**Ga2O3 Nanowire Synthesis and Device Applications**

DOI: 10.5772/intechopen.72464

 **Nanowire Synthesis and Device Applications**

Badriyah Alhalaili, Howard Mao and Saif Islam Additional information is available at the end of the chapter

Badriyah Alhalaili, Howard Mao and Saif Islam

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.72464

#### **Abstract**

**Ga<sup>2</sup>**

**O3**

In recent years, gallium oxide nanowires have been used in many scientific disciplines due to their outstanding and unique properties. Several applications have focused on incorporating gallium oxide nanowires in devices to improve their performance and efficiency. These distinctive structures bring new opportunities to several research fields and applications such as optoelectronics, electronics, and chemistry. This chapter provides a basic overview of gallium oxide's properties and the growth process of gallium oxide nanowires, with an emphasis on varied applications and future challenges.

**Keywords:** β-Ga<sup>2</sup> O3 nanowires, properties, growth, nanodevices, applications, challenges

#### **1. Introduction**

In the last few years, interest in the use of wide-bandgap materials in semiconductor devices has grown. Ideally, such materials would be abundant, inexpensive, and easy to fabricate, and have high thermal and chemical stability. However, reducing the size of these materials and the devices that utilize them down to the micro/nanoscale remains a big challenge. Furthermore, manipulating their growth to create more directional nanowires could be useful in several applications. In addition, a limited number of devices such as transistors and sensors are capable of operating at temperatures up to 500°C. Wide-bandgap materials can operate in harsh environmental conditions with high temperature, pressure, mechanical vibration and radiation. Therefore, sensors and electronic devices incorporating these materials are attractive choices for increasing the operating temperature to above 500°C while effectively integrating electronics operating even at 1000°C [1].

Preliminary studies demonstrate that some robust wide-bandgap materials can enable a number of devices and sensors that can be used in extreme environments [1, 2]. These benefits can

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 IntechOpen. 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

Recently, researchers have showed interest in the fabrication of low-dimensional β-Ga<sup>2</sup>

sensors, photodetectors, and nanophotonic switches have been fabricated using β-Ga<sup>2</sup>

and thus greater interaction with the surroundings. β-Ga<sup>2</sup>

β-Ga<sup>2</sup> O3

ires. These β-Ga<sup>2</sup>

In 1950, Ga<sup>2</sup>

β-Ga<sup>2</sup> O3

β-Ga<sup>2</sup> O3 . β-Ga<sup>2</sup> O3

As Ga<sup>2</sup> O3

β-Ga<sup>2</sup> O3

**2.1. Crystal structure**

double-chained GaO<sup>6</sup>

**2.2. Electrical properties**

*2.2.2. Dopants for Ga2*

*2.2.1. Energy bandgap for Ga2*

O3

O3

parameters (**Table 1**). The unit cell of β-Ga<sup>2</sup>

make up the unit cell are octahedral (GaO<sup>6</sup>

*O3*

the electrical and optical properties of Ga<sup>2</sup>

posed in an irregular cubic array and are close to one another.

*O3*

ires because of their outstanding properties as compared to bulk structures. The main advantage of

a variety of approaches, and various parameters such as size, length and electronic properties can be controlled during their growth. Recently, a wide array of nanoscale devices such as FETs, gas

[7], solar cell fabrication [8], ultraviolet (UV) limiters [9], and high-temperature gas sensors [10].

O3

and three different O atoms [O(I), O(II) and O(III)] (**Figure 2**). The crystalline structures that

rahedra that share vertices. The unit cell consists of three different oxygen ions that are juxta-

Density functional theory (DFT) has been used to calculate the electronic band structure of

located at the M point, which is marginally less than the direct bandgap of 4.87 eV detected

bandgap material, but due to the weakness of the indirect transitions and the small energy difference between indirect and direct gaps, it is effectively a direct bandgap material. Peak absorption occurs around 4.9 eV. **Table 1** summarizes the electronic properties of β-Ga<sup>2</sup>

trol of sheet resistance [12] and facilitates the formation of ohmic contacts. It can also change many properties of a material. For example, it can increase the lattice constant [13], and change crystallinity [14] and electron mobility [15]. Doping can even be used to tune the optical bandgap [16] and introduce defect levels [17]. It is typically done at the same time nanowires are grown, using methods such as chemical vapor deposition, pulsed laser deposition, and evaporation. Alternatively, dopants can be added directly into the material through ion implantation.

is an insulator, it needs to be doped to increase its conductivity. Doping enables con-

is assumed to be an n-type semiconductor due to its shallow donor oxygen vacancies

and ionization energy of 30–40 meV [9]. Its electrical conduction and free carrier concentration can be adjusted by doping it with Si. The effects of oxygen vacancies and other impurities on

O3

at the Γ-point, as shown in **Figure 3** [11]. Research has confirmed that β-Ga<sup>2</sup>

nanowires is a higher surface-to-volume ratio, offering more surface states at the interface

O3

nanowire devices are also popular for diverse applications such as displays

was observed, and in 1960, its crystal structure was discovered by Geller [7].

) and tetrahedral (GaO<sup>4</sup>

octahedra that share edges are connected by single chains of GaO<sup>4</sup>

has an indirect bandgap of 4.83 eV, with a valence-band maximum (VBM)

is part of the C2/m space group and has a base-centered arrangement with four lattice

O3

O3

). Along the b-axis, the

O3

have also been examined [11]. It was found that

is an indirect

O3 .

nanowires can be synthesized using

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

has two different Ga atoms [Ga(I) and Ga(II)]

nanow-

9

nanow-

tet-

**Figure 1.** (a) Relationship between the bandgap and breakdown field for various semiconductor materials, (b) On-resistance as a function of breakdown voltage for major semiconductors. (Reprinted with permission from [82]. Copyright 2012, American Institute of Physics.).

be realized by exploiting the electrical, optical, thermal, physical and chemical properties of wide-bandgap semiconductors including group III nitrides such as gallium nitride (GaN), aluminum nitride (AlN) and their ternary and quaternary alloys; and silicon carbide (SiC), diamond and some wide-bandgap oxides (such as Ga<sup>2</sup> O3 ). These materials offer superior electrical, optical, mechanical, and chemical properties when compared to other widely used semiconductor materials such as Si, GaAs and InP [1, 3].

Gallium oxide has a promising future with respect to other semiconductors based on its material properties (**Figure 1a**). It has the potential to be more widely used in power and optoelectronics applications than not only Si but also SiC and GaN. This is due to its high breakdown voltage and low on-resistance (**Figure 1b**). These features have led to the evaluation of Ga<sup>2</sup> O3 nanowires as a novel material in semiconductor research fields; it is also less expensive and easier to fabricate than SiC and GaN. Compared to thin films, nanowires have a higher surface-to-volume ratio, which increases their sensitivity in detection. Ga<sup>2</sup> O3 nanowires have an established place in many applications, such as optical and sensing studies [4, 5]. For example, β-Ga<sup>2</sup> O3 nanowirebased devices are very attractive for use as gas sensors [6] due to their stability, moisture resistance, fast response, and long lifetime. Research interest in Ga<sup>2</sup> O3 nanowires has been increasing, and it has obtained recognition among other wide-bandgap materials. This chapter will mainly focus on synthesis techniques, device applications, and future challenges for β-Ga<sup>2</sup> O3 nanowires.

#### **2. Fundamental properties of β-Ga2 O3 nanowires**

Gallium (III) oxide is an insulating metal oxide with five different polymorphs (*α*-, *β*- *γ*-, *δ*-, ϵ-). β-Ga<sup>2</sup> O3 has a monoclinic structure and is the most stable form both chemically and thermally. The other phases exist in a metastable state and are converted into β-phase at temperatures above 600°C. β-Ga<sup>2</sup> O3 has exceptional properties which have attracted the attention of scientists, including a wide bandgap of 4.9 eV, a high melting point of 1900°C, excellent electrical conductivity, and photoluminescence. Despite these wonderful properties, it is only in recent years that interest in this material has grown.

Recently, researchers have showed interest in the fabrication of low-dimensional β-Ga<sup>2</sup> O3 nanowires because of their outstanding properties as compared to bulk structures. The main advantage of β-Ga<sup>2</sup> O3 nanowires is a higher surface-to-volume ratio, offering more surface states at the interface and thus greater interaction with the surroundings. β-Ga<sup>2</sup> O3 nanowires can be synthesized using a variety of approaches, and various parameters such as size, length and electronic properties can be controlled during their growth. Recently, a wide array of nanoscale devices such as FETs, gas sensors, photodetectors, and nanophotonic switches have been fabricated using β-Ga<sup>2</sup> O3 nanowires. These β-Ga<sup>2</sup> O3 nanowire devices are also popular for diverse applications such as displays [7], solar cell fabrication [8], ultraviolet (UV) limiters [9], and high-temperature gas sensors [10].

#### **2.1. Crystal structure**

be realized by exploiting the electrical, optical, thermal, physical and chemical properties of wide-bandgap semiconductors including group III nitrides such as gallium nitride (GaN), aluminum nitride (AlN) and their ternary and quaternary alloys; and silicon carbide (SiC),

**Figure 1.** (a) Relationship between the bandgap and breakdown field for various semiconductor materials, (b) On-resistance as a function of breakdown voltage for major semiconductors. (Reprinted with permission from [82]. Copyright 2012,

electrical, optical, mechanical, and chemical properties when compared to other widely used

Gallium oxide has a promising future with respect to other semiconductors based on its material properties (**Figure 1a**). It has the potential to be more widely used in power and optoelectronics applications than not only Si but also SiC and GaN. This is due to its high breakdown voltage

as a novel material in semiconductor research fields; it is also less expensive and easier to fabricate than SiC and GaN. Compared to thin films, nanowires have a higher surface-to-volume

based devices are very attractive for use as gas sensors [6] due to their stability, moisture resis-

and it has obtained recognition among other wide-bandgap materials. This chapter will mainly

**O3**

Gallium (III) oxide is an insulating metal oxide with five different polymorphs (*α*-, *β*- *γ*-, *δ*-, ϵ-).

The other phases exist in a metastable state and are converted into β-phase at temperatures

tists, including a wide bandgap of 4.9 eV, a high melting point of 1900°C, excellent electrical conductivity, and photoluminescence. Despite these wonderful properties, it is only in recent

has a monoclinic structure and is the most stable form both chemically and thermally.

and low on-resistance (**Figure 1b**). These features have led to the evaluation of Ga<sup>2</sup>

in many applications, such as optical and sensing studies [4, 5]. For example, β-Ga<sup>2</sup>

focus on synthesis techniques, device applications, and future challenges for β-Ga<sup>2</sup>

O3

O3

 **nanowires**

has exceptional properties which have attracted the attention of scien-

O3

). These materials offer superior

nanowires have an established place

O3

O3

O3

nanowires has been increasing,

nanowires

nanowire-

nanowires.

diamond and some wide-bandgap oxides (such as Ga<sup>2</sup>

American Institute of Physics.).

8 Novel Nanomaterials - Synthesis and Applications

semiconductor materials such as Si, GaAs and InP [1, 3].

ratio, which increases their sensitivity in detection. Ga<sup>2</sup>

**2. Fundamental properties of β-Ga2**

O3

years that interest in this material has grown.

β-Ga<sup>2</sup> O3

above 600°C. β-Ga<sup>2</sup>

tance, fast response, and long lifetime. Research interest in Ga<sup>2</sup>

In 1950, Ga<sup>2</sup> O3 was observed, and in 1960, its crystal structure was discovered by Geller [7]. β-Ga<sup>2</sup> O3 is part of the C2/m space group and has a base-centered arrangement with four lattice parameters (**Table 1**). The unit cell of β-Ga<sup>2</sup> O3 has two different Ga atoms [Ga(I) and Ga(II)] and three different O atoms [O(I), O(II) and O(III)] (**Figure 2**). The crystalline structures that make up the unit cell are octahedral (GaO<sup>6</sup> ) and tetrahedral (GaO<sup>4</sup> ). Along the b-axis, the double-chained GaO<sup>6</sup> octahedra that share edges are connected by single chains of GaO<sup>4</sup> tetrahedra that share vertices. The unit cell consists of three different oxygen ions that are juxtaposed in an irregular cubic array and are close to one another.

#### **2.2. Electrical properties**

#### *2.2.1. Energy bandgap for Ga2 O3*

Density functional theory (DFT) has been used to calculate the electronic band structure of β-Ga<sup>2</sup> O3 . β-Ga<sup>2</sup> O3 has an indirect bandgap of 4.83 eV, with a valence-band maximum (VBM) located at the M point, which is marginally less than the direct bandgap of 4.87 eV detected at the Γ-point, as shown in **Figure 3** [11]. Research has confirmed that β-Ga<sup>2</sup> O3 is an indirect bandgap material, but due to the weakness of the indirect transitions and the small energy difference between indirect and direct gaps, it is effectively a direct bandgap material. Peak absorption occurs around 4.9 eV. **Table 1** summarizes the electronic properties of β-Ga<sup>2</sup> O3 .

#### *2.2.2. Dopants for Ga2 O3*

As Ga<sup>2</sup> O3 is an insulator, it needs to be doped to increase its conductivity. Doping enables control of sheet resistance [12] and facilitates the formation of ohmic contacts. It can also change many properties of a material. For example, it can increase the lattice constant [13], and change crystallinity [14] and electron mobility [15]. Doping can even be used to tune the optical bandgap [16] and introduce defect levels [17]. It is typically done at the same time nanowires are grown, using methods such as chemical vapor deposition, pulsed laser deposition, and evaporation. Alternatively, dopants can be added directly into the material through ion implantation.

β-Ga<sup>2</sup> O3 is assumed to be an n-type semiconductor due to its shallow donor oxygen vacancies and ionization energy of 30–40 meV [9]. Its electrical conduction and free carrier concentration can be adjusted by doping it with Si. The effects of oxygen vacancies and other impurities on the electrical and optical properties of Ga<sup>2</sup> O3 have also been examined [11]. It was found that


to UV wavelengths. Doping causes the photocurrent to be much larger than that of undoped devices. This is also true for the ratio of photocurrent to dark current. Indium doping can also

tation of Si atoms and annealing [12]. Si doping also helps in creating an ohmic contact

, significantly reducing its resistance after ion implan-

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464 11

. (Reprinted with permission from Ref. [11]. Copyright 2010, American

reduce rise time and improve spectral responsivity and external quantum efficiency.

O3

O3

*2.2.2.2. Silicon (Si)*

**Figure 2.** Unit cell of β-Ga<sup>2</sup>

Si has also been used to dope Ga<sup>2</sup>

**Figure 3.** Electronic band structure of β-Ga<sup>2</sup>

Institute of Physics.).

O3 .

**Table 1.** Summary of the basic properties of β-Ga<sup>2</sup> O3 [18, 29].

it cannot be assumed that oxygen vacancies are the main cause for conduction and luminance because Si is the main impurity behind electrical conduction [18]. Furthermore, these oxygen vacancies are deep donors and hence cannot describe the n-type conductivity. Other dopants such as Si, Ge, Sn, F, and Cl behave like shallow donors. In the next sub-section, we describe some dopants and their properties. **Table 2** summarizes the materials that have been used as dopants for Ga<sup>2</sup> O3 and the changes they make to it.

#### *2.2.2.1. Indium (In)*

Indium is one dopant that can be used with Ga<sup>2</sup> O3 to improve its photoelectrical properties, especially in the context of photodetection [13]. Doping with indium narrows the bandgap, which increases the range of the photoresponse but makes the photodetector less sensitive

**Figure 2.** Unit cell of β-Ga<sup>2</sup> O3 .

to UV wavelengths. Doping causes the photocurrent to be much larger than that of undoped devices. This is also true for the ratio of photocurrent to dark current. Indium doping can also reduce rise time and improve spectral responsivity and external quantum efficiency.

#### *2.2.2.2. Silicon (Si)*

it cannot be assumed that oxygen vacancies are the main cause for conduction and luminance because Si is the main impurity behind electrical conduction [18]. Furthermore, these oxygen vacancies are deep donors and hence cannot describe the n-type conductivity. Other dopants such as Si, Ge, Sn, F, and Cl behave like shallow donors. In the next sub-section, we describe some dopants and their properties. **Table 2** summarizes the materials that have been used as

O3 [18, 29].

O3

especially in the context of photodetection [13]. Doping with indium narrows the bandgap, which increases the range of the photoresponse but makes the photodetector less sensitive

to improve its photoelectrical properties,

and the changes they make to it.

dopants for Ga<sup>2</sup>

*2.2.2.1. Indium (In)*

O3

Indium is one dopant that can be used with Ga<sup>2</sup>

**Property Attributes** Crystal structure Monoclinic Group of symmetry C2/m

10 Novel Nanomaterials - Synthesis and Applications

*a* 12.214 Å *b* 3.037 Å *c* 5.798 Å β 103.83 Å

Density 5.863 g/cm<sup>3</sup> Dopants (n/p type) Ti/Zn, Ge, Mg Electron affinity (*χ*) 3.50 eV Electron effective mass 0.342*me*

Carrier density 1013–1016 cm−3 Rel. dielectric constant (ϵ) 9.9–10.2 Melting point 1800 ° *C*

Specific heat 0.49–0.56 Jg−1 K−1

[−201] 13.3 ± 1.0 W/mK [001] 14.7 ±1.5 W/mK

∥[100] 1.9523 ⊥(100) 1.9201

**Table 1.** Summary of the basic properties of β-Ga<sup>2</sup>

) 4.83 eV (indirect) & 4.87 (direct) at RT

) 10 cm<sup>2</sup> V−1 s−1

[100] (10.9 ±1.0 W/mK)–(13 W/mK)

[110] (27.0 ±2.0 W/mK)–(21 W/mK)

Lattice parameters

Bandgap (E<sup>g</sup>

Electron mobility (*μ*<sup>e</sup>

Thermal conductivity

Refractive index (*n*) @ 532 nm

Si has also been used to dope Ga<sup>2</sup> O3 , significantly reducing its resistance after ion implantation of Si atoms and annealing [12]. Si doping also helps in creating an ohmic contact

**Figure 3.** Electronic band structure of β-Ga<sup>2</sup> O3 . (Reprinted with permission from Ref. [11]. Copyright 2010, American Institute of Physics.).


