the correlation between the voltage applied and the nanofiber morphology is not well defined. Nevertheless, this process parameter is quite important in order to establish, for each solution, the threshold value, above which the charged polymeric jet is originated, thus ensuring the nanofiber deposition on the counter electrode. Another fundamental parameter is the working distance, whose value can influence the completely evaporation of the solvent. Indeed, it is needed to define the minimum value of distance, able to provide the fiber's sufficient time to dry before depositing on the collector [15–18].

The flow rate is known as the rate at which the polymer solution is injected to the tip of the needle, defining then the flowing mass of solution and consequently the position of Taylor's cone related to the syringe needle. Moreover, a direct correlation between flow rate and the length of liner path, which precedes the bending instabilities, can be observed during the process [3]. At low values of flow rate, Taylor's cone is formed inside the tip of the needle, thus leading to an intermittent polymeric jet; however, at too high values of flow rates, the

collected on.

164 Semiconductors - Growth and Characterization

the tip of the needle could be compromised.

porosity distributions and the surface area of nanofibers.

3.2. Electrospinning process parameters

One of the main important aspects of the electrospinning process is represented by the possibility to provide different types of nanofiber morphology, obtained by modifying the electrospinning technique. As an example, coaxial electrospinning is applied to the preparation of polymer core-shell nanofibers and hollow nanofibers composed not only of polymers but also of ceramics. Coaxial electrospinning is obtained by using two syringe supports disposed in a concentric configuration, and each syringe contains different spinning solutions, as sketched in Figure 4(A). All parameters, described above, which influence the formation of polymeric jet during the electrospinning process, are the same. Coaxial electrospinning provides further advantages, when the molecular weight of polymer is too low to ensure the fiber formations, avoiding the droplets and consequently the electrospray process. Incorporating these kinds of polymer as the core into a core-shell nanofibers, it is possible to ensure the formation of a continuous jet and consequently the collection of nanofibers on the counter electrode. Moreover, core-shell fibers can offer a solution when it is needed to keep the functional components (proteins, enzymes, bacteria, viruses) maintaining their functionality. Core-shell nanofibers are characterized by a shell, based on solid materials, such as natural or synthetic polymers, and by a core, which is commonly a solvent (like water) with bio-systems. However, the hollow nanofibers are carried out when a wall is based on inorganic polymeric composites or ceramic materials, and the core results to be empty (as sketched in Figure 4(B–D)). There are two different approaches implemented to obtain hollow nanofibers. The first one is based on the concept of sacrificial polymer templates that is then removed. Choi et al. [23] fabricated hollow ZnO nanofibers by using polyvinyl alcohol (PVA) as polymeric template.

(i) flat-plate collector, (ii) rotating drum collector, (iii) rotating wheel with edge, and (iv)

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In the majority of electrospinning setup, a flat-plate collector is used, thus leading to collect a non-woven nanofibers mats, namely as a random distribution of nanofibers on the counter electrode and on all substrate positioned on the top of it. The formation of a non-woven mat of nanofibers is induced by a layer-by-layer deposition on the planar surface. However, some applications required a certain alignment among the nanofibers. In order to induce certain fiber orientations, specific geometries of counter electrode, combined with its motion, are

Rotating cylindrical collectors combined with high rotating speed (up to 1000 rpm), as represented in Figure 5(a), enhance a distribution of parallel nanofibers on it. In this configuration, the two components of velocity (rotating velocity and linear tangential velocity) play a key role in the alignment of nanofibers. The linear tangential velocity of each point on the collector surface is directly proportional to angular velocity and radius of cylinder. Therefore, when the tangential velocity assumes a threshold value able to guarantee the solvent evaporation of the jet, the nanofibers assumed a circular shape on the collector. However, if the tangential velocity results to be too low, random distribution of nanofibers occurred. Finally, if the tangential velocity is too high, the fiber jet will be broken, and the continuous nanofibers will no collect. Different works in the literature used drum collector in order to obtain aligned metal oxide nanofibers, based on TiO2 [25, 26] and ZnO [23]. Aligned ZnO nanofibers are

Figure 5. A schematic representation of different electrospinning collectors. In (a) and (b), the rotating drum and disk

collectors are shown, allowing the aligned nanofiber mats.

parallel strips [4].