*2.2.2.5. Copper (Cu)*

bandgap than Ga<sup>2</sup>

*2.2.2.6. Zinc (Zn)*

undoped β-Ga<sup>2</sup>

*2.2.2.7. Titanium (Ti)*

*2.2.2.8. Nitrogen (N)*

O3

**2.3. Optical properties**

*2.3.1. Absorption of Ga2*

The absorption spectra of Ga<sup>2</sup>

O3

O3

ear I-V curve. This suggests that doping Ga<sup>2</sup>

*O3*

 *nanowires*

O3

titanium ions high solubility in Ga<sup>2</sup>

Doping Ga<sup>2</sup>

levels in Ga<sup>2</sup>

than Ga<sup>2</sup>

Ga<sup>2</sup> O3 O3

O3

copper has the potential to be a p-type dopant for Ga<sup>2</sup>

Zn has potential to be used as a p-type dopant for Ga<sup>2</sup>

and O [20]. After doping, the acceptor impurity level in Ga<sup>2</sup>

has also been doped with Cu, which reduces the bandgap because CuO has a narrower

similar ionic size (0.073 nm and 0.062 nm, respectively), and copper is a group IB element, so

trons than Ga does, there is a lower density of electrons between Cu and O than between Ga

clearly determined because there were both electrons and holes and their concentrations were approximately equal. High transmittance was observed, indicating that crystalline quality was high. Adding Zn also caused the carrier concentration to drop from 1.4×1014 to 7.2×1011 cm−3.

due to their similar ionic radii (0.047 and 0.042 nm for Ga3+ and Ti4+, respectively). This gives

and reduces leakage current by increasing the activation energy and narrowing the bandgap.

N3− has an ionic radius about the same as that of O2−, so it can be expected to form shallow acceptor levels [23], which is significant because the difficulty of forming shallow acceptor

gap decreased and the valence band maximum and conduction band maximum met at the

ires were doped with N, they exhibited p-type properties [24]. Pt has a larger work function

larger than that of the semiconductor, a Schottky barrier is created when the semiconductor is n-type and an ohmic contact is exhibited for a p-type semiconductor. While an undoped

O3

slightly around 270 nm [18]. The band at 260 nm was due to the intrinsic band-to-band transition while the one at 270 nm was caused by Ga3+ vacancies in the conduction band. The

nanowire had an I-V characteristic that was rectifying, a N-doped nanowire had a lin-

and it was used for the metal contacts. When a contact metal has a work function

with N can make it p-type.

show cutoff absorption edges at around 255–260 nm and

is the main challenge in making it p-type. On doping Ga<sup>2</sup>

same point. The acceptor impurity levels were above the valence band. After Ga<sup>2</sup>

annealing and belongs to group IIB with one fewer valence electron than Ga<sup>2</sup>

O3

[17]. The bandgap widens again after annealing. Ga3+ and Cu2+ have a

O3

. Because Cu has fewer valence elec-

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

as it can add holes to the material after

O3

with N, the band-

O3

nanow-

is higher than the valence band.

O3

[21]. While

13

O3

O3

O3

with Ti lowers its dielectric permittivity. [22]. Ti can dope about 10% of Ga<sup>2</sup>

. Inclusion of Ti increases the optosensitivity of samples

is clearly n-type, after doping with Zn the conductivity type could not be

Ga<sup>2</sup> O3

**Table 2.** Some materials have been used as a dopant and their properties.

on the surface with Ti and Au. It was seen to increase the rejection ratio, photocurrent, responsivity, and quantum efficiency of a photodetector at high temperatures. *γ*-Ga<sup>2</sup> O3 has been doped n-type by ablation of a Ga<sup>2</sup> O3 :Si ceramic target with laser pulses [19]. In that case, doping capability was independent of crystal structure, and carrier concentration did not vary much with temperature, meaning the Ga<sup>2</sup> O3 had been degenerately doped.

#### *2.2.2.3. Tin (Sn)*

Ga<sup>2</sup> O3 has also been doped with Sn to increase its conductivity [14]. Interestingly, when deposition was performed at temperatures above 410°C, doping caused the phase of the Ga<sup>2</sup> O3 to change from β-Ga<sup>2</sup> O3 to ε-Ga<sup>2</sup> O3 . Another method of doping Ga<sup>2</sup> O3 with Sn is heating SnO<sup>2</sup> powder in a K-cell with ozone as an oxidizer [15]. However, it was observed that as the temperature of the SnO<sup>2</sup> powder was raised, the carrier concentration increased, which lowered electron mobility.

#### *2.2.2.4. Magnesium (Mg)*

Mg is one of a few materials believed to hold promise in making Ga<sup>2</sup> O3 p-type. Mg-doped Ga<sup>2</sup> O3 has been deposited into an MOCVD system with a Mg-containing source, and as the Mg concentration was increased from 1–10%, the crystalline quality degraded, with films becoming amorphous [16]. Annealing improved the crystallinity and lowered the resistivity, but the Ga<sup>2</sup> O3 did not become p-type as expected after doping.

#### *2.2.2.5. Copper (Cu)*

Ga<sup>2</sup> O3 has also been doped with Cu, which reduces the bandgap because CuO has a narrower bandgap than Ga<sup>2</sup> O3 [17]. The bandgap widens again after annealing. Ga3+ and Cu2+ have a similar ionic size (0.073 nm and 0.062 nm, respectively), and copper is a group IB element, so copper has the potential to be a p-type dopant for Ga<sup>2</sup> O3 . Because Cu has fewer valence electrons than Ga does, there is a lower density of electrons between Cu and O than between Ga and O [20]. After doping, the acceptor impurity level in Ga<sup>2</sup> O3 is higher than the valence band.

#### *2.2.2.6. Zinc (Zn)*

Zn has potential to be used as a p-type dopant for Ga<sup>2</sup> O3 as it can add holes to the material after annealing and belongs to group IIB with one fewer valence electron than Ga<sup>2</sup> O3 [21]. While undoped β-Ga<sup>2</sup> O3 is clearly n-type, after doping with Zn the conductivity type could not be clearly determined because there were both electrons and holes and their concentrations were approximately equal. High transmittance was observed, indicating that crystalline quality was high. Adding Zn also caused the carrier concentration to drop from 1.4×1014 to 7.2×1011 cm−3.

#### *2.2.2.7. Titanium (Ti)*

Doping Ga<sup>2</sup> O3 with Ti lowers its dielectric permittivity. [22]. Ti can dope about 10% of Ga<sup>2</sup> O3 due to their similar ionic radii (0.047 and 0.042 nm for Ga3+ and Ti4+, respectively). This gives titanium ions high solubility in Ga<sup>2</sup> O3 . Inclusion of Ti increases the optosensitivity of samples and reduces leakage current by increasing the activation energy and narrowing the bandgap.

#### *2.2.2.8. Nitrogen (N)*

O3

O3 to

:Si ceramic target with laser pulses [19]. In

had been degenerately

O3

with Sn is heating SnO<sup>2</sup>

p-type. Mg-doped

O3

O3

O3

on the surface with Ti and Au. It was seen to increase the rejection ratio, photocurrent, responsivity, and quantum efficiency of a photodetector at high temperatures. *γ*-Ga<sup>2</sup>

O3

that case, doping capability was independent of crystal structure, and carrier concentra-

has also been doped with Sn to increase its conductivity [14]. Interestingly, when depo-

. Another method of doping Ga<sup>2</sup>

 has been deposited into an MOCVD system with a Mg-containing source, and as the Mg concentration was increased from 1–10%, the crystalline quality degraded, with films becoming amorphous [16]. Annealing improved the crystallinity and lowered the resistivity,

powder was raised, the carrier concentration increased, which lowered

O3

O3

O3

O3 p-type

• To support creating an ohmic contact on the surface

O3 p-type

Cu • To decrease bandgap because CuO has a narrower bandgap than Ga<sup>2</sup>

• To decrease carrier concentration

• To increase the optosensitivity • To reduce leakage current

• A potential to be a p-type dopant for Ga<sup>2</sup>

powder in a K-cell with ozone as an oxidizer [15]. However, it was observed that as the tem-

sition was performed at temperatures above 410°C, doping caused the phase of the Ga<sup>2</sup>

has been doped n-type by ablation of a Ga<sup>2</sup>

**Materials Properties**

12 Novel Nanomaterials - Synthesis and Applications

In To improve the photoelectrical properties

Zn • A potential to be a p-type dopant for Ga<sup>2</sup>

N Doping with N has the potential to make Ga<sup>2</sup>

Ti • To decrease dielectric permittivity

**Table 2.** Some materials have been used as a dopant and their properties.

Si • To reduce sheet resistance of Ga<sup>2</sup>

Sn To increase conductivity. Mg A promise in making Ga<sup>2</sup>

O3

to ε-Ga<sup>2</sup>

doped.

Ga<sup>2</sup> O3

Ga<sup>2</sup> O3

but the Ga<sup>2</sup>

*2.2.2.3. Tin (Sn)*

change from β-Ga<sup>2</sup>

perature of the SnO<sup>2</sup>

*2.2.2.4. Magnesium (Mg)*

O3

electron mobility.

tion did not vary much with temperature, meaning the Ga<sup>2</sup>

O3

Mg is one of a few materials believed to hold promise in making Ga<sup>2</sup>

did not become p-type as expected after doping.

N3− has an ionic radius about the same as that of O2−, so it can be expected to form shallow acceptor levels [23], which is significant because the difficulty of forming shallow acceptor levels in Ga<sup>2</sup> O3 is the main challenge in making it p-type. On doping Ga<sup>2</sup> O3 with N, the bandgap decreased and the valence band maximum and conduction band maximum met at the same point. The acceptor impurity levels were above the valence band. After Ga<sup>2</sup> O3 nanowires were doped with N, they exhibited p-type properties [24]. Pt has a larger work function than Ga<sup>2</sup> O3 and it was used for the metal contacts. When a contact metal has a work function larger than that of the semiconductor, a Schottky barrier is created when the semiconductor is n-type and an ohmic contact is exhibited for a p-type semiconductor. While an undoped Ga<sup>2</sup> O3 nanowire had an I-V characteristic that was rectifying, a N-doped nanowire had a linear I-V curve. This suggests that doping Ga<sup>2</sup> O3 with N can make it p-type.

#### **2.3. Optical properties**

#### *2.3.1. Absorption of Ga2 O3 nanowires*

The absorption spectra of Ga<sup>2</sup> O3 show cutoff absorption edges at around 255–260 nm and slightly around 270 nm [18]. The band at 260 nm was due to the intrinsic band-to-band transition while the one at 270 nm was caused by Ga3+ vacancies in the conduction band. The absorption range (255–260 nm) can be obtained because of the transition from the valence band to the conduction band [25]. The absorption of Ga<sup>2</sup> O3 is influenced by polarization of incident light [26]. The difference between the absorption edge at 260 and at 270 nm was attained using the transition from the valence band.

These studies showed that the spectra characteristic of undoped samples are independent of the excitation and emission wavelengths. However, the Si-doped sample produced a blue emission which was highly related to the excitation and emission wavelengths. Finally, it

 is considered a weak thermal conductor compared to other semiconductors, with a thermal conductivity lower than that of other wide-bandgap materials such as SiC and

has the highest thermal conductivity along the [010] direction and the lowest

O3

O3

talline anisotropy. Directions with a smaller lattice constant have higher thermal conduc-

along the [100] direction. This has been measured at temperatures between 80 and 495 K (**Figure 4**). At high temperatures, thermal transport is primarily dominated by phonon scat-

oxidation [36], vapor–liquid–solid mechanism [37], pulsed laser deposition [38], sputtering [39], thermal evaporation [40–42], molecular beam epitaxy [43], laser ablation [44], arc discharge [45], carbothermal reduction [46], microwave plasma [47], metalorganic chemical vapor deposition [48] and the hydrothermal method [49, 50]. **Table 3** summarizes the advantages and disadvantages of these methods. Research is ongoing to determine the best method for grow-

nanowires have been fabricated by oxidizing GaAs in a furnace [36]. At 1050°C and

nanowires are grown using this method, the material first begins as a thin film

O3

surface. As the GaAs was heated, it dissociated and the arsenic evaporated, leaving behind

began to grow. The roughness of the surface was believed to be caused by oxygen-deficiency

before nanowire formation if conditions are favorable. The steps in this process have been clearly described for GaAs oxidation in a furnace [51]. At high temperatures, phase separation of the GaAs occurs, and the constituents decompose. Arsenic diffuses through the substrate and evaporates from the surface, and gallium melts, forming clusters of liquid gallium on the surface. As the clusters form, the arsenic evaporates faster because it is easier for it to dissociate through the clusters than from the surface of the GaAs. Nanowire growth occurs from these clusters. Due to liquid Ga's ability to wet GaAs, the droplets are hemispherical. Oxygen

O3

a liquid Ga-rich surface which was then oxidized. At high temperatures, Ga<sup>2</sup>

defects and the difference in crystal structure between GaAs and Ga<sup>2</sup>

varies depending on crystal direction due to its crys-

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464 15

nanowires. These include thermal

nanowires with a nanotextured

thin film, with voids where the As

, Ga is no longer able to wet it, resulting in

O3 . O3

nanowires

has been assumed that the red emission is caused by nitrogen impurities [32, 34].

O3

 **nanowire growth mechanisms**

A wide variety of methods have been used to grow Ga<sup>2</sup>

at atmospheric pressure, argon gas was used to grow Ga<sup>2</sup>

in the chamber reacts with the Ga clusters to form a Ga<sup>2</sup>

evaporated. As the GaAs is converted into Ga<sup>2</sup>

*2.3.3. Thermal properties*

O3

GaN. Thermal conductivity in Ga<sup>2</sup>

ing high-quality nanowires at low cost.

**3.1. Thermal oxidation**

O3

Ga<sup>2</sup> O3

tivity., Ga<sup>2</sup>

tering [35].

**O3**

**3. Ga2**

Ga<sup>2</sup> O3

When β-Ga<sup>2</sup>

#### *2.3.2. Luminescence properties of Ga2 O3 nanowires*

Ga<sup>2</sup> O3 can produce four emissions: UV (3.2–3.6 eV), blue (2.8–3.0 eV), green (2.4 eV) [27, 28] and red [29]. Due to its wide bandgap, there are a number of defect states, which cause its different emissions from infrared to UV because its wide bandgap has a number of defect states that are responsible for them. The bandgap of Ga<sup>2</sup> O3 is known to be 4.8 eV, suggesting that UV emission initiated from the edge of the band is due to free electrons and recombination of self-trapped holes [30]. However, the blue and deep-blue luminescence are caused by the recombination of an electron on a defect donor state formed by oxygen vacancies and a hole on an acceptor state formed by a gallium vacancy or a gallium–oxygen vacancy pair [28, 31]. Green emission could only be achieved when the samples were doped with certain impurities such as Be, Ge, Sn, Li, Zr, and Si [18, 32]. It has been suggested that green luminescence is related to self-trapped or bound excitons [26]. Several studies have been performed on the excitation and photoluminescence of pure and Si-doped β-Ga<sup>2</sup> O3 single crystals [33].

**Figure 4.** Thermal conductivity as a function of temperature for β-Ga<sup>2</sup> O3 along different crystal directions. (Reprinted with permission from Ref. [35]. Copyright 2015, American Institute of Physics.).

These studies showed that the spectra characteristic of undoped samples are independent of the excitation and emission wavelengths. However, the Si-doped sample produced a blue emission which was highly related to the excitation and emission wavelengths. Finally, it has been assumed that the red emission is caused by nitrogen impurities [32, 34].

#### *2.3.3. Thermal properties*

absorption range (255–260 nm) can be obtained because of the transition from the valence

incident light [26]. The difference between the absorption edge at 260 and at 270 nm was

 can produce four emissions: UV (3.2–3.6 eV), blue (2.8–3.0 eV), green (2.4 eV) [27, 28] and red [29]. Due to its wide bandgap, there are a number of defect states, which cause its different emissions from infrared to UV because its wide bandgap has a number of defect

that UV emission initiated from the edge of the band is due to free electrons and recombination of self-trapped holes [30]. However, the blue and deep-blue luminescence are caused by the recombination of an electron on a defect donor state formed by oxygen vacancies and a hole on an acceptor state formed by a gallium vacancy or a gallium–oxygen vacancy pair [28, 31]. Green emission could only be achieved when the samples were doped with certain impurities such as Be, Ge, Sn, Li, Zr, and Si [18, 32]. It has been suggested that green luminescence is related to self-trapped or bound excitons [26]. Several studies have been performed

O3

O3

O3

along different crystal directions. (Reprinted

is influenced by polarization of

is known to be 4.8 eV, suggesting

O3

single crystals [33].

band to the conduction band [25]. The absorption of Ga<sup>2</sup>

states that are responsible for them. The bandgap of Ga<sup>2</sup>

**Figure 4.** Thermal conductivity as a function of temperature for β-Ga<sup>2</sup>

with permission from Ref. [35]. Copyright 2015, American Institute of Physics.).

*O3*

on the excitation and photoluminescence of pure and Si-doped β-Ga<sup>2</sup>

 *nanowires*

attained using the transition from the valence band.

*2.3.2. Luminescence properties of Ga2*

14 Novel Nanomaterials - Synthesis and Applications

Ga<sup>2</sup> O3 Ga<sup>2</sup> O3 is considered a weak thermal conductor compared to other semiconductors, with a thermal conductivity lower than that of other wide-bandgap materials such as SiC and GaN. Thermal conductivity in Ga<sup>2</sup> O3 varies depending on crystal direction due to its crystalline anisotropy. Directions with a smaller lattice constant have higher thermal conductivity., Ga<sup>2</sup> O3 has the highest thermal conductivity along the [010] direction and the lowest along the [100] direction. This has been measured at temperatures between 80 and 495 K (**Figure 4**). At high temperatures, thermal transport is primarily dominated by phonon scattering [35].

#### **3. Ga2 O3 nanowire growth mechanisms**

A wide variety of methods have been used to grow Ga<sup>2</sup> O3 nanowires. These include thermal oxidation [36], vapor–liquid–solid mechanism [37], pulsed laser deposition [38], sputtering [39], thermal evaporation [40–42], molecular beam epitaxy [43], laser ablation [44], arc discharge [45], carbothermal reduction [46], microwave plasma [47], metalorganic chemical vapor deposition [48] and the hydrothermal method [49, 50]. **Table 3** summarizes the advantages and disadvantages of these methods. Research is ongoing to determine the best method for growing high-quality nanowires at low cost.

#### **3.1. Thermal oxidation**

Ga<sup>2</sup> O3 nanowires have been fabricated by oxidizing GaAs in a furnace [36]. At 1050°C and at atmospheric pressure, argon gas was used to grow Ga<sup>2</sup> O3 nanowires with a nanotextured surface. As the GaAs was heated, it dissociated and the arsenic evaporated, leaving behind a liquid Ga-rich surface which was then oxidized. At high temperatures, Ga<sup>2</sup> O3 nanowires began to grow. The roughness of the surface was believed to be caused by oxygen-deficiency defects and the difference in crystal structure between GaAs and Ga<sup>2</sup> O3 .

When β-Ga<sup>2</sup> O3 nanowires are grown using this method, the material first begins as a thin film before nanowire formation if conditions are favorable. The steps in this process have been clearly described for GaAs oxidation in a furnace [51]. At high temperatures, phase separation of the GaAs occurs, and the constituents decompose. Arsenic diffuses through the substrate and evaporates from the surface, and gallium melts, forming clusters of liquid gallium on the surface. As the clusters form, the arsenic evaporates faster because it is easier for it to dissociate through the clusters than from the surface of the GaAs. Nanowire growth occurs from these clusters. Due to liquid Ga's ability to wet GaAs, the droplets are hemispherical. Oxygen in the chamber reacts with the Ga clusters to form a Ga<sup>2</sup> O3 thin film, with voids where the As evaporated. As the GaAs is converted into Ga<sup>2</sup> O3 , Ga is no longer able to wet it, resulting in


1015°C. The average lengths of the nanowires were 1–2 μm for 900°C, 1–3 μm for 1000°C, and 2–4 μm for 1015°C. The differences in nanowire length caused by temperature are more apparent in these cross-sections. The nanowires are sharp and have apparently random orientations.

the surface, particularly at higher temperatures. This causes the orientation of the nanowires to be more random, as can be seen in **Figure 6**. Another problem with growing nanowires using thermal oxidation is that at some temperatures, the surface is not uniformly covered with nanowires but instead contains gallium clusters [51]. **Figure 7** shows focused ion beam (FIB) images of these bubbles on the surface that were formed at high temperatures due to the growth mechanism. If a nanowire device is used as a sensor, this means that the effective sensing surface area is reduced. The clusters could also negatively affect the performance of devices in other ways. This method is simple and inexpensive, but it does not allow precise

The VLS mechanism (**Figure 8**) is commonly used to grow nanowires. It requires the presence of impurities, which act as catalysts at the sites where nanowires will be grown [29]. The choice of impurity affects the diameter and growth direction of the nanowires and thus is an important growth parameter. Being impurities, the catalysts become part of the material and

**Figure 7.** (a) SEM image of a Ga cluster on the surface of n-type GaAs oxidized at 942°C, (b) Closeup of the same cluster

after etching with a focused ion beam, showing nanowires surrounding the cluster but none on top of it.

O3

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

film on

17

One problem seen with furnace oxidation is buckling and delamination of the Ga<sup>2</sup>

**Figure 6.** SEM image of cross-sections of GaAs oxidized at (a) 900, (b) 1000, (c) 1015°C.

control of nanowire growth, hence their lack of directionality.

**3.2. Vapor–liquid–solid (VLS) approach**

**Table 3.** Advantages and disadvantages of β-Ga<sup>2</sup> O3 NWs growth techniques.

droplets with large contact angles. The droplets are unable to expand laterally and are thus forced to grow in one dimension, resulting in nanowire growth.

At oxidation temperatures above 750°C, nanowires begin to appear on the surface of Ga<sup>2</sup> O3 [51]. As the growth temperature is increased, the nanowires increase in length and density. The nanowires grown at 870°C (**Figure 5a**) were shorter and less dense than the ones grown at 900 and 942°C (**Figure 5b** and **c**). The samples were obtained by heating n-type (111) silicon-doped GaAs. **Figure 6** shows SEM images for nanowire growth at temperatures of 900, 1000 and

**Figure 5.** SEM images of n-type GaAs oxidized at (a) 870, (b) 900, (c) 942°C.

**Figure 6.** SEM image of cross-sections of GaAs oxidized at (a) 900, (b) 1000, (c) 1015°C.

1015°C. The average lengths of the nanowires were 1–2 μm for 900°C, 1–3 μm for 1000°C, and 2–4 μm for 1015°C. The differences in nanowire length caused by temperature are more apparent in these cross-sections. The nanowires are sharp and have apparently random orientations.

One problem seen with furnace oxidation is buckling and delamination of the Ga<sup>2</sup> O3 film on the surface, particularly at higher temperatures. This causes the orientation of the nanowires to be more random, as can be seen in **Figure 6**. Another problem with growing nanowires using thermal oxidation is that at some temperatures, the surface is not uniformly covered with nanowires but instead contains gallium clusters [51]. **Figure 7** shows focused ion beam (FIB) images of these bubbles on the surface that were formed at high temperatures due to the growth mechanism. If a nanowire device is used as a sensor, this means that the effective sensing surface area is reduced. The clusters could also negatively affect the performance of devices in other ways. This method is simple and inexpensive, but it does not allow precise control of nanowire growth, hence their lack of directionality.

#### **3.2. Vapor–liquid–solid (VLS) approach**

droplets with large contact angles. The droplets are unable to expand laterally and are thus

NWs growth techniques.

At oxidation temperatures above 750°C, nanowires begin to appear on the surface of Ga<sup>2</sup>

[51]. As the growth temperature is increased, the nanowires increase in length and density. The nanowires grown at 870°C (**Figure 5a**) were shorter and less dense than the ones grown at 900 and 942°C (**Figure 5b** and **c**). The samples were obtained by heating n-type (111) silicon-doped GaAs. **Figure 6** shows SEM images for nanowire growth at temperatures of 900, 1000 and

O3

forced to grow in one dimension, resulting in nanowire growth.

**Growth Mechanisms Advantages Disadvantages**

many materials

impurity levels

Arc-discharge Faster growth • Low quality NWs

variety of structures

nanowires

Thermal evaporation Inexpensive and compatible with

Carbothermal reduction Simple process for growing a wide

Microwave plasma Can grow a wide variety of

**Table 3.** Advantages and disadvantages of β-Ga<sup>2</sup>

Metalorganic chemical vapor

deposition

Molecular beam epitaxy high deposition rate and low

16 Novel Nanomaterials - Synthesis and Applications

Thermal oxidation Simple and inexpensive Low degree of control over growth Vapor–liquid–solid Good control of growth Can interfere with doping levels Pulsed laser deposition High deposition rate More defects and dislocations RF magnetron sputtering Non-stoichiometric deposition • Incorporation of gas into material

Laser ablation High quality film Expensive and require longer time

Hydrothermal method Simple, cheap and efficient Require precise temperature control Sol-gel method Simple and inexpensive The thickness is nonuniform

O3

• It produces low quality NWs

More surface roughness at higher

Contamination of material

temperatures

• Many defects

High levels of impurities

Radial nonuniformity

High growth rate Expensive and uses toxic gases

**Figure 5.** SEM images of n-type GaAs oxidized at (a) 870, (b) 900, (c) 942°C.

The VLS mechanism (**Figure 8**) is commonly used to grow nanowires. It requires the presence of impurities, which act as catalysts at the sites where nanowires will be grown [29]. The choice of impurity affects the diameter and growth direction of the nanowires and thus is an important growth parameter. Being impurities, the catalysts become part of the material and

**Figure 7.** (a) SEM image of a Ga cluster on the surface of n-type GaAs oxidized at 942°C, (b) Closeup of the same cluster after etching with a focused ion beam, showing nanowires surrounding the cluster but none on top of it.

**3.5. Thermal evaporation**

Thermal evaporation has been used to grow Ga<sup>2</sup>

core and a different material (ZnO and SnO<sup>2</sup>

sensor. The entire structure consisted of Ga<sup>2</sup>

sensors as compared to simply using Ga<sup>2</sup>

substances present in the chamber.

**3.6. Molecular beam epitaxy (MBE)**

O3

sapphire using both a Ga<sup>2</sup>

source, Ga<sup>2</sup>

and high vacuum conditions.

**3.7. Laser ablation**

grow Ga<sup>2</sup>

Ga<sup>2</sup> O3 O3

of 121°C. Rather than a Ga<sup>2</sup>

the material grown into Ga<sup>2</sup>

nanowires.

decompose it into Ga<sup>2</sup>

Ga<sup>2</sup> O3

Molecular beam epitaxy is another option for growing Ga<sup>2</sup>

O3

O and O2

completely oxidize the Ga atoms, while with the Ga<sup>2</sup>

Laser ablation involves the use of a laser on a Ga<sup>2</sup>

faults, attesting to the high quality of material that can be grown.

O3

O3

O3

tions [41, 42] at lower temperatures. In both cases, the nanowires were grown using a two-step process involving thermal evaporation and atomic layer deposition. This resulted in a Ga<sup>2</sup>

rial on the outside. The addition of the second material improves the performance of the gas

at which the sensor can operate. This method of growth is inexpensive and compatible with a wide variety of materials, but the evaporation process often leads to contamination by other

than 5 × 10−10 torr, and oxygen gas was used to protect the substrate from ions. For the Ga source, Ga was evaporated from an effusion cell through an oxygen plasma, while for the

substrates [15]. The rate of growth was highly dependent on the substrate orientation, with the quickest growth occurring on the (010) and (310) planes and the growth on the (100) plane being very slow. At lower temperatures, surface roughness decreased. Molecular beam epitaxy has a high deposition rate and gives low impurity levels, but requires high temperatures

O3

them onto a substrate. Parameters such as pulse width, pulse power, distance between target and substrate, and the ablation target can all be controlled, giving this method high versatility and control over the material grown. The nanostructures it produces are high-quality but it takes a long time to grow them and it is expensive. For example, it took 5 hours to grow nanowires with diameters between 15 and 50 nm and lengths of several micrometers [52]. When a single nanowire was examined, it was found to be free of dislocations and stacking

Laser ablation can also be used as part of a two-step process along with solution refluxing to

using the laser was only 20 minutes, the reflux process took a day, and calcination to convert

nanowires [53]. This approach allows nanowires to grow at the low temperature

target, a high-purity Ga plate was used. While the time spent

took 18 hours, so this method requires more time to fabricate

oxide layer was thicker. Molecular beam epitaxy has also been used to grow Ga<sup>2</sup>

O3

O3

nanowires, mainly for gas-sensing applica-

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

) as the shell on the outside, which acted as the

nanowires that were coated with another mate-

nanowires, mainly by lowering the temperature

. It has used to grow Ga<sup>2</sup>

source, at lower temperatures, the

target to extract its contents and deposit

O3

on β-Ga<sup>2</sup>

O3

O3

source and an elemental Ga source [43]. The pressure was less

. Growth using elemental Ga required high temperatures to

was evaporated from an iridium crucible with an oxygen plasma to

O3

O3

19

O3 on

**Figure 8.** Vapor-liquid-solid growth of semiconductor nanowires.

can interfere with doping. Commonly used catalysts include Au, Ni, and Fe. These catalysts are commonly deposited on the substrate using the sol–gel method [37]. Nanowire growth occurs at the locations where the catalysts are deposited.

#### **3.3. Pulsed laser deposition (PLD)**

Pulsed laser deposition has been used to grow Ga<sup>2</sup> O3 nanostructures on α-Al<sup>2</sup> O3 at 900°C. The nanostructure was doped with Cr by alternating laser pulses to deposit both Ga<sup>2</sup> O3 and Cr [38]. This resulted in nanowires with a controlled length and diameter. The nanowires had tips larger than their bodies. PLD has a high deposition rate and is a low-cost growth method. There are many parameters involved in optimizing it for a specific process, such as the thickness of the catalyst, sample orientation, pulse intensity and number of pulses, and distance between laser and material, giving good control over the process. PLD produces more defects and dislocations due to the high energies used.

#### **3.4. RF (radio frequency) magnetron sputtering**

RF magnetron sputtering is another way of growing Ga<sup>2</sup> O3 nanowires [39]. If doping is needed, Ga<sup>2</sup> O3 powder can be mixed with the dopant (also in powder form) to form a sputter target. The fact that it produces non-stoichiometric films in an oxygen-deficient atmosphere is important in growing nanowires using Ga seeds. Growth time affects film thickness, which in turn affects the nanowires. The dopant only appears at the tip of the nanowires. Sputtering can be performed at room temperature, is low-cost, and provides good step coverage, but the films produced are not very high-quality because of damage caused during the process.

#### **3.5. Thermal evaporation**

Thermal evaporation has been used to grow Ga<sup>2</sup> O3 nanowires, mainly for gas-sensing applications [41, 42] at lower temperatures. In both cases, the nanowires were grown using a two-step process involving thermal evaporation and atomic layer deposition. This resulted in a Ga<sup>2</sup> O3 core and a different material (ZnO and SnO<sup>2</sup> ) as the shell on the outside, which acted as the sensor. The entire structure consisted of Ga<sup>2</sup> O3 nanowires that were coated with another material on the outside. The addition of the second material improves the performance of the gas sensors as compared to simply using Ga<sup>2</sup> O3 nanowires, mainly by lowering the temperature at which the sensor can operate. This method of growth is inexpensive and compatible with a wide variety of materials, but the evaporation process often leads to contamination by other substances present in the chamber.

#### **3.6. Molecular beam epitaxy (MBE)**

Molecular beam epitaxy is another option for growing Ga<sup>2</sup> O3 . It has used to grow Ga<sup>2</sup> O3 on sapphire using both a Ga<sup>2</sup> O3 source and an elemental Ga source [43]. The pressure was less than 5 × 10−10 torr, and oxygen gas was used to protect the substrate from ions. For the Ga source, Ga was evaporated from an effusion cell through an oxygen plasma, while for the Ga<sup>2</sup> O3 source, Ga<sup>2</sup> O3 was evaporated from an iridium crucible with an oxygen plasma to decompose it into Ga<sup>2</sup> O and O2 . Growth using elemental Ga required high temperatures to completely oxidize the Ga atoms, while with the Ga<sup>2</sup> O3 source, at lower temperatures, the oxide layer was thicker. Molecular beam epitaxy has also been used to grow Ga<sup>2</sup> O3 on β-Ga<sup>2</sup> O3 substrates [15]. The rate of growth was highly dependent on the substrate orientation, with the quickest growth occurring on the (010) and (310) planes and the growth on the (100) plane being very slow. At lower temperatures, surface roughness decreased. Molecular beam epitaxy has a high deposition rate and gives low impurity levels, but requires high temperatures and high vacuum conditions.

#### **3.7. Laser ablation**

can interfere with doping. Commonly used catalysts include Au, Ni, and Fe. These catalysts are commonly deposited on the substrate using the sol–gel method [37]. Nanowire growth

[38]. This resulted in nanowires with a controlled length and diameter. The nanowires had tips larger than their bodies. PLD has a high deposition rate and is a low-cost growth method. There are many parameters involved in optimizing it for a specific process, such as the thickness of the catalyst, sample orientation, pulse intensity and number of pulses, and distance between laser and material, giving good control over the process. PLD produces more defects

target. The fact that it produces non-stoichiometric films in an oxygen-deficient atmosphere is important in growing nanowires using Ga seeds. Growth time affects film thickness, which in turn affects the nanowires. The dopant only appears at the tip of the nanowires. Sputtering can be performed at room temperature, is low-cost, and provides good step coverage, but the films produced are not very high-quality because of damage caused during the process.

nanostructure was doped with Cr by alternating laser pulses to deposit both Ga<sup>2</sup>

O3

nanostructures on α-Al<sup>2</sup>

O3

powder can be mixed with the dopant (also in powder form) to form a sputter

O3

nanowires [39]. If doping is

at 900°C. The

and Cr

O3

occurs at the locations where the catalysts are deposited.

**Figure 8.** Vapor-liquid-solid growth of semiconductor nanowires.

Pulsed laser deposition has been used to grow Ga<sup>2</sup>

and dislocations due to the high energies used.

**3.4. RF (radio frequency) magnetron sputtering**

RF magnetron sputtering is another way of growing Ga<sup>2</sup>

**3.3. Pulsed laser deposition (PLD)**

18 Novel Nanomaterials - Synthesis and Applications

needed, Ga<sup>2</sup>

O3

Laser ablation involves the use of a laser on a Ga<sup>2</sup> O3 target to extract its contents and deposit them onto a substrate. Parameters such as pulse width, pulse power, distance between target and substrate, and the ablation target can all be controlled, giving this method high versatility and control over the material grown. The nanostructures it produces are high-quality but it takes a long time to grow them and it is expensive. For example, it took 5 hours to grow nanowires with diameters between 15 and 50 nm and lengths of several micrometers [52]. When a single nanowire was examined, it was found to be free of dislocations and stacking faults, attesting to the high quality of material that can be grown.

Laser ablation can also be used as part of a two-step process along with solution refluxing to grow Ga<sup>2</sup> O3 nanowires [53]. This approach allows nanowires to grow at the low temperature of 121°C. Rather than a Ga<sup>2</sup> O3 target, a high-purity Ga plate was used. While the time spent using the laser was only 20 minutes, the reflux process took a day, and calcination to convert the material grown into Ga<sup>2</sup> O3 took 18 hours, so this method requires more time to fabricate Ga<sup>2</sup> O3 nanowires.

#### **3.8. Arc discharge**

Arc discharge involves applying a direct-current arc voltage across an anode and cathode, both usually made of graphite, in an inert gas. The anode usually has a hole drilled into it to hold powders, which are used to grow the material desired, and the energy of the plasma generated is used to create nanostructures. In one study, the anode was a graphite electrode filled with GaN, graphite, and Ni powders while the cathode was a slightly larger graphite rod [45]. The growth took place in a helium environment. The Ni powder acted as a catalyst for the reaction. Nanostructures were deposited on the cathode, and most of them were nanorods. The nanorods were most commonly oriented in the [111] and [200] directions. Defects such as microtwins and stacking faults were found on some of the nanorods, suggesting that this method of growth does not produce very high-quality nanostructures.

for growing Ga<sup>2</sup>

SiO<sup>2</sup>

Ga<sup>2</sup> O3

tal (α-Ga<sup>2</sup>

O3

and Al<sup>2</sup>

cursor and H<sup>2</sup>

O3

**3.12. Hydrothermal Method**

considered the simplest method for growing β-Ga<sup>2</sup>

when GaN reacts with different alkalis such as NaOH, NH<sup>4</sup>

) and monoclinic crystal (β-Ga<sup>2</sup>

O and N<sup>2</sup>

**3.10. Microwave plasma**

O3

are high levels of carbon impurities in the final material.

but use of a plasma is associated with radial nonuniformity.