5.1. Flat-plate collector

required (as proposed in Figure 5).

5.2. Rotating drum collector

Figure 4. Core-shell nanofibers obtained by coaxial electrospinning (A) and hollow nanofibers (B-C-D) are proposed. The figure is adapted and reprinted with the permission from (Li and Xia [25]. Copyright (2004) American Chemical Society).

Another approach is based on coaxial electrospinning, starting from two immiscible liquids through the coaxial spinneret, followed by a selective removal of the core. Li et al. [24] studied hollow nanofibers, obtained by coaxial electrospinning, and used a polymeric solution of polyvinylpyrrolidone (PVP) and titania precursor (Ti (OiPr)4) as shell and mineral oil as core. An example of the resulting hollow nanofiber is reported in Figure 4(D).

Du et al. [22] used coaxial electrospinning in order to design TiO2/ZnO core-shell nanofibers as photo-anodes in dye-sensitized solar cells (DSSCs). The resulting DSSC efficiency was close to 5%. This improvement can be related to the enhanced light-harvesting efficiency and electron collection efficiency.

## 5. Nanofiber deposition controlled by counter electrode and by patterning

The final step of electrospinning process is represented by the deposition of dried nanofiber mats on the counter electrode (collector). The collector is a conductive electrode, connected to the ground potential in order to provide a stable potential difference between the first electrode (tip of needle) and the second one (counter electrode). In electrospinning process, the deposition texture depends on the electrode configurations. Different works in the literature demonstrated the correlation between the morphological and physical properties of nanofibers with different types of counter electrodes. Indeed, different collectors can be divided into (i) flat-plate collector, (ii) rotating drum collector, (iii) rotating wheel with edge, and (iv) parallel strips [4].

#### 5.1. Flat-plate collector

In the majority of electrospinning setup, a flat-plate collector is used, thus leading to collect a non-woven nanofibers mats, namely as a random distribution of nanofibers on the counter electrode and on all substrate positioned on the top of it. The formation of a non-woven mat of nanofibers is induced by a layer-by-layer deposition on the planar surface. However, some applications required a certain alignment among the nanofibers. In order to induce certain fiber orientations, specific geometries of counter electrode, combined with its motion, are required (as proposed in Figure 5).

#### 5.2. Rotating drum collector

Another approach is based on coaxial electrospinning, starting from two immiscible liquids through the coaxial spinneret, followed by a selective removal of the core. Li et al. [24] studied hollow nanofibers, obtained by coaxial electrospinning, and used a polymeric solution of polyvinylpyrrolidone (PVP) and titania precursor (Ti (OiPr)4) as shell and mineral oil as core.

Figure 4. Core-shell nanofibers obtained by coaxial electrospinning (A) and hollow nanofibers (B-C-D) are proposed. The figure is adapted and reprinted with the permission from (Li and Xia [25]. Copyright (2004) American Chemical Society).

Du et al. [22] used coaxial electrospinning in order to design TiO2/ZnO core-shell nanofibers as photo-anodes in dye-sensitized solar cells (DSSCs). The resulting DSSC efficiency was close to 5%. This improvement can be related to the enhanced light-harvesting efficiency and electron

5. Nanofiber deposition controlled by counter electrode and by patterning

The final step of electrospinning process is represented by the deposition of dried nanofiber mats on the counter electrode (collector). The collector is a conductive electrode, connected to the ground potential in order to provide a stable potential difference between the first electrode (tip of needle) and the second one (counter electrode). In electrospinning process, the deposition texture depends on the electrode configurations. Different works in the literature demonstrated the correlation between the morphological and physical properties of nanofibers with different types of counter electrodes. Indeed, different collectors can be divided into

An example of the resulting hollow nanofiber is reported in Figure 4(D).

collection efficiency.