**3.11. Metalorganic chemical vapor deposition (MOCVD)**

Metalorganic chemical vapor deposition has been used to grow Ga<sup>2</sup>

morphology of the nanostructures can be controlled. With a longer pulse of H<sup>2</sup>

high-quality films, but it is expensive and requires the use of highly toxic gases.

nanostructures and can be used to create a wide variety of structures, there

, CH<sup>4</sup>

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

nanostructures on Si/

O and using O2

CO<sup>3</sup>

[49]. The

O, the nanow-

, the

O3

[49, 50]. The crystal of GaOOH is obtained

OH, KOH, and Na<sup>2</sup>

). The morphologies of the gallium oxide were

, and O2 . 21

The microwave plasma method of growing nanowires involves using low-melting metals as a solvent medium for the growth of nanowires [47]. Droplets or thin films of gallium on the substrate cause nucleation of nanowires with high density (>1011 cm−2). This technique was successfully used to grow unique "nanopaintbrushes." A variety of substrates were used with

After thin films developed, gallium droplets were added to grow nanowires. Stacking faults were not observed on the nanostructures. Four different nanostructures were seen on different parts of the substrate—nanowires, nanotubes, nanorods, and nanopaintbrushes—and it was suggested that the type of nanostructure formed depended on the initial state of the nuclei on the surface. Hydrogen appeared to etch the nuclei on the surface, preventing them from growing laterally and agglomerating with other nuclei, thus promoting vertical nanowire growth. This method of nanowire growth can produce a wide variety of nanostructures,

temperature (450°C in this case). The growth process consists of alternating pulses of a Ga pre-

ires were about a micron long and had large particles at their tips. If the water pulse was shortened, the wires would taper from the tip down, being widest at the tip and becoming thinner close to the substrate. At higher deposition temperatures, more nanostructures were observed on the surface. Metalorganic chemical vapor deposition provides relatively fast growth and

to purge the chamber. By changing the pulses of H<sup>2</sup>

nanowires can be obtained by calcination of gallium oxide hydroxide (GaOOH). This is

GaOOH morphologies were changed into two different crystal structures: rhombohedral crys-

influenced by changing the temperatures [50]. **Figure 9** shows FE-SEM images of the GaOOH nanostructures grown by the hydrothermal method. Different growth temperatures were tried, such as room temperature, 50, 75, and 95°C. Temperature has a significant impact on the growth of gallium oxide. The number of shapes seen and the size of the nanowires increased as

O3

the temperature was increased. This growth method can be used at low temperatures.

O3

[48]. As in furnace oxidation, nanostructures are only observed above a certain

molten gallium and exposed to plasma in a microwave plasma reactor with H<sup>2</sup>

A similar study has been carried out using a mix of Ni and Co powder as the catalyst inside the graphite anode and using a mixture of Ar and O<sup>2</sup> instead of He [54]. The pressure was varied, and it was found that no nanowires were present under 450 torr, suggesting that oxygen gas was necessary for their formation. Furthermore, without the Ni and Co powder there was no nanowire growth either. The nanowires obtained showed twin defects and rough step edges. The surfaces of the nanowires were covered with amorphous layers, and jog defects were also seen. Arc discharge can grow nanowires very quickly, but it produces nanowires that contain many defects.

#### **3.9. Carbothermal reduction (CTR)**

Carbothermal reduction is a simple method for growing Ga<sup>2</sup> O3 nanowires that involves mixing Ga<sup>2</sup> O3 and graphite powder and heating them in a furnace in the presence of a substrate [46]. This process can result in cactus-like nanostructures on the surface of the substrate. These nanostructures were very thin and had small spherical structures at their tips. Dense nanowire growth with a radial distribution was observed all over the surface of the samples. The cactus-like structures varied in diameter from a few micrometers to several tens of micrometers, which is much larger than conventional nanostructures. Nanowires with a vertical orientation were also seen, but their tips were often bent and a few had multiple nanorods with a random orientation at their tips. They were also tapered and had rough surfaces on closer examination. Growth time was correlated with nanowire density, and nanospheres were the first structures to grow on the surface, on top of which nanowires subsequently grew.

In another study, nanowires with lengths of tens of micrometers to hundreds of micrometers were grown using a similar process [29]. Different nanostructures were associated with different growth temperatures and substrates, with nanowires found at 900°C on Si substrates and nanosheets found at 800°C on quartz substrates. These nanosheets were large, with areas in the order of several tens of square micrometers. Their diffraction patterns were similar to that of thin films. When Si substrates were used at 800°C, Ga<sup>2</sup> O3 nanoribbons were observed. At high temperatures, the nanowires had sawtooth-like structures, which was believed to be due to oxygen deficiencies during growth. Although carbothermal reduction is a simple method for growing Ga<sup>2</sup> O3 nanostructures and can be used to create a wide variety of structures, there are high levels of carbon impurities in the final material.

#### **3.10. Microwave plasma**

**3.8. Arc discharge**

20 Novel Nanomaterials - Synthesis and Applications

that contain many defects.

Ga<sup>2</sup> O3

**3.9. Carbothermal reduction (CTR)**

on top of which nanowires subsequently grew.

of thin films. When Si substrates were used at 800°C, Ga<sup>2</sup>

Arc discharge involves applying a direct-current arc voltage across an anode and cathode, both usually made of graphite, in an inert gas. The anode usually has a hole drilled into it to hold powders, which are used to grow the material desired, and the energy of the plasma generated is used to create nanostructures. In one study, the anode was a graphite electrode filled with GaN, graphite, and Ni powders while the cathode was a slightly larger graphite rod [45]. The growth took place in a helium environment. The Ni powder acted as a catalyst for the reaction. Nanostructures were deposited on the cathode, and most of them were nanorods. The nanorods were most commonly oriented in the [111] and [200] directions. Defects such as microtwins and stacking faults were found on some of the nanorods, suggesting that this

A similar study has been carried out using a mix of Ni and Co powder as the catalyst inside

varied, and it was found that no nanowires were present under 450 torr, suggesting that oxygen gas was necessary for their formation. Furthermore, without the Ni and Co powder there was no nanowire growth either. The nanowires obtained showed twin defects and rough step edges. The surfaces of the nanowires were covered with amorphous layers, and jog defects were also seen. Arc discharge can grow nanowires very quickly, but it produces nanowires

 and graphite powder and heating them in a furnace in the presence of a substrate [46]. This process can result in cactus-like nanostructures on the surface of the substrate. These nanostructures were very thin and had small spherical structures at their tips. Dense nanowire growth with a radial distribution was observed all over the surface of the samples. The cactus-like structures varied in diameter from a few micrometers to several tens of micrometers, which is much larger than conventional nanostructures. Nanowires with a vertical orientation were also seen, but their tips were often bent and a few had multiple nanorods with a random orientation at their tips. They were also tapered and had rough surfaces on closer examination. Growth time was correlated with nanowire density, and nanospheres were the first structures to grow on the surface,

In another study, nanowires with lengths of tens of micrometers to hundreds of micrometers were grown using a similar process [29]. Different nanostructures were associated with different growth temperatures and substrates, with nanowires found at 900°C on Si substrates and nanosheets found at 800°C on quartz substrates. These nanosheets were large, with areas in the order of several tens of square micrometers. Their diffraction patterns were similar to that

high temperatures, the nanowires had sawtooth-like structures, which was believed to be due to oxygen deficiencies during growth. Although carbothermal reduction is a simple method

O3

O3

instead of He [54]. The pressure was

nanowires that involves mixing

nanoribbons were observed. At

method of growth does not produce very high-quality nanostructures.

the graphite anode and using a mixture of Ar and O<sup>2</sup>

Carbothermal reduction is a simple method for growing Ga<sup>2</sup>

The microwave plasma method of growing nanowires involves using low-melting metals as a solvent medium for the growth of nanowires [47]. Droplets or thin films of gallium on the substrate cause nucleation of nanowires with high density (>1011 cm−2). This technique was successfully used to grow unique "nanopaintbrushes." A variety of substrates were used with molten gallium and exposed to plasma in a microwave plasma reactor with H<sup>2</sup> , CH<sup>4</sup> , and O2 . After thin films developed, gallium droplets were added to grow nanowires. Stacking faults were not observed on the nanostructures. Four different nanostructures were seen on different parts of the substrate—nanowires, nanotubes, nanorods, and nanopaintbrushes—and it was suggested that the type of nanostructure formed depended on the initial state of the nuclei on the surface. Hydrogen appeared to etch the nuclei on the surface, preventing them from growing laterally and agglomerating with other nuclei, thus promoting vertical nanowire growth. This method of nanowire growth can produce a wide variety of nanostructures, but use of a plasma is associated with radial nonuniformity.

#### **3.11. Metalorganic chemical vapor deposition (MOCVD)**

Metalorganic chemical vapor deposition has been used to grow Ga<sup>2</sup> O3 nanostructures on Si/ SiO<sup>2</sup> and Al<sup>2</sup> O3 [48]. As in furnace oxidation, nanostructures are only observed above a certain temperature (450°C in this case). The growth process consists of alternating pulses of a Ga precursor and H<sup>2</sup> O and N<sup>2</sup> to purge the chamber. By changing the pulses of H<sup>2</sup> O and using O2 , the morphology of the nanostructures can be controlled. With a longer pulse of H<sup>2</sup> O, the nanowires were about a micron long and had large particles at their tips. If the water pulse was shortened, the wires would taper from the tip down, being widest at the tip and becoming thinner close to the substrate. At higher deposition temperatures, more nanostructures were observed on the surface. Metalorganic chemical vapor deposition provides relatively fast growth and high-quality films, but it is expensive and requires the use of highly toxic gases.

#### **3.12. Hydrothermal Method**

Ga<sup>2</sup> O3 nanowires can be obtained by calcination of gallium oxide hydroxide (GaOOH). This is considered the simplest method for growing β-Ga<sup>2</sup> O3 [49, 50]. The crystal of GaOOH is obtained when GaN reacts with different alkalis such as NaOH, NH<sup>4</sup> OH, KOH, and Na<sup>2</sup> CO<sup>3</sup> [49]. The GaOOH morphologies were changed into two different crystal structures: rhombohedral crystal (α-Ga<sup>2</sup> O3 ) and monoclinic crystal (β-Ga<sup>2</sup> O3 ). The morphologies of the gallium oxide were influenced by changing the temperatures [50]. **Figure 9** shows FE-SEM images of the GaOOH nanostructures grown by the hydrothermal method. Different growth temperatures were tried, such as room temperature, 50, 75, and 95°C. Temperature has a significant impact on the growth of gallium oxide. The number of shapes seen and the size of the nanowires increased as the temperature was increased. This growth method can be used at low temperatures.

*3.13.1. Aluminum oxide (Al2*

a few layers of α-Ga<sup>2</sup>

better-matched [57].

GaN. ε-Ga<sup>2</sup>

Ga<sup>2</sup> O3

*3.13.2. Gallium Nitride (GaN)*

The lattices of GaN and Ga<sup>2</sup>

O3

*3.13.3. Silicon carbide (SiC)*

because it is semiconducting.

*3.13.4. Gallium oxide (Ga2*

the high cost of Ga<sup>2</sup>

high [55]. When ε-Ga<sup>2</sup>

Al<sup>2</sup> O3 *O3 )*

O3

4.0 eV, so there is almost no offset in the conduction band when Ga<sup>2</sup>

Depending on the orientation used, lattice mismatch between SiC and Ga<sup>2</sup>

insulator, this provides electrical isolation for Ga<sup>2</sup>

O3

O3

*O3 )*

O3

*3.13.5. Magnesium aluminum oxide (MgAl6*

As an alternative to using Al<sup>2</sup>

types of point defects in the Ga<sup>2</sup>

Gallium oxide is also an occasionally used substrate for Ga<sup>2</sup>

O3

O3

formed by pairs of gallium vacancies and gallium-oxygen vacancies [61].

Ga vacancies, and gallium-oxygen vacancy pairs. When MgAl<sup>6</sup>

the monoclinic structure of β- Ga<sup>2</sup>

is one of the most commonly used materials for Ga<sup>2</sup>

O3

O3

O3

was grown on (001) SiC, morphological and structural disorder was

O3

O3

were observed, such as oxygen vacancies, interstitial Ga,

on MgAl<sup>6</sup>

O10 is used, the donor band is

O3

O3

O3

keeps its morphology and has a low defect density when grown on GaN [55].

between the two materials. Lattice mismatches of 4.2% and 10.7% have been reported [44]. Sapphire is fairly inexpensive and available in different-sized wafers but there are always

between a sapphire substrate and β-Ga<sup>2</sup>

reported that there is dissimilarity between the corundum crystal structure of sapphire and

[56], but α-Ga<sup>2</sup>

are well matched, and thus Ga<sup>2</sup>

Furthermore, GaN has a wide bandgap, so using it as a substrate means that a photodetector will be less sensitive to longer wavelengths [58]. Both materials also have electron affinities of

very high, due to the difference between the crystal structures of the two materials, and the

was observed. Because it is fairly easy to dope SiC, a p-n junction can be created using SiC as a substrate [59]. Using SiC as a substrate also enables fabrication of vertically structured devices

lattice mismatch, high-quality crystal growth is possible [15]. To fully oxidize the gallium sub-oxides during pulsed laser deposition, an oxygen-radical-rich atmosphere is needed [60]. Without oxygen radicals, gallium tends to sublimate and its surfaces tend to be rough. Also,

*O10)*

has high thermal stability and a lattice mismatch of only 2.9%, less than that of Al<sup>2</sup>

Chemically, the material is not very different from sapphire. When using MgAl<sup>6</sup>

wafers makes them less attractive as a growth substrate.

, it is also possible to grow Ga<sup>2</sup>

was polycrystalline. Interestingly, in the [111] direction, high-quality Ga<sup>2</sup>

growth. As sapphire is an

is also commonly grown on

is grown on top of GaN.

can be fairly

growth

O3

O10. MgAl<sup>6</sup>

O3 [44].

O10, several

O<sup>10</sup>

O3

growth. Due to the lack of

[55]. It has also been

23

. However, there is still lattice mismatch

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

has a corundum structure, so it is

O3

**Figure 9.** Effect of reaction temperature on the morphology of GaOOH nanostructures. Ref. [50] is licensed CC BY 4.0.

#### **3.13. Ga2O3 growth substrates**

Different substrates have been used to examine and produce low-cost and high-quality nanowires. The next sub-sections provide an overview of the substrates that have been utilized recently for the growth of gallium oxide nanowires. **Table 4** describes some advantages and disadvantages of different substrates on which β-Ga<sup>2</sup> O3 nanowires are grown.


**Table 4.**An overview of some substrates advantages and disadvantages were applied of β-Ga<sup>2</sup> O3 NWs growth techniques.

#### *3.13.1. Aluminum oxide (Al2 O3 )*

Al<sup>2</sup> O3 is one of the most commonly used materials for Ga<sup>2</sup> O3 growth. As sapphire is an insulator, this provides electrical isolation for Ga<sup>2</sup> O3 . However, there is still lattice mismatch between the two materials. Lattice mismatches of 4.2% and 10.7% have been reported [44]. Sapphire is fairly inexpensive and available in different-sized wafers but there are always a few layers of α-Ga<sup>2</sup> O3 between a sapphire substrate and β-Ga<sup>2</sup> O3 [55]. It has also been reported that there is dissimilarity between the corundum crystal structure of sapphire and the monoclinic structure of β- Ga<sup>2</sup> O3 [56], but α-Ga<sup>2</sup> O3 has a corundum structure, so it is better-matched [57].

#### *3.13.2. Gallium Nitride (GaN)*

The lattices of GaN and Ga<sup>2</sup> O3 are well matched, and thus Ga<sup>2</sup> O3 is also commonly grown on GaN. ε-Ga<sup>2</sup> O3 keeps its morphology and has a low defect density when grown on GaN [55]. Furthermore, GaN has a wide bandgap, so using it as a substrate means that a photodetector will be less sensitive to longer wavelengths [58]. Both materials also have electron affinities of 4.0 eV, so there is almost no offset in the conduction band when Ga<sup>2</sup> O3 is grown on top of GaN.

#### *3.13.3. Silicon carbide (SiC)*

**3.13. Ga2O3 growth substrates**

22 Novel Nanomaterials - Synthesis and Applications

O<sup>10</sup> Thermally stable

MgO It has wide bandgap and can shift

wavelengths

Al<sup>2</sup>

Ga<sup>2</sup>

MgAl<sup>6</sup>

and disadvantages of different substrates on which β-Ga<sup>2</sup>

Small lattice mismatch

the absorption edge to shorter

Si Abundant & inexpensive Lattice mismatch

**Table 4.**An overview of some substrates advantages and disadvantages were applied of β-Ga<sup>2</sup>

**Materials Advantages Disadvantages**

O3 Inexpensive and available Lattice mismatch GaN Common and well-matched Few defects observed

Different substrates have been used to examine and produce low-cost and high-quality nanowires. The next sub-sections provide an overview of the substrates that have been utilized recently for the growth of gallium oxide nanowires. **Table 4** describes some advantages

**Figure 9.** Effect of reaction temperature on the morphology of GaOOH nanostructures. Ref. [50] is licensed CC BY 4.0.

SiC Large-scale fabrication processes Lattice mismatch leads to high morphological and

O3 Lack of lattice mismatch Expensive and requires atmosphere with oxygen

O3

structural disorder

Oxygen vacancies

radicals to obtain a smooth surface.

vacancies and Ga-O vacancy pairs

Require annealing to improve crystallinity

nanowires are grown.

Point defects: Oxygen vacancies, interstitial Ga, Ga

O3

NWs growth techniques.

Depending on the orientation used, lattice mismatch between SiC and Ga<sup>2</sup> O3 can be fairly high [55]. When ε-Ga<sup>2</sup> O3 was grown on (001) SiC, morphological and structural disorder was very high, due to the difference between the crystal structures of the two materials, and the Ga<sup>2</sup> O3 was polycrystalline. Interestingly, in the [111] direction, high-quality Ga<sup>2</sup> O3 growth was observed. Because it is fairly easy to dope SiC, a p-n junction can be created using SiC as a substrate [59]. Using SiC as a substrate also enables fabrication of vertically structured devices because it is semiconducting.

#### *3.13.4. Gallium oxide (Ga2 O3 )*

Gallium oxide is also an occasionally used substrate for Ga<sup>2</sup> O3 growth. Due to the lack of lattice mismatch, high-quality crystal growth is possible [15]. To fully oxidize the gallium sub-oxides during pulsed laser deposition, an oxygen-radical-rich atmosphere is needed [60]. Without oxygen radicals, gallium tends to sublimate and its surfaces tend to be rough. Also, the high cost of Ga<sup>2</sup> O3 wafers makes them less attractive as a growth substrate.

#### *3.13.5. Magnesium aluminum oxide (MgAl6 O10)*

As an alternative to using Al<sup>2</sup> O3 , it is also possible to grow Ga<sup>2</sup> O3 on MgAl<sup>6</sup> O10. MgAl<sup>6</sup> O<sup>10</sup> has high thermal stability and a lattice mismatch of only 2.9%, less than that of Al<sup>2</sup> O3 [44]. Chemically, the material is not very different from sapphire. When using MgAl<sup>6</sup> O10, several types of point defects in the Ga<sup>2</sup> O3 were observed, such as oxygen vacancies, interstitial Ga, Ga vacancies, and gallium-oxygen vacancy pairs. When MgAl<sup>6</sup> O10 is used, the donor band is formed by pairs of gallium vacancies and gallium-oxygen vacancies [61].

#### *3.13.6. Magnesium oxide (MgO)*

MgO has also been used as a substrate for growing Ga<sup>2</sup> O3 . It tends to absorb water from the air, which affects the structural properties of the Ga<sup>2</sup> O3 on top [57]. When Ga<sup>2</sup> O3 is grown on MgO, it is initially amorphous and requires annealing to achieve good crystallinity [16]. However, this also decreases the bandgap, making the material less sensitive to UV light. The Ga<sup>2</sup> O3 film is also highly resistive, although it remains n-type due to oxygen vacancies.

**4.4. Titanium/aluminum (Ti/Al)**

**4.5. Graphene**

**5. Ga2**

These Ga<sup>2</sup>

such as Ga<sup>2</sup>

Ga<sup>2</sup> O3 **O3**

O3

O3

**5.1. UV photodetectors**

through an interdigitated shadow mask onto a Ga<sup>2</sup>

Graphene electrodes have also been deposited on a β-Ga<sup>2</sup>

responsivity, larger Ilight/Idark ratio, and a shorter rise time.

is still a great challenge. In spite of these difficulties, various Ga<sup>2</sup>

 **nanowire applications**

found to be non-ohmic, which was the goal.

Electron beam evaporation was used to deposit Ti and Al onto the back side of a Ga<sup>2</sup>

diode [66]. On the other side of the diode, Schottky contacts of Au were also deposited. These contacts were rectifying. A contact consisting of alternating layers, Ti/Al/Ti/Al, was deposited

phene was grown by chemical vapor deposition on Cu and then transferred to the substrate. The graphene was transparent and thus allowed most incident light to reach the contact area. It also acted as a channel when the electrons and the holes were separated. On the back of the substrate, a Au/Ti electrode was deposited. A device with a gold electrode on top was also fabricated. The device with the graphene electrode showed a larger photoresponse, higher

nanowires, nanobelts and nanorods have been extensively studied in recent years.

nanowire-based FETs, gas sensors and UV photodetectors are reported.

The spectrum of radiation between 200 and 280 nm is called solar-blind due to the absorption of solar radiation in that spectrum by the ozone layer. Because of this lack of background noise, it is possible to detect very weak signals in this spectrum. In recent years, the need for solarblind UV photodetectors has increased, due to their potential in several applications such as flame detection and missile warning [68]. When photons strike the surface of a UV photodetector, they generate electron–hole pairs. These electron–hole pairs change the conductivity of the material, which shows that light has been detected. For a device to detect signals in the UV spectrum, a material with a wide bandgap should be used. It is important that the material

Photomultiplier tubes are used for UV detection in the solar-blind region, but they have several problems. They are bulky, require high bias voltages, and have high leakage current. Siliconbased photodiodes are also used, but due to their narrow bandgap, they require additional

detect UV light but is not affected by light from other parts of the optical spectrum.

nanostructure-based devices are limited to a single nanowire, and the integration of individual devices on a single chip is still a real challenge. Another crucial issue is controlled doping in nanostructures and the formation of high-quality ohmic contacts between nanostructures and electrodes. Therefore, the construction of a device via a simple and cost-effective method

nanostructures provide a path to a new generation of devices, but most Ga<sup>2</sup>

O3

O3

O3

photodetector [4]. The contacts were

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

/SiC photodetector [67]. The gra-

O3

photo-

25

O3

nanowire applications

#### *3.13.7. Silicon (Si)*

Ga<sup>2</sup> O3 can also be grown on silicon. It has been deposited on (100) Si [57]. A minimum temperature of 600°C was required for crystallization to occur. Because silicon, like silicon carbide, can be easily doped p-type, growing Ga<sup>2</sup> O3 on Si is an option to make a p-n junction. Ga<sup>2</sup> O3 grown on (100) Si using molecular beam epitaxy has good crystallinity and produces smooth surfaces [62].

#### **4. Contacts for Ga2 O3**

Contacts to Ga<sup>2</sup> O3 can be either Schottky or ohmic. To make an ohmic contact, the barrier height of the metal should be low. Gold and titanium are the most commonly used metals for contacts. Metal contacts can be deposited using shadow masks, which involve depositing metal onto the substrate while it is covered by a thin plate with gaps in the desired shape.

#### **4.1. Titanium/gold (Ti/Au)**

Ti/Au interdigitated electrodes have been deposited on both β-Ga<sup>2</sup> O3 [63] and chromiumdoped Ga<sup>2</sup> O3 nanowire thin films [38]. In both cases, a shadow mask was used and the contacts exhibited ohmic behavior. It was suggested that this was due to the abundance of surface states in Ga<sup>2</sup> O3 , which makes it easy for carriers to tunnel.

#### **4.2. Gold (Au)**

A shadow mask has been used to deposit Au onto the back of a β-Ga<sup>2</sup> O3 /Si heterojunction photodetector [62]. The deposition method was radio frequency magnetron sputtering. A Ti/Au contact was also sputtered on top of the β-Ga<sup>2</sup> O3 . The contacts formed were ohmic. Standard photolithography techniques and liftoff have been used to deposit interdigitated Au electrodes on β-Ga<sup>2</sup> O3 [64]. Carrier multiplication that occurred in the area under the electrodes resulted in high responsivity and quantum efficiency.

#### **4.3. Indium (In)**

Thermal evaporation was used to deposit an In electrode on the SnO<sup>2</sup> part of a β-Ga<sup>2</sup> O3 / SnO<sup>2</sup> heterojunction photodetector [65]. A bilayer of Ti and Au was also deposited on the β-Ga<sup>2</sup> O3 thin film using evaporation. The I-V curve of the device showed that the contacts were rectifying.

#### **4.4. Titanium/aluminum (Ti/Al)**

Electron beam evaporation was used to deposit Ti and Al onto the back side of a Ga<sup>2</sup> O3 photodiode [66]. On the other side of the diode, Schottky contacts of Au were also deposited. These contacts were rectifying. A contact consisting of alternating layers, Ti/Al/Ti/Al, was deposited through an interdigitated shadow mask onto a Ga<sup>2</sup> O3 photodetector [4]. The contacts were found to be non-ohmic, which was the goal.

#### **4.5. Graphene**

*3.13.6. Magnesium oxide (MgO)*

24 Novel Nanomaterials - Synthesis and Applications

easily doped p-type, growing Ga<sup>2</sup>

O3

**4.1. Titanium/gold (Ti/Au)**

O3

O3

**4. Contacts for Ga2**

Contacts to Ga<sup>2</sup>

doped Ga<sup>2</sup>

states in Ga<sup>2</sup>

**4.2. Gold (Au)**

on β-Ga<sup>2</sup>

SnO<sup>2</sup>

β-Ga<sup>2</sup> O3 O3

**4.3. Indium (In)**

were rectifying.

Ga<sup>2</sup> O3

Ga<sup>2</sup> O3

*3.13.7. Silicon (Si)*

MgO has also been used as a substrate for growing Ga<sup>2</sup>

O3

substrate while it is covered by a thin plate with gaps in the desired shape.

Ti/Au interdigitated electrodes have been deposited on both β-Ga<sup>2</sup>

, which makes it easy for carriers to tunnel.

A shadow mask has been used to deposit Au onto the back of a β-Ga<sup>2</sup>

Thermal evaporation was used to deposit an In electrode on the SnO<sup>2</sup>

contact was also sputtered on top of the β-Ga<sup>2</sup>

in high responsivity and quantum efficiency.

air, which affects the structural properties of the Ga<sup>2</sup>

**O3**

O3

on Si is an option to make a p-n junction. Ga<sup>2</sup>

can be either Schottky or ohmic. To make an ohmic contact, the barrier height

nanowire thin films [38]. In both cases, a shadow mask was used and the con-

O3

on MgO, it is initially amorphous and requires annealing to achieve good crystallinity [16]. However, this also decreases the bandgap, making the material less sensitive to UV light. The

film is also highly resistive, although it remains n-type due to oxygen vacancies.

(100) Si using molecular beam epitaxy has good crystallinity and produces smooth surfaces [62].

of the metal should be low. Gold and titanium are the most commonly used metals for contacts. Metal contacts can be deposited using shadow masks, which involve depositing metal onto the

tacts exhibited ohmic behavior. It was suggested that this was due to the abundance of surface

todetector [62]. The deposition method was radio frequency magnetron sputtering. A Ti/Au

O3

heterojunction photodetector [65]. A bilayer of Ti and Au was also deposited on the

thin film using evaporation. The I-V curve of the device showed that the contacts

[64]. Carrier multiplication that occurred in the area under the electrodes resulted

photolithography techniques and liftoff have been used to deposit interdigitated Au electrodes

 can also be grown on silicon. It has been deposited on (100) Si [57]. A minimum temperature of 600°C was required for crystallization to occur. Because silicon, like silicon carbide, can be

. It tends to absorb water from the

O3

O3

[63] and chromium-

/Si heterojunction pho-

part of a β-Ga<sup>2</sup>

O3 /

is grown

grown on

on top [57]. When Ga<sup>2</sup>

O3

O3

. The contacts formed were ohmic. Standard

Graphene electrodes have also been deposited on a β-Ga<sup>2</sup> O3 /SiC photodetector [67]. The graphene was grown by chemical vapor deposition on Cu and then transferred to the substrate. The graphene was transparent and thus allowed most incident light to reach the contact area. It also acted as a channel when the electrons and the holes were separated. On the back of the substrate, a Au/Ti electrode was deposited. A device with a gold electrode on top was also fabricated. The device with the graphene electrode showed a larger photoresponse, higher responsivity, larger Ilight/Idark ratio, and a shorter rise time.

#### **5. Ga2 O3 nanowire applications**

Ga<sup>2</sup> O3 nanowires, nanobelts and nanorods have been extensively studied in recent years. These Ga<sup>2</sup> O3 nanostructures provide a path to a new generation of devices, but most Ga<sup>2</sup> O3 nanostructure-based devices are limited to a single nanowire, and the integration of individual devices on a single chip is still a real challenge. Another crucial issue is controlled doping in nanostructures and the formation of high-quality ohmic contacts between nanostructures and electrodes. Therefore, the construction of a device via a simple and cost-effective method is still a great challenge. In spite of these difficulties, various Ga<sup>2</sup> O3 nanowire applications such as Ga<sup>2</sup> O3 nanowire-based FETs, gas sensors and UV photodetectors are reported.

#### **5.1. UV photodetectors**

The spectrum of radiation between 200 and 280 nm is called solar-blind due to the absorption of solar radiation in that spectrum by the ozone layer. Because of this lack of background noise, it is possible to detect very weak signals in this spectrum. In recent years, the need for solarblind UV photodetectors has increased, due to their potential in several applications such as flame detection and missile warning [68]. When photons strike the surface of a UV photodetector, they generate electron–hole pairs. These electron–hole pairs change the conductivity of the material, which shows that light has been detected. For a device to detect signals in the UV spectrum, a material with a wide bandgap should be used. It is important that the material detect UV light but is not affected by light from other parts of the optical spectrum.

Photomultiplier tubes are used for UV detection in the solar-blind region, but they have several problems. They are bulky, require high bias voltages, and have high leakage current. Siliconbased photodiodes are also used, but due to their narrow bandgap, they require additional filters to block light from unwanted parts of the optical spectrum. Wide-bandgap materials are preferred for use in solar-blind UV photodetectors because they are transparent to deep UV. Of the different materials available, Ga<sup>2</sup> O3 is particularly promising because of its wide bandgap of 4.9 eV, high melting point, and chemical stability.

**5.5. Schottky photodiode**

**5.6. p-n photodiode**

Ga<sup>2</sup> O3

p-type Ga<sup>2</sup>

material.

O3

**5.7. p-i-n photodiode**

responsivity cutoff, and fast response time.

**5.8. Avalanche photodiode**

A Schottky diode consists of a metal layer in contact with a semiconductor layer that exhibits rectifying behavior due to the difference in the work function between the two layers [69]. It has high quantum efficiency, high response speed, and low dark current and exhibits good contrast between UV and visible light. A Schottky diode can operate at zero bias, and doping levels are important for controlling barrier height. Unlike photoconductors, Schottky diodes have a very low persistent photoconductivity [71]. However, a lack of uniformity in metal films leads to large leakage currents, and external quantum efficiency is reduced by transparent Schottky contacts. A back-illuminated Schottky contact does not have this problem, but those devices are more difficult to fabricate. Schottky diodes have a constant responsivity for excitations above the bandgap, independent of power and temperature [72], and the time response is limited by the RC time constant. They have some limitations. For example, the maximum responsivity is limited by reflection in the transparent top contact [75]. The leakage current of the contacts is strongly dependent on the Schottky barrier height, with a higher barrier corresponding to lower

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464 27

A p-n photodiode is a p-n diode made with materials that allow light to penetrate the p-n junction [69]. It has a fast response speed and low dark current, and can operate without an applied bias, so it consumes less power. The photoresponse is also linear with optical power [72]. For

and SiC [59] used on the p side. With heterojunctions, it is possible to tune the bandgap of each

p-i-n photodiodes are similar to p-n photodiodes, except for the addition of an intrinsic layer between the p and n materials. In a p-i-n junction photodiode, absorbed photons generate electron–hole pairs, which are collected by the n and p layers due to the reverse bias. Carriers generated in the junction experience a high electric field and are separated rapidly, giving the detector a fast response [75]. The addition of an intrinsic layer improves absorption and increases the quantum efficiency. To reduce the capacitance, the thickness of the intrinsic layer can be increased. However, a thicker intrinsic layer increases the transit time [74]. The junctions are critical to device performance, because if recombination occurs in the junctions, device performance is degraded. There is a good UV-to-visible rejection ratio, sharp spectral

An avalanche photodiode operates at high speeds and multiplies photocurrent internally, resulting in high sensitivity. Above the breakdown voltage, electron–hole pairs are accelerated by a

 devices, there is good rejection of light from the visible spectrum. The time response is limited by p-doping. A p-n photodiode can be either a homojunction, with both sides made of the same material, or a heterojunction, with different materials. Due to the difficulty of making

, only heterojunction photodiodes have been made, with materials such as Si [62]

leakage current. Defects will also result in a large leakage current.

#### **5.2. Types of photodetector structures**

Many different types of photodetectors can be made. Photoconductors, MSM photodiodes, Schottky diodes, p-n junctions, p-i-n junctions, and avalanche photodiodes have all been fabricated with Ga<sup>2</sup> O3 .

#### **5.3. Photoconductors**

A photoconductor is basically a radiation-sensitive resistor. When a photon with more energy than the bandgap of the material is absorbed, an electron–hole pair is formed, which makes the material more conductive [69]. Photoconductors have a few key advantages, such as high internal gain at room temperature, large photoresponsivity, and lack of a need for amplifying equipment. They are also compatible with planar IC technologies [70], so it is not difficult to integrate them with other devices on a chip. However, a photoconductor has strong persistent photoconductivity, meaning that photocurrent persists even after removal of illumination [71]. In addition, the response speed is very slow, and the responsivity depends on the amount of time it is kept in the dark. It is not possible for a photoconductor to operate at zero bias, meaning that it consumes more power. Photoconductors also exhibit sublinear behavior for incident power [72] and poor contrast between UV and visible light.

#### **5.4. MSM photodiode**

MSM (metal–semiconductor–metal) photodiodes are a commonly used structure for Ga<sup>2</sup> O3 photodetectors. They consist of two back-to-back Schottky diodes with an interdigitated electrode on top of an active light-collecting region [69]. An MSM photodiode has fast operation due to a low capacitance per unit area. The speed of the device is limited by transit time rather than by RC time constant. Using electron beam lithography, electrode width and spacing can be made very small, improving the speed of the device. The device also has a simple structure that is easy to fabricate and integrate. However, there is intrinsic low responsivity due to the interdigitated electrodes covering the active region.

MSM photodiodes have high gain and are easy to integrate with read-out circuitry [73]. The interdigitated electrodes need to be close together to maintain device performance, but this lowers responsivity because it blocks part of the incoming light. It is also possible to illuminate the device from the back, but this makes fabrication much harder. Fabrication of an MSM photodiode requires only a single photolithography step [74] because it only needs a single active layer of dopants. The fabrication process is compatible with that of FETs, but fine feature sizes are required. It is difficult to reliably control the metal–semiconductor interface, and the reflection of light from surface metals is a problem. It is possible to make the dark current in an MSM photodiode very low [72]. The photoresponse is linear with optical power. There is good contrast between the visible and the UV spectrum, and the bandwidth is wide. However, noise is a significant problem [71] and the material has high resistivity [75].