166 Semiconductors - Growth and Characterization

Rotating cylindrical collectors combined with high rotating speed (up to 1000 rpm), as represented in Figure 5(a), enhance a distribution of parallel nanofibers on it. In this configuration, the two components of velocity (rotating velocity and linear tangential velocity) play a key role in the alignment of nanofibers. The linear tangential velocity of each point on the collector surface is directly proportional to angular velocity and radius of cylinder. Therefore, when the tangential velocity assumes a threshold value able to guarantee the solvent evaporation of the jet, the nanofibers assumed a circular shape on the collector. However, if the tangential velocity results to be too low, random distribution of nanofibers occurred. Finally, if the tangential velocity is too high, the fiber jet will be broken, and the continuous nanofibers will no collect. Different works in the literature used drum collector in order to obtain aligned metal oxide nanofibers, based on TiO2 [25, 26] and ZnO [23]. Aligned ZnO nanofibers are

Figure 5. A schematic representation of different electrospinning collectors. In (a) and (b), the rotating drum and disk collectors are shown, allowing the aligned nanofiber mats.

electrospun starting from a gel containing the precursor of metal oxides and zinc acetate. The electrospun ordered nanofibers were then calcined at 450–500C in oxygen atmosphere to induce the nucleation and growth of ZnO [23]. The aligned ceramic nanofibers show a higher surface-area-to-volume ratio, leading to enhance charge collection and their transport.

6. Application of semiconducting nanofibers

focusing on photovoltaic systems and lithium-ion batteries.

electrons resulting in an electric current generation.

morphology as needed by the final application.

in plastic substrates for flexible devices [39, 40].

6.1. Energy production

For the last decades, tremendous efforts have been devoted to the exploitation of renewable, efficient, and low environmental impact of energy sources, able to contribute to the rise of new models for a sustainable human development [30]. Nanostructured materials have demonstrated a huge potential in energy devices, significantly contributing to improve the final performance of the systems [31, 32]. Nanofibers by electrospinning belong to this intriguing class of materials. As described in the previous paragraphs, nanofibers offer a wide range of strategies to fine tune their morphology in order to meet the requirements of the final application. This versatility of the process provides the nanofibers with a huge potential for energy-related applications [33, 34]. Good examples exist, showing the integration of nanofibers in energy systems for both energy production and storage. In the following paragraphs, some examples are provided, especially

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In photovoltaic devices, photons from solar light are directly converted into electrons thanks to the presence of proper materials. In traditional solar cells, adsorption/conversion is granted by a semiconducting material in a thin-film form, as GaAs, InP, and Si. In the most recently proposed systems, conversion is performed by organic molecules, as in dye-sensitized solar cells (DSSCs) [35, 36] or metalorganic lead halide perovskites in the so-called perovskite solar cells (PSCs) [37]. In both cases, a semiconductor is then needed to capture the generated

The leading wide bandgap, mesoporous semiconductor in these devices is TiO2. Since the photo-generated charges have to be efficiently injected in the conduction band of the nanostructured semiconductor, the higher the injection efficiency, the lower the losses associated to the process. For this reason, the design of the semiconductor at the nanoscale plays a key role to obtain high-performing photo-electrodes. Nanofibers are quite promising candidates for obtaining well-performing devices since they offer several strategies to control and tune their

As an example, the porosity of nanofiber mats can be considered: in DSSC, the possibility to control the dye uptake and the penetration of viscous, solid, or semisolid is a quite important feature. It is possible to change and control the nanofiber mat porosity, tuning the electrospinning process parameters, in order to optimize nanofiber mats for the design of photo-anodes in DSSC [38]. Nanofibers by electrospinning offer several strategies for low-temperature processing of photo-anodes: this is quite important for an easy integration of nanostructured semiconductors

Several metal oxide semiconductors have been successfully fabricated by electrospinning, as

Core-shell nanofibers [41]; TiO2-graphene composite nanofibers [42]; electrospun ZnO photoelectrodes made of ZnO nanofibers with a dense, twisted structure [43]; and SnO2/TiO2 core-shell nanofiber-based photo-anodes have been successfully integrated in DSSC devices. Similarly,

TiO2, ZnO, and SnO2 [39, 40], for the design of well-performing photo-anodes.

TiO2- and Au-decorated TiO2 nanofibers [44–46] have been proposed in PCSs.

#### 5.3. Rotating disk collector

The rotating disk can be defined as a thinner drum collector, on which the nanofibers are deposited on its edge. For this approach, collected nanofibers appear more aligned than the ones obtained with drum collector, described above [12, 27, 28]. Figure 5(b) represents a schematic view of rotating disk, and the nanofibers intercept the edge of counter electrode. With this kind of architecture of counter electrode, the rotation of the disk generates a tangential force, which acts on the polymeric jet, carrying out the nanofiber deposition only on the edge of the disk. This force reduces their diameter, stretching the nanofibers. In this configuration, the alignment of nanofibers results to be better than the one obtained by using rotating drum collector. However, the main limitation of this collector is that only a small quantity of aligned fibers can be obtained. In order to overcome the limitations induced by these kinds of counter electrode to obtain aligned nanofibers and to guarantee the formation of oriented nanofibers, several methods can be implemented. In particular, they mainly involve the modulation of external field provided by a specific geometry of counter electrode.

#### 5.4. Patterning designed on counter electrode

Since some applications in energy field requires highly ordered structure, different works in the literature designed different patternings on a planar counter electrode, able to overcome all limitations introduced by different types of counter electrode (drum or rotating disk counter electrode) and enhance the aligned of nanofibers [29]. An example is shown in Figure 6(A). Two gold bars are placed on the planar counter electrode, and their disposition breaks the asymmetry of the deposition, ensuring the deposition of parallel fibers. In a similar way, a quadripolar arrangement of isolated strips of electrodes, as represented in Figure 6(B), induces a cross grating type of fiber deposition. Wu et al. [12] used two silver plate placed on the planar counter electrode, inducing a final aligned nanofiber mat with a highly ordered structure. The initial polymeric solution was made of CuO precursor mixed with PVA, dissolved in deionized water. The aligned CuO nanofibers, obtained after the calcination treatment conducted in air at a temperature of 500C, enhanced their electrical transfer properties.

Figure 6. (A) Parallel arrangement of nanofibers induced by two gold strips, placed on planar counter electrode and (B) quadripolar arrangement of isolated strips of electrodes.

## 6. Application of semiconducting nanofibers

For the last decades, tremendous efforts have been devoted to the exploitation of renewable, efficient, and low environmental impact of energy sources, able to contribute to the rise of new models for a sustainable human development [30]. Nanostructured materials have demonstrated a huge potential in energy devices, significantly contributing to improve the final performance of the systems [31, 32]. Nanofibers by electrospinning belong to this intriguing class of materials. As described in the previous paragraphs, nanofibers offer a wide range of strategies to fine tune their morphology in order to meet the requirements of the final application. This versatility of the process provides the nanofibers with a huge potential for energy-related applications [33, 34].

Good examples exist, showing the integration of nanofibers in energy systems for both energy production and storage. In the following paragraphs, some examples are provided, especially focusing on photovoltaic systems and lithium-ion batteries.