#### **5.5. Schottky photodiode**

filters to block light from unwanted parts of the optical spectrum. Wide-bandgap materials are preferred for use in solar-blind UV photodetectors because they are transparent to deep UV. Of

Many different types of photodetectors can be made. Photoconductors, MSM photodiodes, Schottky diodes, p-n junctions, p-i-n junctions, and avalanche photodiodes have all been fab-

A photoconductor is basically a radiation-sensitive resistor. When a photon with more energy than the bandgap of the material is absorbed, an electron–hole pair is formed, which makes the material more conductive [69]. Photoconductors have a few key advantages, such as high internal gain at room temperature, large photoresponsivity, and lack of a need for amplifying equipment. They are also compatible with planar IC technologies [70], so it is not difficult to integrate them with other devices on a chip. However, a photoconductor has strong persistent photoconductivity, meaning that photocurrent persists even after removal of illumination [71]. In addition, the response speed is very slow, and the responsivity depends on the amount of time it is kept in the dark. It is not possible for a photoconductor to operate at zero bias, meaning that it consumes more power. Photoconductors also exhibit sublinear behavior

MSM (metal–semiconductor–metal) photodiodes are a commonly used structure for Ga<sup>2</sup>

photodetectors. They consist of two back-to-back Schottky diodes with an interdigitated electrode on top of an active light-collecting region [69]. An MSM photodiode has fast operation due to a low capacitance per unit area. The speed of the device is limited by transit time rather than by RC time constant. Using electron beam lithography, electrode width and spacing can be made very small, improving the speed of the device. The device also has a simple structure that is easy to fabricate and integrate. However, there is intrinsic low responsivity due to the

MSM photodiodes have high gain and are easy to integrate with read-out circuitry [73]. The interdigitated electrodes need to be close together to maintain device performance, but this lowers responsivity because it blocks part of the incoming light. It is also possible to illuminate the device from the back, but this makes fabrication much harder. Fabrication of an MSM photodiode requires only a single photolithography step [74] because it only needs a single active layer of dopants. The fabrication process is compatible with that of FETs, but fine feature sizes are required. It is difficult to reliably control the metal–semiconductor interface, and the reflection of light from surface metals is a problem. It is possible to make the dark current in an MSM photodiode very low [72]. The photoresponse is linear with optical power. There is good contrast between the visible and the UV spectrum, and the bandwidth is wide.

However, noise is a significant problem [71] and the material has high resistivity [75].

is particularly promising because of its wide bandgap

O3

O3

for incident power [72] and poor contrast between UV and visible light.

interdigitated electrodes covering the active region.

the different materials available, Ga<sup>2</sup>

26 Novel Nanomaterials - Synthesis and Applications

**5.2. Types of photodetector structures**

O3 .

ricated with Ga<sup>2</sup>

**5.3. Photoconductors**

**5.4. MSM photodiode**

of 4.9 eV, high melting point, and chemical stability.

A Schottky diode consists of a metal layer in contact with a semiconductor layer that exhibits rectifying behavior due to the difference in the work function between the two layers [69]. It has high quantum efficiency, high response speed, and low dark current and exhibits good contrast between UV and visible light. A Schottky diode can operate at zero bias, and doping levels are important for controlling barrier height. Unlike photoconductors, Schottky diodes have a very low persistent photoconductivity [71]. However, a lack of uniformity in metal films leads to large leakage currents, and external quantum efficiency is reduced by transparent Schottky contacts. A back-illuminated Schottky contact does not have this problem, but those devices are more difficult to fabricate. Schottky diodes have a constant responsivity for excitations above the bandgap, independent of power and temperature [72], and the time response is limited by the RC time constant. They have some limitations. For example, the maximum responsivity is limited by reflection in the transparent top contact [75]. The leakage current of the contacts is strongly dependent on the Schottky barrier height, with a higher barrier corresponding to lower leakage current. Defects will also result in a large leakage current.

#### **5.6. p-n photodiode**

A p-n photodiode is a p-n diode made with materials that allow light to penetrate the p-n junction [69]. It has a fast response speed and low dark current, and can operate without an applied bias, so it consumes less power. The photoresponse is also linear with optical power [72]. For Ga<sup>2</sup> O3 devices, there is good rejection of light from the visible spectrum. The time response is limited by p-doping. A p-n photodiode can be either a homojunction, with both sides made of the same material, or a heterojunction, with different materials. Due to the difficulty of making p-type Ga<sup>2</sup> O3 , only heterojunction photodiodes have been made, with materials such as Si [62] and SiC [59] used on the p side. With heterojunctions, it is possible to tune the bandgap of each material.

#### **5.7. p-i-n photodiode**

p-i-n photodiodes are similar to p-n photodiodes, except for the addition of an intrinsic layer between the p and n materials. In a p-i-n junction photodiode, absorbed photons generate electron–hole pairs, which are collected by the n and p layers due to the reverse bias. Carriers generated in the junction experience a high electric field and are separated rapidly, giving the detector a fast response [75]. The addition of an intrinsic layer improves absorption and increases the quantum efficiency. To reduce the capacitance, the thickness of the intrinsic layer can be increased. However, a thicker intrinsic layer increases the transit time [74]. The junctions are critical to device performance, because if recombination occurs in the junctions, device performance is degraded. There is a good UV-to-visible rejection ratio, sharp spectral responsivity cutoff, and fast response time.

#### **5.8. Avalanche photodiode**

An avalanche photodiode operates at high speeds and multiplies photocurrent internally, resulting in high sensitivity. Above the breakdown voltage, electron–hole pairs are accelerated by a large applied field, causing impact ionization. The avalanche effect also amplifies noise. The gain changes with temperature as well as applied bias. Because large voltages are applied, the device consumes large amounts of power. There is no persistent photoconductivity [65] and the quantum efficiency is high. The detectivity and selectivity are also high, and the response and decay times are fast. An avalanche photodiode requires less chip area compared to other type of photodetectors [70], has high gain, and can operate with a high bandwidth. Its fabrication is also compatible with IC fabrication technology. There are uniform junction regions to handle high applied fields [75]. The thickness of the multiplication layer affects the electric field profile and spectral response.

**6. Ga2**

the Ga<sup>2</sup>

**6.1. Ga2**

β-Ga<sup>2</sup> O3

Ga<sup>2</sup> O3

β-Ga<sup>2</sup> O3

Thus, Ga<sup>2</sup>

O3

of challenges associated with Ga<sup>2</sup>

O3

**O3**

**O3**

A single Ga<sup>2</sup>

 **FETs**

O3

cies and extra gallium atoms in the lattice.

based gas sensors [5, 79, 80]. β-Ga<sup>2</sup>

 **nanowire-based temperature/gas sensor**

affected by its surroundings. To fabricate a gas sensor, Ga<sup>2</sup>

**6.2. Photoelectrical generation of hydrogen**

GaN thin films used for the same purpose. Ga<sup>2</sup>

**7. Challenges & future perspectives**

nanowire-based FET has been fabricated to measure the electrical properties

) film

29

O3

O3

nanowires have been grown using

Ga2O3 Nanowire Synthesis and Device Applications http://dx.doi.org/10.5772/intechopen.72464

is particularly attractive because of the

nanowire-

of the as-grown nanowires [77, 78]. The device was fabricated on a silicon dioxide (SiO<sup>2</sup>

on top of an n-type silicon (Si) wafer with parallel pairs of Au contacts. These parallel Au electrodes acted as the source and drain, and an n-type Si layer served as the back gate. The Ga<sup>2</sup>

nanowires were first dispersed in ethanol, and then dried onto an n-type silicon chip. It was found that the conductance increased as the back-gate voltage increased. This suggests that

nanowires had n-type characteristics, likely due to the presence of oxygen vacan-

nanowire-based temperature sensors have unique conductivity behavior up to high

appears to be an insulating material at room temperature;

O3

temperatures. This property distinguishes this material from other semiconductors for nanode-

however, at high temperatures, it behaves as an n-type material and its conductance is greatly

the VLS method [80]. The developed sensor was cheap, easily fabricated and able to detect various chemical constituents at room temperature. As the target chemical substance approaches the nanowires, the gas may be physically adsorbed onto defects on the surface of the nanowire. This changes the dielectric constant of the nanowires, and the device detects the change in capacitance. As the concentration of the gas increases, the capacitance also increases. Because this sensor does not require an external heat source to recover quickly, it can operate at low power.

 nanowires show promise as a photocatalyst for splitting water into hydrogen and oxygen [81]. Efficiency was measured as 0.906%, much higher than the 0.581% efficiency of

O3

nanowires have unique properties that distinguish them from other semiconductors.

could be excellent for nanodevice applications in the future. However, a number

still exist and need to be addressed.

For instance, these nanowires have a wide bandgap and high chemical and thermal stability.

O3

tunability of its optoelectronic properties through doping and alloying, and enhancement of photoelectrochemical efficiency due to the presence of defect bands. However, when the nanowires were grown using a GaN substrate at high temperatures, an interfacial layer was observed between the nanowires and the GaN. This interfacial layer degraded the photoelectrochemical performance of the nanowires. Because the coverage of the substrate with nanowires is not complete, the substrate also contributes to the photoelectrochemical process.

vice applications. Many studies have shown different methods for fabricating β-Ga<sup>2</sup>

O3

#### **5.9. Ga2 O3 nanowire-based field ionization**

The high aspect ratio of Ga<sup>2</sup> O3 nanowires gives them potential to be used as field emission devices [76, 77]. For this application, it is desirable for the material to exhibit emission at a low electric field and remain stable at high current densities [51]. The field enhancement factor is also an important figure of merit for determining whether a material is suitable for field emission. If a nanowire with a sharp tip and a high aspect ratio is used, there is a large electric field at the tip, which reduces the potential barrier for field emission and increases the field emission current. Ga<sup>2</sup> O3 nanowires with ultra-sharp tips 3.5 nm in radius (**Figure 10**) have been fabricated [51], demonstrating that Ga<sup>2</sup> O3 has potential in this area. In fact, the field emission characteristics of turn-on field, threshold electric field, and geometrical field enhancement factor of these nanowires were comparable to those of diamond nanostructures and single-wall carbon nanotubes.

**Figure 10.** Transmission electron microscopy (TEM) image of the sharp tip. An ensemble of Ga or GaxO droplets is visible on the tip.

#### **6. Ga2 O3 FETs**

large applied field, causing impact ionization. The avalanche effect also amplifies noise. The gain changes with temperature as well as applied bias. Because large voltages are applied, the device consumes large amounts of power. There is no persistent photoconductivity [65] and the quantum efficiency is high. The detectivity and selectivity are also high, and the response and decay times are fast. An avalanche photodiode requires less chip area compared to other type of photodetectors [70], has high gain, and can operate with a high bandwidth. Its fabrication is also compatible with IC fabrication technology. There are uniform junction regions to handle high applied fields [75]. The thickness of the multiplication layer affects the electric field profile and spectral response.

[76, 77]. For this application, it is desirable for the material to exhibit emission at a low electric field and remain stable at high current densities [51]. The field enhancement factor is also an important figure of merit for determining whether a material is suitable for field emission. If a nanowire with a sharp tip and a high aspect ratio is used, there is a large electric field at the tip, which reduces the

ultra-sharp tips 3.5 nm in radius (**Figure 10**) have been fabricated [51], demonstrating that Ga<sup>2</sup>

has potential in this area. In fact, the field emission characteristics of turn-on field, threshold electric field, and geometrical field enhancement factor of these nanowires were comparable to those

**Figure 10.** Transmission electron microscopy (TEM) image of the sharp tip. An ensemble of Ga or GaxO droplets is

potential barrier for field emission and increases the field emission current. Ga<sup>2</sup>

nanowires gives them potential to be used as field emission devices

O3

nanowires with

O3

**5.9. Ga2**

**O3**

visible on the tip.

The high aspect ratio of Ga<sup>2</sup>

28 Novel Nanomaterials - Synthesis and Applications

 **nanowire-based field ionization**

O3

of diamond nanostructures and single-wall carbon nanotubes.

A single Ga<sup>2</sup> O3 nanowire-based FET has been fabricated to measure the electrical properties of the as-grown nanowires [77, 78]. The device was fabricated on a silicon dioxide (SiO<sup>2</sup> ) film on top of an n-type silicon (Si) wafer with parallel pairs of Au contacts. These parallel Au electrodes acted as the source and drain, and an n-type Si layer served as the back gate. The Ga<sup>2</sup> O3 nanowires were first dispersed in ethanol, and then dried onto an n-type silicon chip. It was found that the conductance increased as the back-gate voltage increased. This suggests that the Ga<sup>2</sup> O3 nanowires had n-type characteristics, likely due to the presence of oxygen vacancies and extra gallium atoms in the lattice.

#### **6.1. Ga2 O3 nanowire-based temperature/gas sensor**

β-Ga<sup>2</sup> O3 nanowire-based temperature sensors have unique conductivity behavior up to high temperatures. This property distinguishes this material from other semiconductors for nanodevice applications. Many studies have shown different methods for fabricating β-Ga<sup>2</sup> O3 nanowirebased gas sensors [5, 79, 80]. β-Ga<sup>2</sup> O3 appears to be an insulating material at room temperature; however, at high temperatures, it behaves as an n-type material and its conductance is greatly affected by its surroundings. To fabricate a gas sensor, Ga<sup>2</sup> O3 nanowires have been grown using the VLS method [80]. The developed sensor was cheap, easily fabricated and able to detect various chemical constituents at room temperature. As the target chemical substance approaches the nanowires, the gas may be physically adsorbed onto defects on the surface of the nanowire. This changes the dielectric constant of the nanowires, and the device detects the change in capacitance. As the concentration of the gas increases, the capacitance also increases. Because this sensor does not require an external heat source to recover quickly, it can operate at low power.

#### **6.2. Photoelectrical generation of hydrogen**

Ga<sup>2</sup> O3 nanowires show promise as a photocatalyst for splitting water into hydrogen and oxygen [81]. Efficiency was measured as 0.906%, much higher than the 0.581% efficiency of GaN thin films used for the same purpose. Ga<sup>2</sup> O3 is particularly attractive because of the tunability of its optoelectronic properties through doping and alloying, and enhancement of photoelectrochemical efficiency due to the presence of defect bands. However, when the nanowires were grown using a GaN substrate at high temperatures, an interfacial layer was observed between the nanowires and the GaN. This interfacial layer degraded the photoelectrochemical performance of the nanowires. Because the coverage of the substrate with nanowires is not complete, the substrate also contributes to the photoelectrochemical process.

#### **7. Challenges & future perspectives**

β-Ga<sup>2</sup> O3 nanowires have unique properties that distinguish them from other semiconductors. For instance, these nanowires have a wide bandgap and high chemical and thermal stability. Thus, Ga<sup>2</sup> O3 could be excellent for nanodevice applications in the future. However, a number of challenges associated with Ga<sup>2</sup> O3 still exist and need to be addressed.

#### **7.1. Dopants**

One issue is the introduction of specific dopants to enhance the electrical properties of nanowires. Specifically, a dopant that can reliably make Ga<sup>2</sup> O3 p-type needs to be found. More investigation is also required to evaluate the unintentional doping of Ga<sup>2</sup> O3 nanowires during their growth. The influence of these impurities on device performance also requires further study. Some procedural challenges for the growth of Ga<sup>2</sup> O3 nanowires also need to be solved. For instance, more directional Ga<sup>2</sup> O3 nanowire growth is highly desirable and the nanowires should also be more uniform in size and length. Finally, techniques to pattern and etch the Ga<sup>2</sup> O3 nanowires would be highly valuable.

**Author details**

**References**

Badriyah Alhalaili\*, Howard Mao and Saif Islam

\*Address all correspondence to: balhalaili@ucdavis.edu

Electron Device Letters. Jul 2012;**33**:985-987

[4] Weng WY, Hsueh TJ, Chang SJ, Huang GJ, Hsueh HT. A beta-Ga<sup>2</sup>

O3

. Journal of Physics and Chemistry of Solids. 1967;**28**:403-&

. Applied Physics Letters. Oct 4 2010;**97**(14):4-6

O3

[5] Mazeina L, Perkins FK, Bermudez VM, Arnold SP, Prokes SM.Functionalized Ga<sup>2</sup>

Scientific Reports. Nov 27 2015;**5**:2

Aug 17 2010;**26**:13722-13726

Academic Publishers; 2003 [7] Geller S. Structure of beta-Ga<sup>2</sup>

Publishers; 2000

**11**:999-1003

Ga<sup>2</sup> O3

19 2008;**92**

in beta-Ga<sup>2</sup>

O3

Technology B. Jul 2016;**34**:4-7

of Si-implanted solar-blind beta-Ga<sup>2</sup>

Department of Electrical and Computer Engineering, University of California, Davis, USA

[1] Maier D, Alomari M, Grandjean N, Carlin JF, Diforte-Poisson MA, Dua C, et al. InAlN/ GaN HEMTs for operation in the 1000 degrees C regime: A first experiment. IEEE

[2] Tanaka A, Chen RJ, Jungjohann KL, Dayeh SA. Strong geometrical effects in submillimeter selective area growth and light extraction of GaN light emitting diodes on sapphire.

[3] Pearton SJ. GaN and Related Materials II. Australia: Gordon and Breach Science

tector prepared by furnace oxidization of GaN thin film. IEEE Sensors Journal. Apr 2011;

ires as active material in room temperature capacitance-based gas sensors. Langmuir.

[6] Wang ZL. Nanowires and Nanobelts: Materials, Properties, and Devices. Boston: Kluwer

[8] Ahman J, Svensson G, Albertsson J.A reinvestigation of beta-gallium oxide. Acta Crystallographica, Section C: Crystal Structure Communications. Jun 15 1996;**52**:1336-1338 [9] Lorenz MR, Woods JF, Gambino RJ. Some electrical properties of semiconductor beta-

[10] Villora EG, Shimamura K, Yoshikawa Y, Ujiie T, Aoki K. Electrical conductivity and carrier concentration control in beta-Ga(2)O(3) by Si doping. Applied Physics Letters. May

[11] Varley JB, Weber JR, Janotti A, Van de Walle CG. Oxygen vacancies and donor impurities

[12] Ahn S, Ren F, Oh S, Jung Y, Kim J, Mastro MA, et al. Elevated temperature performance

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. Journal of Solid State Chemistry. 1977;**20**:209-210

photodetectors. Journal of Vacuum Science &

solar-blind photode-

O3

nanow-

#### **7.2. Contacts**

Different metals have been tested as contacts in Ga<sup>2</sup> O3 devices including titanium/gold, gold, indium, and graphene. These contacts are a key component in the performance of nanowire devices. Thus, more investigation into the ideal materials for ohmic and Schottky contacts in Ga<sup>2</sup> O3 nanowire devices is required. Integrating particular components of these nanowire devices to build a complete system at low cost is still a challenge. Research should also be performed on ways to improve carrier mobility and reduce contact resistance at the device level, and to lower costs to enable large-scale production.