#### 6.1. Energy production

electrospun starting from a gel containing the precursor of metal oxides and zinc acetate. The electrospun ordered nanofibers were then calcined at 450–500C in oxygen atmosphere to induce the nucleation and growth of ZnO [23]. The aligned ceramic nanofibers show a higher

The rotating disk can be defined as a thinner drum collector, on which the nanofibers are deposited on its edge. For this approach, collected nanofibers appear more aligned than the ones obtained with drum collector, described above [12, 27, 28]. Figure 5(b) represents a schematic view of rotating disk, and the nanofibers intercept the edge of counter electrode. With this kind of architecture of counter electrode, the rotation of the disk generates a tangential force, which acts on the polymeric jet, carrying out the nanofiber deposition only on the edge of the disk. This force reduces their diameter, stretching the nanofibers. In this configuration, the alignment of nanofibers results to be better than the one obtained by using rotating drum collector. However, the main limitation of this collector is that only a small quantity of aligned fibers can be obtained. In order to overcome the limitations induced by these kinds of counter electrode to obtain aligned nanofibers and to guarantee the formation of oriented nanofibers, several methods can be implemented. In particular, they mainly involve the mod-

Since some applications in energy field requires highly ordered structure, different works in the literature designed different patternings on a planar counter electrode, able to overcome all limitations introduced by different types of counter electrode (drum or rotating disk counter electrode) and enhance the aligned of nanofibers [29]. An example is shown in Figure 6(A). Two gold bars are placed on the planar counter electrode, and their disposition breaks the asymmetry of the deposition, ensuring the deposition of parallel fibers. In a similar way, a quadripolar arrangement of isolated strips of electrodes, as represented in Figure 6(B), induces a cross grating type of fiber deposition. Wu et al. [12] used two silver plate placed on the planar counter electrode, inducing a final aligned nanofiber mat with a highly ordered structure. The initial polymeric solution was made of CuO precursor mixed with PVA, dissolved in deionized water. The aligned CuO nanofibers, obtained after the calcination treatment conducted in air at a temperature of 500C, enhanced their electrical

Figure 6. (A) Parallel arrangement of nanofibers induced by two gold strips, placed on planar counter electrode and (B)

surface-area-to-volume ratio, leading to enhance charge collection and their transport.

ulation of external field provided by a specific geometry of counter electrode.

5.4. Patterning designed on counter electrode

quadripolar arrangement of isolated strips of electrodes.

transfer properties.

5.3. Rotating disk collector

168 Semiconductors - Growth and Characterization

In photovoltaic devices, photons from solar light are directly converted into electrons thanks to the presence of proper materials. In traditional solar cells, adsorption/conversion is granted by a semiconducting material in a thin-film form, as GaAs, InP, and Si. In the most recently proposed systems, conversion is performed by organic molecules, as in dye-sensitized solar cells (DSSCs) [35, 36] or metalorganic lead halide perovskites in the so-called perovskite solar cells (PSCs) [37]. In both cases, a semiconductor is then needed to capture the generated electrons resulting in an electric current generation.

The leading wide bandgap, mesoporous semiconductor in these devices is TiO2. Since the photo-generated charges have to be efficiently injected in the conduction band of the nanostructured semiconductor, the higher the injection efficiency, the lower the losses associated to the process. For this reason, the design of the semiconductor at the nanoscale plays a key role to obtain high-performing photo-electrodes. Nanofibers are quite promising candidates for obtaining well-performing devices since they offer several strategies to control and tune their morphology as needed by the final application.

As an example, the porosity of nanofiber mats can be considered: in DSSC, the possibility to control the dye uptake and the penetration of viscous, solid, or semisolid is a quite important feature. It is possible to change and control the nanofiber mat porosity, tuning the electrospinning process parameters, in order to optimize nanofiber mats for the design of photo-anodes in DSSC [38]. Nanofibers by electrospinning offer several strategies for low-temperature processing of photo-anodes: this is quite important for an easy integration of nanostructured semiconductors in plastic substrates for flexible devices [39, 40].

Several metal oxide semiconductors have been successfully fabricated by electrospinning, as TiO2, ZnO, and SnO2 [39, 40], for the design of well-performing photo-anodes.

Core-shell nanofibers [41]; TiO2-graphene composite nanofibers [42]; electrospun ZnO photoelectrodes made of ZnO nanofibers with a dense, twisted structure [43]; and SnO2/TiO2 core-shell nanofiber-based photo-anodes have been successfully integrated in DSSC devices. Similarly, TiO2- and Au-decorated TiO2 nanofibers [44–46] have been proposed in PCSs.

Another important class of devices for energy conversion is represented by fuel cells. They are electrochemical devices, able to convert the chemical energy stored in several classes of molecules, acting as fuels (e.g., H2, methanol, ethanol) into electricity in the presence of a catalyst. Different types of fuel cells exist according to the fuel that is converted (i.e., direct methanol fuel cells (DMFC)), the electrolyte that they use (i.e., solid oxide fuel cells (SOFCs)), or the catalyst that controls the oxidation process (i.e., microbial fuel cells (MFCs)).

High-capacity anodes for Li-ion batteries can be also designed using SnOx, but strategies are needed for this semiconductor to improve its stability over cycles. Indeed, the variation of the volume induced by the intercalation process is detrimental for its mechanical stability, resulting in reduced lifetime of SnOx-based anodes. In order to significantly improve the cycling durability of the resulting anodes, the electrospinning method is used to synthetize

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The possibility to use carbon-based nanofibers in the area of energy storage offers new interesting possibility to design flexible devices. Indeed, carbon-based mats can be processed to be freestanding and usually exhibit very high bendability, offering several possibilities of integration as electrodes in devices for smart electronics. In this area, several processes are developed to decorate the starting carbon mats with metal oxides to design new, well-performing anodes. Samuel et al. [53] decorated carbon-based nanofibers with MnO nanoparticles, demonstrating the possibility to couple the high performances achievable by this semiconducting oxide (923 mAh g�<sup>1</sup> at a

carbon-based nanofibers decorated with small-size SnOx nanoparticles [52].

current rate of 123 mA g�<sup>1</sup> after 90 cycles) with optimal flexibility of the carbon mats.

1 Center For Sustainable Future Technologies@Polito, Istituto Italiano Di Tecnologia, Torino,

2 Department of Applied Science and Technology (DISAT), Politecnico di Torino, Torino, Italy

[1] Jilani A, Abdel-Wahab MS, Hammad AH. Advance deposition techniques for thin film and. In: IntechOpen, editor. Modern Technologies for Creating the Thin-film Systems and

[2] Reneker DH, Yarin AL. Electrospinning jets and polymer nanofibers. Polymer. 2008;49:

[3] Wendorff J, Agarwal S, Greiner A. Electrospinning: Materials, Processing and Applica-

[4] Huang ZM, Zhang YZ, Kotaki M, Ramakrishna S. A review on polymer nanofibers by electrospinning and their applications in nanocomposites. Composites Science and Tech-

[5] Teo WE, Ramakrishna S. A review on electrospinning design and nanofibre assemblies.

coatings. Croatia: Intech; 2017. p. 137-149. DOI: 10.5772/65702

nology. 2003;63:2223-2253. DOI: 10.1016/S0266-3538(03)00178-7

Nanotechnology. 2006;17:89-106. DOI: 10.1088/0957-4484/17/14/R01

Author details

Italy

References

Giulia Massaglia1,2\* and Marzia Quaglio1

\*Address all correspondence to: giulia.massaglia@polito.it

2387-2425. DOI: 10.1016/j.polymer.2008.02.002

tions. Weiheim Germany: Wiley-VCH; 201. 241 pp

In this area of energy, semiconducting nanofibers are especially proposed to design new cathode, when the oxygen reduction reaction (ORR) occurs at the cathode. As an example, manganese oxide nanofibers are successfully proposed to catalyze the ORR as an alternative to platinum [47]. One of the main disadvantages of metal oxides as catalysts to drive the ORR is related to their low electrical conductivity. To overcome this issue, different strategies have been proposed. A successful method is based on the use of composite nanofibers made of doped semiconductors, as proposed by Alvar et al. that optimized a process to embed carbon nanoparticles into mesoporous Nb-doped TiO2 nanofibers [48].