#### **7.3. Growth techniques**

More research is required to identify the different factors that influence Ga<sup>2</sup> O3 nanowire growth and how these affect their optical, structural, thermal and electrical properties. Lattice mismatch needs to be taken into account when selecting a substrate for nanowire growth. High-quality fabrication at low cost for Ga<sup>2</sup> O3 nanowire growth is needed to make various nanowire devices such as FETs and UV photodetectors more attractive. The growth techniques for Ga<sup>2</sup> O3 nanowires should also be more closely examined to see if they can be improved on. Currently, different techniques such as molecular bean epitaxy (MBE), pulsed laser deposition (PLD), metal–organic chemical vapor deposition (MOCVD) and the hydrothermal method show promise in growing vertical Ga<sup>2</sup> O3 nanowires with high precision and control. Because Ga<sup>2</sup> O3 nanowires are a new material that has only recently been used in devices, integration of these devices with other devices on a chip is still necessary.

#### **8. Conclusions**

Ga<sup>2</sup> O3 nanowires exhibit unique electrical, thermal and optical properties that make them particularly attractive for use in future devices in sensing and optical applications. Using nanowires improves on the performance of thin films. These nanowires have a high surface-tovolume ratio, for high detection sensitivity. From a materials point of view, the wide bandgap of Ga<sup>2</sup> O3 also makes it less sensitive to visible light than other wide-bandgap materials such as GaN and SiC and allows it to handle higher power and temperatures. Thus, Ga<sup>2</sup> O3 nanowires are expected to play a key role in devices of the next generation.

### **Author details**

**7.1. Dopants**

**7.2. Contacts**

**7.3. Growth techniques**

Ga<sup>2</sup> O3

in Ga<sup>2</sup> O3

for Ga<sup>2</sup> O3

Ga<sup>2</sup> O3

Ga<sup>2</sup> O3

of Ga<sup>2</sup> O3

**8. Conclusions**

Specifically, a dopant that can reliably make Ga<sup>2</sup>

challenges for the growth of Ga<sup>2</sup>

30 Novel Nanomaterials - Synthesis and Applications

is also required to evaluate the unintentional doping of Ga<sup>2</sup>

and length. Finally, techniques to pattern and etch the Ga<sup>2</sup>

Different metals have been tested as contacts in Ga<sup>2</sup>

level, and to lower costs to enable large-scale production.

High-quality fabrication at low cost for Ga<sup>2</sup>

show promise in growing vertical Ga<sup>2</sup>

O3

One issue is the introduction of specific dopants to enhance the electrical properties of nanowires.

influence of these impurities on device performance also requires further study. Some procedural

indium, and graphene. These contacts are a key component in the performance of nanowire devices. Thus, more investigation into the ideal materials for ohmic and Schottky contacts

devices to build a complete system at low cost is still a challenge. Research should also be performed on ways to improve carrier mobility and reduce contact resistance at the device

growth and how these affect their optical, structural, thermal and electrical properties. Lattice mismatch needs to be taken into account when selecting a substrate for nanowire growth.

O3

nanowire devices such as FETs and UV photodetectors more attractive. The growth techniques

Currently, different techniques such as molecular bean epitaxy (MBE), pulsed laser deposition (PLD), metal–organic chemical vapor deposition (MOCVD) and the hydrothermal method

nanowires should also be more closely examined to see if they can be improved on.

nanowires are a new material that has only recently been used in devices, integration

 nanowires exhibit unique electrical, thermal and optical properties that make them particularly attractive for use in future devices in sensing and optical applications. Using nanowires improves on the performance of thin films. These nanowires have a high surface-tovolume ratio, for high detection sensitivity. From a materials point of view, the wide bandgap

also makes it less sensitive to visible light than other wide-bandgap materials such as

More research is required to identify the different factors that influence Ga<sup>2</sup>

O3

GaN and SiC and allows it to handle higher power and temperatures. Thus, Ga<sup>2</sup>

of these devices with other devices on a chip is still necessary.

are expected to play a key role in devices of the next generation.

nanowire growth is highly desirable and the nanowires should also be more uniform in size

O3

O3

O3

O3

nanowire devices is required. Integrating particular components of these nanowire

nanowires also need to be solved. For instance, more directional

p-type needs to be found. More investigation

nanowires during their growth. The

nanowires would be highly valuable.

devices including titanium/gold, gold,

nanowire growth is needed to make various

nanowires with high precision and control. Because

O3

O3

nanowires

nanowire

Badriyah Alhalaili\*, Howard Mao and Saif Islam

\*Address all correspondence to: balhalaili@ucdavis.edu

Department of Electrical and Computer Engineering, University of California, Davis, USA

#### **References**


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**Chapter 3**

**Provisional chapter**

phase (needle

B matrix. Ternary

phases,

(Hex) as the main phases after heat treat-

B and TiB2

**Synthesis of TiB2-Ni3B Nanocomposite Powders by**

DOI: 10.5772/intechopen.73593

Combination of good oxidation resistance, thermal stability, hardness and high strength are great interest properties in engineering and, that are possible to obtain with the Ni-Ti-B ternary system. Mechanical alloying (MA) is an alternative method and cheapest for the synthesis of this kind of metal-ceramic materials with respect to the traditional melt and quench process. The transformation sequence of all the mixtures reported the formation of (ɣ Ni) phase

morphology), which was more evident with the increase of titanium content (M2 and M3 mixtures) after 24 h of milling. Thermal activation of the milled powders showed the nucle-

phase (thin flakes morphology) was nucleated onto Ni<sup>3</sup>

alloy by MA took place under a metastable equilibrium, offering the possibility to form glassy alloys for compositions, which are not accessible by melting or quenching techniques.

Metal-ceramic materials (glassy metals) from ternary-eutectic alloys experiment several crystalline transformations where more than two phases or two stages of crystallization have been reported because they are a complex transformation. Different ternary alloys prepared

with a nodular morphology and identified the additional presence of the TiB<sup>2</sup>

B (O boride) and TiB2

**Keywords:** mechanical alloying, metal-ceramic materials, Ni3

**B Nanocomposite Powders by** 

© 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 IntechOpen. 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.

**Mechanical Alloying**

**Mechanical Alloying**

**Synthesis of TiB<sup>2</sup>**

Jorge Morales Hernández,

Jorge Morales Hernández,

Héctor Herrera Hernández,

and Joel Moreno Palmerin

**Abstract**

Verónica N. Martínez Escobedo,

Verónica N. Martínez Escobedo,

http://dx.doi.org/10.5772/intechopen.73593

ation and growth of the Ni3

composite powder, high density

ment, where the TiB2

**1. Introduction**

José M. Juárez García and Joel Moreno Palmerin

Héctor Herrera Hernández, José M. Juárez García

**-Ni<sup>3</sup>**

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter


#### **Synthesis of TiB2-Ni3B Nanocomposite Powders by Mechanical Alloying Synthesis of TiB<sup>2</sup> -Ni<sup>3</sup> B Nanocomposite Powders by Mechanical Alloying**

DOI: 10.5772/intechopen.73593

Jorge Morales Hernández, Verónica N. Martínez Escobedo, Héctor Herrera Hernández, José M. Juárez García and Joel Moreno Palmerin Jorge Morales Hernández, Verónica N. Martínez Escobedo, Héctor Herrera Hernández, José M. Juárez García and Joel Moreno Palmerin

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.73593

#### **Abstract**

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36 Novel Nanomaterials - Synthesis and Applications

rated Ga<sup>2</sup>

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ACS Nano. Apr 2010;**4**:2452-2458

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ires. Science of Advanced Materials. May 2017;**9**:810-814

Physics. 2012;**112**:034307

room-temperature Ga<sup>2</sup>

AIP Advances. Apr 2016;**6**:21

Combination of good oxidation resistance, thermal stability, hardness and high strength are great interest properties in engineering and, that are possible to obtain with the Ni-Ti-B ternary system. Mechanical alloying (MA) is an alternative method and cheapest for the synthesis of this kind of metal-ceramic materials with respect to the traditional melt and quench process. The transformation sequence of all the mixtures reported the formation of (ɣ Ni) phase with a nodular morphology and identified the additional presence of the TiB<sup>2</sup> phase (needle morphology), which was more evident with the increase of titanium content (M2 and M3 mixtures) after 24 h of milling. Thermal activation of the milled powders showed the nucleation and growth of the Ni3 B (O boride) and TiB2 (Hex) as the main phases after heat treatment, where the TiB2 phase (thin flakes morphology) was nucleated onto Ni<sup>3</sup> B matrix. Ternary alloy by MA took place under a metastable equilibrium, offering the possibility to form glassy alloys for compositions, which are not accessible by melting or quenching techniques.

**Keywords:** mechanical alloying, metal-ceramic materials, Ni3 B and TiB2 phases, composite powder, high density

#### **1. Introduction**

Metal-ceramic materials (glassy metals) from ternary-eutectic alloys experiment several crystalline transformations where more than two phases or two stages of crystallization have been reported because they are a complex transformation. Different ternary alloys prepared

© 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 IntechOpen. 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.

by melt spinning with a high content of Ni and B (78 and 18.2 at.% respectively, Ti as complement) have reported the formation of the orthorhombic Ni3 B (O boride), complex cubic boride (NiTi)23B6 (ɣ phase), and f.c.c. nickel (ɣ Ni) nucleates onto it, where the variation in the number of crystals and eutectic colonies by nucleation and growth processes were activated thermally [1].

powders, induction melting, arc melting, and milling of commercial powders, have reported

at temperature lower than 1073 K, making it difficult to obtain the Ƭ phase as a single phase at high temperature. Transformation sequence from the liquid phase reported four invariant equilibria in the Ni-Ti-B system where the Ƭ phase is present in all time, such equilibria are:

Some authors have determined the possible phases theoretically calculated in the Ni-Ti-B system, where the results shows the transformation sequence according the most negative value

The self-propagating high temperature synthesis (SHS) reaction has been used for the forma-

control during the heat transfer and solidification are associated with the concentration of strong internal stresses and the subsequent cracks formation. For these reasons, the addition of a metallic element like Ni as diluent of the Ti-B system was necessary to reduce the heat

Ti-TiB2 obtained by SHS showed that during the combustion reaction, a eutectic liquid

phase was formed by reducing the NiTi with boron. So, the reaction sequence from the

characteristic that NiTi phase was observed at the combustion front during the quenching.

4*Ni* + 2*Ti* → *Ni*<sup>3</sup> *Ti* + *NiTi* (2)

3*NiTi* + 4*B* → *Ni*<sup>3</sup> *Ti* + 2 *TiB*<sup>2</sup> (3)

The glassy metallic alloys obtained by rapid quenching from the liquid phase have special features and different from those of crystalline alloys. Some problems in the production of bulk metallic glasses by casting techniques, are the changes in composition by the occurrence of many eutectic reactions in equilibrium during the rapid quenching, which affect the glass forming ability [5]. Metallic glass formation is not restricted to quenched process because a variety of methods based on the disordering process result in a crystal-to-glass transition such as the mechanical alloying (MA) of elemental powders, where a crystalline phase can be obtained from a disordered

amorphous state to transform after that, in a crystalline stable phase activated thermally.

B composites with low titanium content. (Ni)-TiB2

Through high temperature melting processes has been reported the formation of (Ni)-TiB2

ites were identified at high titanium content. The synthesis of the Ni-Ti-B system by high

B, NiB and Ni3

70 wt.%, showing a minimum particle size of 0.6 μm with 70 wt.% Ni [3].

B3 , Ni3

compound. Due to the high exothermic energy, some difficulties in the growth

, NiB, Ni4

are easiest to occur thermodynamically than the other phases [2].

evolution during SHS reaction and synthesize finer TiB<sup>2</sup>

the Ni-Ti-B system by SHS reported the formation of TiB2

B3 , Ni3

corresponding with Ni-76 at.% Ti transforms into NiTi and Ni3

eutectic liquid at high titanium concentration can be indicated as [4]:

phase at 1073 K (800°C) [2]. Because the B solubility in the Ni-Ti phases is very small

, *L* ↔ *τ* + *TiB*<sup>2</sup> and *L* ↔ *τ* + *Ni*<sup>3</sup> *B* (1)

Synthesis of TiB2-Ni3B Nanocomposite Powders by Mechanical Alloying

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39

Ti, observing that (Ni) and

particles (4–6 μm) [2]. Studies in

and (Ni) as main products, with

Ti at 1215 K (940°C) with the

and (Ni)-Ni3

Ti compos-

Ti with a Ni content in the range from 30 to

B, and Ni3

the TiB2

TiB2

Ni3

TiB2

and (Ni)-Ni3

tion of TiB2

*L* ↔ *τ* + (*Ni*), *L* ↔ *τ* + *TiNi*<sup>3</sup>

of free energy ΔG°, as follows: (Ni), TiB<sup>2</sup>

transient phases such as Ni4

B-Ni-Ti system is of great interest for scientist and technologist due to the combination of good physical properties of TiB2 such as low density, high hardness, good electrical and thermal conductivity, high strength and good oxidation resistance. TiB2 is a reinforced second phase with a great potential in the development of cutting tools that request high speed and the minimum dimensional variation (tolerances) and associate with their thermal stability. In combination with the Ni3 B phase in the nickel matrix, this metal matrix composite was studied to work out the requirements of the heat-resistant alloys. Physical and chemical stability at high temperature of this combination suggests the use of this compound as a corrosion and wear-resistant material under aggressive environments between other industrial applications.

Eutectic phase Ƭ,Ti3 Ni20B6 , at high temperature transform in the compounds of (Ni)-Ni3 B, (Ni)-TiB2 and (Ni)-TiNi3 , which are surrounding to the Ƭ phase during the solidification as shown in **Figure 1** [2]. Investigations since 1958 by different methods, such as sintering of

**Figure 1.** Partial liquidus surface projection of Ni-Ti-B in at.% [2].

powders, induction melting, arc melting, and milling of commercial powders, have reported the TiB2 phase at 1073 K (800°C) [2]. Because the B solubility in the Ni-Ti phases is very small at temperature lower than 1073 K, making it difficult to obtain the Ƭ phase as a single phase at high temperature. Transformation sequence from the liquid phase reported four invariant equilibria in the Ni-Ti-B system where the Ƭ phase is present in all time, such equilibria are:

by melt spinning with a high content of Ni and B (78 and 18.2 at.% respectively, Ti as com-

number of crystals and eutectic colonies by nucleation and growth processes were activated

B-Ni-Ti system is of great interest for scientist and technologist due to the combination of

mal conductivity, high strength and good oxidation resistance. TiB2 is a reinforced second phase with a great potential in the development of cutting tools that request high speed and the minimum dimensional variation (tolerances) and associate with their thermal stability. In

ied to work out the requirements of the heat-resistant alloys. Physical and chemical stability at high temperature of this combination suggests the use of this compound as a corrosion and wear-resistant material under aggressive environments between other industrial applications.

shown in **Figure 1** [2]. Investigations since 1958 by different methods, such as sintering of

(ɣ phase), and f.c.c. nickel (ɣ Ni) nucleates onto it, where the variation in the

such as low density, high hardness, good electrical and ther-

B phase in the nickel matrix, this metal matrix composite was stud-

, at high temperature transform in the compounds of (Ni)-Ni3

, which are surrounding to the Ƭ phase during the solidification as

B (O boride), complex cubic

B,

plement) have reported the formation of the orthorhombic Ni3

boride (NiTi)23B6

good physical properties of TiB2

38 Novel Nanomaterials - Synthesis and Applications

and (Ni)-TiNi3

Ni20B6

**Figure 1.** Partial liquidus surface projection of Ni-Ti-B in at.% [2].

combination with the Ni3

Eutectic phase Ƭ,Ti3

(Ni)-TiB2

thermally [1].

$$L \rightsquigarrow \tau \text{+ (Ni)}, L \rightsquigarrow \tau \text{+ TiNi}\_{\text{g}}\\L \rightsquigarrow \tau \text{+ TiB}\_{\text{z}} \text{and} L \rightsquigarrow \tau \text{+Ni}\_{\text{g}}\\B \tag{1}$$

Some authors have determined the possible phases theoretically calculated in the Ni-Ti-B system, where the results shows the transformation sequence according the most negative value of free energy ΔG°, as follows: (Ni), TiB<sup>2</sup> , NiB, Ni4 B3 , Ni3 B, and Ni3 Ti, observing that (Ni) and TiB2 are easiest to occur thermodynamically than the other phases [2].

The self-propagating high temperature synthesis (SHS) reaction has been used for the formation of TiB2 compound. Due to the high exothermic energy, some difficulties in the growth control during the heat transfer and solidification are associated with the concentration of strong internal stresses and the subsequent cracks formation. For these reasons, the addition of a metallic element like Ni as diluent of the Ti-B system was necessary to reduce the heat evolution during SHS reaction and synthesize finer TiB<sup>2</sup> particles (4–6 μm) [2]. Studies in the Ni-Ti-B system by SHS reported the formation of TiB2 and (Ni) as main products, with transient phases such as Ni4 B3 , Ni3 B, NiB and Ni3 Ti with a Ni content in the range from 30 to 70 wt.%, showing a minimum particle size of 0.6 μm with 70 wt.% Ni [3].