#### 6.2. Energy storage

In the area of energy storage, lithium-ion batteries (LIBs) play a crucial role as a promising technology toward sustainability. In a LIB, a negative electrode and a positive electrode are present, both able of reversibly intercalate Li+ions, and separated by a nonaqueous lithium-ion conducting electrolyte. During discharge, Li+ions carry the current from the negative to the positive electrode, through the nonaqueous electrolyte. During charge, an external high voltage is applied that forces lithium ions to migrate from the positive to the negative electrode, where the process known as intercalation occurs, during which they are embedded in the porous electrode material [31].

In this field, semiconducting nanofibers by electrospinning have been especially proposed for the fabrication of high-efficiency anodes. Good examples are represented by TiO2 nanofibers. Han et al. fabricated TiO2 hollow nanofibers sheathed with TiOxNy/TiN layers with the aim to optimize capability diffusion of lithium ions and electronic conductivity. The fabrication process was based on electrospinning to fabricate hollow nanofibers, followed by a thermal treatment in NH3 atmosphere [49]. Another possibility is represented by the synthesis of composite TiO2 based nanofibers. Zhang et al. proposed the fabrication by electrospinning, followed by a calcination step of TiO2-graphene composite nanofibers able to behave as highly durable anodes [50].

Another interesting possibility offered by electrospinning is the decoration of carbon-based nanofibers with metal oxide catalysts, by adding the oxide precursor into the solution already containing the carbon precursor. The nucleation of the semiconducting oxide in the form of nanoparticles can then be achieved by the thermal process, which also permits the carbonization of the nanofibers. An interesting example of this process is offered by the work of Ji et al. [51]. They synthetized carbon nanofibers decorated with α-Fe2O3 nanoparticles, demonstrating homogenous dispersion of the nanoparticles along the carbon-based nanofibers. The composite mats were tested as anodes in Li-ion batteries; the resulting electrodes showed good reversibility and capacity.

High-capacity anodes for Li-ion batteries can be also designed using SnOx, but strategies are needed for this semiconductor to improve its stability over cycles. Indeed, the variation of the volume induced by the intercalation process is detrimental for its mechanical stability, resulting in reduced lifetime of SnOx-based anodes. In order to significantly improve the cycling durability of the resulting anodes, the electrospinning method is used to synthetize carbon-based nanofibers decorated with small-size SnOx nanoparticles [52].

The possibility to use carbon-based nanofibers in the area of energy storage offers new interesting possibility to design flexible devices. Indeed, carbon-based mats can be processed to be freestanding and usually exhibit very high bendability, offering several possibilities of integration as electrodes in devices for smart electronics. In this area, several processes are developed to decorate the starting carbon mats with metal oxides to design new, well-performing anodes. Samuel et al. [53] decorated carbon-based nanofibers with MnO nanoparticles, demonstrating the possibility to couple the high performances achievable by this semiconducting oxide (923 mAh g�<sup>1</sup> at a current rate of 123 mA g�<sup>1</sup> after 90 cycles) with optimal flexibility of the carbon mats.

## Author details

Another important class of devices for energy conversion is represented by fuel cells. They are electrochemical devices, able to convert the chemical energy stored in several classes of molecules, acting as fuels (e.g., H2, methanol, ethanol) into electricity in the presence of a catalyst. Different types of fuel cells exist according to the fuel that is converted (i.e., direct methanol fuel cells (DMFC)), the electrolyte that they use (i.e., solid oxide fuel cells (SOFCs)), or the

In this area of energy, semiconducting nanofibers are especially proposed to design new cathode, when the oxygen reduction reaction (ORR) occurs at the cathode. As an example, manganese oxide nanofibers are successfully proposed to catalyze the ORR as an alternative to platinum [47]. One of the main disadvantages of metal oxides as catalysts to drive the ORR is related to their low electrical conductivity. To overcome this issue, different strategies have been proposed. A successful method is based on the use of composite nanofibers made of doped semiconductors, as proposed by Alvar et al. that optimized a process to embed carbon

In the area of energy storage, lithium-ion batteries (LIBs) play a crucial role as a promising technology toward sustainability. In a LIB, a negative electrode and a positive electrode are present, both able of reversibly intercalate Li+ions, and separated by a nonaqueous lithium-ion conducting electrolyte. During discharge, Li+ions carry the current from the negative to the positive electrode, through the nonaqueous electrolyte. During charge, an external high voltage is applied that forces lithium ions to migrate from the positive to the negative electrode, where the process known as intercalation occurs, during which they are embedded in the

In this field, semiconducting nanofibers by electrospinning have been especially proposed for the fabrication of high-efficiency anodes. Good examples are represented by TiO2 nanofibers. Han et al. fabricated TiO2 hollow nanofibers sheathed with TiOxNy/TiN layers with the aim to optimize capability diffusion of lithium ions and electronic conductivity. The fabrication process was based on electrospinning to fabricate hollow nanofibers, followed by a thermal treatment in NH3 atmosphere [49]. Another possibility is represented by the synthesis of composite TiO2 based nanofibers. Zhang et al. proposed the fabrication by electrospinning, followed by a calcination step of TiO2-graphene composite nanofibers able to behave as highly durable anodes [50]. Another interesting possibility offered by electrospinning is the decoration of carbon-based nanofibers with metal oxide catalysts, by adding the oxide precursor into the solution already containing the carbon precursor. The nucleation of the semiconducting oxide in the form of nanoparticles can then be achieved by the thermal process, which also permits the carbonization of the nanofibers. An interesting example of this process is offered by the work of Ji et al. [51]. They synthetized carbon nanofibers decorated with α-Fe2O3 nanoparticles, demonstrating homogenous dispersion of the nanoparticles along the carbon-based nanofibers. The composite mats were tested as anodes in Li-ion batteries; the resulting electrodes showed good

catalyst that controls the oxidation process (i.e., microbial fuel cells (MFCs)).

nanoparticles into mesoporous Nb-doped TiO2 nanofibers [48].

6.2. Energy storage

porous electrode material [31].

170 Semiconductors - Growth and Characterization

reversibility and capacity.

Giulia Massaglia1,2\* and Marzia Quaglio1

\*Address all correspondence to: giulia.massaglia@polito.it

1 Center For Sustainable Future Technologies@Polito, Istituto Italiano Di Tecnologia, Torino, Italy

2 Department of Applied Science and Technology (DISAT), Politecnico di Torino, Torino, Italy

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174 Semiconductors - Growth and Characterization


## *Edited by Rosalinda Inguanta and Carmelo Sunseri*

Semiconducting materials are widely used in several applications such as photonics, photovoltaics, electronics, and thermoelectrics, because of their optical and electrooptical features. The fundamental and technological importance of these materials is due to the unique physical and chemical properties. Over the years, numerous methods have been developed for the synthesis of high-efficient semiconductors. Moreover, a variety of approach and characterization methods have been used to study the numerous and fascinating properties of the semiconducting materials. This book collects new developments about semiconductors, from the fundamental issues to their synthesis and applications. Special attention has been devoted to electrochemical growth and characterization.

Photo by undefined undefined / iStock

Semiconductors - Growth and Characterization

Semiconductors

Growth and Characterization

*Edited by Rosalinda Inguanta* 

*and Carmelo Sunseri*