Ni3 Ti-TiB2 obtained by SHS showed that during the combustion reaction, a eutectic liquid corresponding with Ni-76 at.% Ti transforms into NiTi and Ni3 Ti at 1215 K (940°C) with the characteristic that NiTi phase was observed at the combustion front during the quenching. TiB2 phase was formed by reducing the NiTi with boron. So, the reaction sequence from the eutectic liquid at high titanium concentration can be indicated as [4]:

$$\mathbf{4Ni} + \mathbf{2Ti} \to \text{Ni}\_3 \text{Ti} + \text{NiTi} \tag{2}$$

$$\text{3NiTi} + \text{4B} \rightarrow \text{Ni}\_3 \text{Ti} + \text{2 } \text{TiB}\_2 \tag{3}$$

The glassy metallic alloys obtained by rapid quenching from the liquid phase have special features and different from those of crystalline alloys. Some problems in the production of bulk metallic glasses by casting techniques, are the changes in composition by the occurrence of many eutectic reactions in equilibrium during the rapid quenching, which affect the glass forming ability [5]. Metallic glass formation is not restricted to quenched process because a variety of methods based on the disordering process result in a crystal-to-glass transition such as the mechanical alloying (MA) of elemental powders, where a crystalline phase can be obtained from a disordered amorphous state to transform after that, in a crystalline stable phase activated thermally.

Through high temperature melting processes has been reported the formation of (Ni)-TiB2 and (Ni)-Ni3 B composites with low titanium content. (Ni)-TiB2 and (Ni)-Ni3 Ti composites were identified at high titanium content. The synthesis of the Ni-Ti-B system by high temperature processes implies some problems during the solidification (out of equilibria), due to the changes in composition by the occurrence of a lot of eutectic reactions during the rapid quenching, that affect the glass forming ability [5]. The glassy metallic alloys obtained by rapid quenching or melting process demand more energy consumption and elevated cost; for this reason, the mechanical alloying process was explored as an alternative method for the synthesis of this important glassy material.

**3. Experimental**

figure shows the composition of the TiB<sup>2</sup>

mixtures; however, the Ni3

each composition used are indicated in **Table 2**, where the TiB2

**Table 1.** Composition of the mixtures obtained by mechanical alloying.

M1 Ni3

M2 Ƭ, TiB2

M3 TiN3

**Mixture Equilibrium phases**

**Table 2.** Equilibrium phases according to the isothermal section at 1073 K (800°C).

Elemental powders of Ni, Ti and B with high purity (99.7, 99.5 and 99.3%, respectively) were

compounds, three mixtures were synthetized by mechanical alloying according to the compositions shown in **Table 1**. Stainless steel vial with the powder mixture and hardened balls of 4.76 and 12.7 mm in diameter was loaded into a glove box, evacuated during 1200 s (20 min) with a mechanical vacuum pump and after filled with argon gas to prevent the powder oxidation. Ethyl alcohol (2 c.c.) was used as process control agent (PCA) to maintain the balance between fracture and cold welding of the powders. High-energy equipment SPEX mill 80 was used with a ball-to-powder weight ratio of 10:1 until a maximum milling time of 86.4 ks (24 h). Structural evolution of the powders in function of milling time (each 21.6 ks) was characterized by X-ray diffraction (XRD) in a Bruker D8 advanced diffractometer with Cu-Kα (λ = 1.542 Å). Morphology of the powders at the maximum milling time (86.4 ks) was characterized by scanning electron microscope (Jeol 2000). Particle size was validated with a Zeta Sizer equipment, model Malvern MPT-2. Determination of powder's density was evaluated at the end of the milling and after heat treatment at 1173 K (900°C), during 7.2 ks (2 h) under argon atmosphere. Isothermal section at 1073 K (800°C) in **Figure 2** shows the position in the ternary diagram of each synthesized composition, indicating the equilibrium phases to obtain. Yellow circle in the

centrated at this point is projected to the main phases from the Ni-B and Ni-Ti binary alloys to indicate the equilibrium phases according to the eutectic reactions. The equilibrium phases for

**wt. % at. % wt. % at. % wt. % at. %**

B, Ni2

B, TiB2

, Ni3 B

, TiB2 , Ƭ

(24.36 at.%), indicating the formation of a phase between titanium and nickel (TiN3

M1 80 54.77 10 8.29 10 36.93 M2 70 47.25 20 16.34 10 36.47 M3 60 39.82 30 24.36 10 35.80

**Mixture Ni content Ti content B content**

phase under equilibrium condition. Each line con-

Synthesis of TiB2-Ni3B Nanocomposite Powders by Mechanical Alloying

http://dx.doi.org/10.5772/intechopen.73593

B phase was not observed with the maximum content of titanium

phase was reported in the three

).

and Ni3

B

41

used as started materials. To evaluate the effect of Ti content in the formation of TiB<sup>2</sup>

### **2. Mechanical alloying of boride compounds**

Since their beginning in the 1970s, mechanical alloying (MA) was used in the synthesis of base nickel alloys, expanding its applications to various alloy systems. MA is a non-equilibrium process similar to rapid solidification and tempering, where is possible to synthesize amorphous phases, intermetallic compounds and hardened materials by second phases, impossible to obtain with other techniques.

Conventional methods for preparing refractory compounds like boride, carbide and nitride phases are expensive with a long time-consuming. Mechanical alloying (MA) has been a one-step successful route in synthesizing these compounds, like an inexpensive and faster way [6].

Thermodynamic studies have demonstrated that MA may reproduce crystalline, quasicrystalline, nanocrystalline and amorphous alloys [6]. High melting point materials have been obtained with high-energy mills, under a strict atmosphere control and trying to reduce the wear at the mill wall and steel balls. Titanium diboride (TiB2 ) is one example of this kind of compound of high melting point (2790 °C) synthesized by MA, with a wide use by their excellent physical stability at high temperature.

Some researchers have reported the synthesis of TiB2 by the mechanochemical reaction between Al-TiO2 -B2 O3 and Al-B2 O3 -Ti, where the reduction of the oxide compounds catalyzes the formation of TiB2 [7]. This possibility has been proposed for the manufacturing of nanocomposites materials in the solid state, with a ceramic reinforcement (TiB2 ) dispersed in an oxide matrix that could be Al2 O3 , TiO2 or SiO2 [8] .

Phase transformation in the Ti-B system synthesized by mechanical alloying (MA) have reported the formation of TiB2 by gradual diffusion reaction (GDR) mechanism [9]. Ni-B system assumes some solubility of B in the Ni matrix with an increase in their lattice parameter to transform in (Ni) + Ni3 B, being more stable than the NiB phase, which need more dissolution of B in their structure [6].

Composition of the mixture, process time, steel ball-to-powder ratio, balls diameter, atmosphere and the process control agent are some important variables in the MA process that determinates the result; so that, more studies about the synthesis of TiB2 and Ni3 B in the solid state are necessary. In this document, the effect of titanium content in the development of titanium and nickel borides by mechanical alloying was analyzed.

### **3. Experimental**

temperature processes implies some problems during the solidification (out of equilibria), due to the changes in composition by the occurrence of a lot of eutectic reactions during the rapid quenching, that affect the glass forming ability [5]. The glassy metallic alloys obtained by rapid quenching or melting process demand more energy consumption and elevated cost; for this reason, the mechanical alloying process was explored as an alterna-

Since their beginning in the 1970s, mechanical alloying (MA) was used in the synthesis of base nickel alloys, expanding its applications to various alloy systems. MA is a non-equilibrium process similar to rapid solidification and tempering, where is possible to synthesize amorphous phases, intermetallic compounds and hardened materials by second phases, impos-

Conventional methods for preparing refractory compounds like boride, carbide and nitride phases are expensive with a long time-consuming. Mechanical alloying (MA) has been a one-step successful route in synthesizing these compounds, like an inexpensive and faster

Thermodynamic studies have demonstrated that MA may reproduce crystalline, quasicrystalline, nanocrystalline and amorphous alloys [6]. High melting point materials have been obtained with high-energy mills, under a strict atmosphere control and trying to reduce the

compound of high melting point (2790 °C) synthesized by MA, with a wide use by their excel-

 [8] . Phase transformation in the Ti-B system synthesized by mechanical alloying (MA) have

tem assumes some solubility of B in the Ni matrix with an increase in their lattice parameter to

Composition of the mixture, process time, steel ball-to-powder ratio, balls diameter, atmosphere and the process control agent are some important variables in the MA process that

state are necessary. In this document, the effect of titanium content in the development of

) is one example of this kind of

) dispersed in an

B in the solid

by the mechanochemical reaction

and Ni3


by gradual diffusion reaction (GDR) mechanism [9]. Ni-B sys-

B, being more stable than the NiB phase, which need more dissolution

[7]. This possibility has been proposed for the manufacturing of nano-

tive method for the synthesis of this important glassy material.

**2. Mechanical alloying of boride compounds**

wear at the mill wall and steel balls. Titanium diboride (TiB2

O3

O3 , TiO2

composites materials in the solid state, with a ceramic reinforcement (TiB2

determinates the result; so that, more studies about the synthesis of TiB2

titanium and nickel borides by mechanical alloying was analyzed.

or SiO2

Some researchers have reported the synthesis of TiB2

and Al-B2

sible to obtain with other techniques.

40 Novel Nanomaterials - Synthesis and Applications

lent physical stability at high temperature.


oxide matrix that could be Al2

reported the formation of TiB2

way [6].

between Al-TiO2

the formation of TiB2

transform in (Ni) + Ni3

of B in their structure [6].

Elemental powders of Ni, Ti and B with high purity (99.7, 99.5 and 99.3%, respectively) were used as started materials. To evaluate the effect of Ti content in the formation of TiB<sup>2</sup> and Ni3 B compounds, three mixtures were synthetized by mechanical alloying according to the compositions shown in **Table 1**. Stainless steel vial with the powder mixture and hardened balls of 4.76 and 12.7 mm in diameter was loaded into a glove box, evacuated during 1200 s (20 min) with a mechanical vacuum pump and after filled with argon gas to prevent the powder oxidation. Ethyl alcohol (2 c.c.) was used as process control agent (PCA) to maintain the balance between fracture and cold welding of the powders. High-energy equipment SPEX mill 80 was used with a ball-to-powder weight ratio of 10:1 until a maximum milling time of 86.4 ks (24 h). Structural evolution of the powders in function of milling time (each 21.6 ks) was characterized by X-ray diffraction (XRD) in a Bruker D8 advanced diffractometer with Cu-Kα (λ = 1.542 Å). Morphology of the powders at the maximum milling time (86.4 ks) was characterized by scanning electron microscope (Jeol 2000). Particle size was validated with a Zeta Sizer equipment, model Malvern MPT-2. Determination of powder's density was evaluated at the end of the milling and after heat treatment at 1173 K (900°C), during 7.2 ks (2 h) under argon atmosphere.

Isothermal section at 1073 K (800°C) in **Figure 2** shows the position in the ternary diagram of each synthesized composition, indicating the equilibrium phases to obtain. Yellow circle in the figure shows the composition of the TiB<sup>2</sup> phase under equilibrium condition. Each line concentrated at this point is projected to the main phases from the Ni-B and Ni-Ti binary alloys to indicate the equilibrium phases according to the eutectic reactions. The equilibrium phases for each composition used are indicated in **Table 2**, where the TiB2 phase was reported in the three mixtures; however, the Ni3 B phase was not observed with the maximum content of titanium (24.36 at.%), indicating the formation of a phase between titanium and nickel (TiN3 ).


**Table 1.** Composition of the mixtures obtained by mechanical alloying.


**Table 2.** Equilibrium phases according to the isothermal section at 1073 K (800°C).

**Figure 2.** Isothermal section at 1073 K (800°C) of Ni-Ti-B in at.%, indicating the position of the M1, M2 and M3 mixtures respectively [2].

Powder evolution during the milling of the mixtures M2 and M3 is shown in **Figures 4** and **5**,

widening and reduction of the relative intensity from the nickel's peaks were characterized. Mixture M3 with more titanium concentration (24.36 at.%) reported the intermetallic com-

intensity from the nickel's peaks with respect to the M1 and M2 mixtures. This condition was

Thermodynamic data of this system based on experimental information show that the most nega-

The structure of the titanium diboride phase is isomorphous [11] with a large homogeneity ranged from 61 to 70 at.% B, characteristic of the M2 and M3 mixtures at the maximum milling time.

According with the **Figure 2** in the binary composition line of Ni-B, are identified the NiB,

therefore, the largest amount of boron in the mixture reacted with titanium, resulting in a

(unstable) and the (Ni) phase after 6 h of milling. With more milling time (12,

B until the maximum milling time (24 h) with a less widening and more relative

B phases with the increase of nickel content respectively; however, the forma-

B in the milled samples was more evident with the increase of titanium content.

was not detected and were characterized the crystalline phases of (Ni),

phase was identified; with

B phases transforma-

phase [10].

B phase [12, 13].

and (Ni) phases with an asymmetric

and Ni2

Synthesis of TiB2-Ni3B Nanocomposite Powders by Mechanical Alloying

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43

phase was more stable than the nickel borides and

respectively. For the composition M2 after 12 h of milling, the TiB2

**Figure 3.** X-ray diffraction patterns for Ni80Ti10B10 at different milling times (composition in wt.%).

because the deformation energy was taken advantage in the TiB2

tion and not in the formation of the solid solution over saturated of nickel (Ni).

tive enthalpy of formation for the three types of Ti borides corresponds with the TiB2

minimum relation of B/Ni in equilibrium; but enough to the formation of Ni2

more milling time and until the 24 h of milling, the TiB2

pound Ti3

TiB2

Ni2

B4

B4

These means that the formation of the TiB2

18 and 24 h), Ti3

and Ni2

B and Ni3

tion of Ni2

#### **4. Results and discussion**

#### **4.1. Mechanical alloying**

X-ray diffraction patterns taken each 21.6 ks (6 h) from the M1 mixture are shown in **Figure 3**. At 0 h of milling, the diffraction peaks of crystalline Ni, Ti and B were characterized. After 6 h of milling, the Ti and B peaks were not identified, observing the reduction in the relative intensity and widening of the nickel's peaks. Between 6 and 18 h of milling, the intermetallic compound Ti3 B4 (unstable) and the (Ni) phase were identified. At the maximum milling time (24 h), only the diffraction peaks corresponding with a solid solution of nickel (Ni) were characterized, identifying an asymmetric widening and reduction of the relative intensity from the nickel's peaks, due to various factors such as the refinement of grain size and the increase of structure deformation, associated with the tendency to the formation of amorphous compound.

**Figure 3.** X-ray diffraction patterns for Ni80Ti10B10 at different milling times (composition in wt.%).

**4. Results and discussion**

42 Novel Nanomaterials - Synthesis and Applications

B4

X-ray diffraction patterns taken each 21.6 ks (6 h) from the M1 mixture are shown in **Figure 3**. At 0 h of milling, the diffraction peaks of crystalline Ni, Ti and B were characterized. After 6 h of milling, the Ti and B peaks were not identified, observing the reduction in the relative intensity and widening of the nickel's peaks. Between 6 and 18 h of milling, the intermetal-

**Figure 2.** Isothermal section at 1073 K (800°C) of Ni-Ti-B in at.%, indicating the position of the M1, M2 and M3 mixtures

time (24 h), only the diffraction peaks corresponding with a solid solution of nickel (Ni) were characterized, identifying an asymmetric widening and reduction of the relative intensity from the nickel's peaks, due to various factors such as the refinement of grain size and the increase of structure deformation, associated with the tendency to the formation of amor-

(unstable) and the (Ni) phase were identified. At the maximum milling

**4.1. Mechanical alloying**

lic compound Ti3

respectively [2].

phous compound.

Powder evolution during the milling of the mixtures M2 and M3 is shown in **Figures 4** and **5**, respectively. For the composition M2 after 12 h of milling, the TiB2 phase was identified; with more milling time and until the 24 h of milling, the TiB2 and (Ni) phases with an asymmetric widening and reduction of the relative intensity from the nickel's peaks were characterized.

Mixture M3 with more titanium concentration (24.36 at.%) reported the intermetallic compound Ti3 B4 (unstable) and the (Ni) phase after 6 h of milling. With more milling time (12, 18 and 24 h), Ti3 B4 was not detected and were characterized the crystalline phases of (Ni), TiB2 and Ni2 B until the maximum milling time (24 h) with a less widening and more relative intensity from the nickel's peaks with respect to the M1 and M2 mixtures. This condition was because the deformation energy was taken advantage in the TiB2 and Ni2 B phases transformation and not in the formation of the solid solution over saturated of nickel (Ni).

Thermodynamic data of this system based on experimental information show that the most negative enthalpy of formation for the three types of Ti borides corresponds with the TiB2 phase [10]. The structure of the titanium diboride phase is isomorphous [11] with a large homogeneity ranged from 61 to 70 at.% B, characteristic of the M2 and M3 mixtures at the maximum milling time.

According with the **Figure 2** in the binary composition line of Ni-B, are identified the NiB, Ni2 B and Ni3 B phases with the increase of nickel content respectively; however, the formation of Ni2 B in the milled samples was more evident with the increase of titanium content. These means that the formation of the TiB2 phase was more stable than the nickel borides and therefore, the largest amount of boron in the mixture reacted with titanium, resulting in a minimum relation of B/Ni in equilibrium; but enough to the formation of Ni2 B phase [12, 13].

Powder morphology after milling (24 h) for each composition is shown in **Figure 6**. M1 mixture has an acicular particles (**Figure 6a** and **b**) with a minimum particle size of 416.12 nm. Powder density determined by the Archimedes' principle was 15.89 g/c.c. With the increase of titanium content (M2 and M3 mixtures), the characterized morphology showed composite

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phases like inserted needles, indicated with the black arrows in **Figure 6d** and **f**. The minimum particle size was 823.07 and 741.4 nm for the M2 and M3 mixtures, respectively. The increase of particle size in these two last mixtures respecting to the M1 mixture was due to the formation of a composite particle conformed by nodules and needles. The average density

**Figure 6.** Characteristic particle shapes after milling (24 h). (a) and (b) Acicular powder particles in the M1 mixture;

phase in the M2 and M3 mixtures.

(c)–(f) nodular powder particles with inserted needles of TiB2

crystalline

45

particles constituied by nodular powder as matrix, with the nucleation of the TiB2

reported for both mixtures (M2 and M3) was 15.5 g/c.c. similar with the M1 mixture.

**Figure 4.** X-ray diffraction patterns for Ni70Ti20B10 at different milling times (composition in wt.%).

**Figure 5.** X-ray diffraction patterns for Ni70Ti30B10 at different milling times (composition in wt.%).

Powder morphology after milling (24 h) for each composition is shown in **Figure 6**. M1 mixture has an acicular particles (**Figure 6a** and **b**) with a minimum particle size of 416.12 nm. Powder density determined by the Archimedes' principle was 15.89 g/c.c. With the increase of titanium content (M2 and M3 mixtures), the characterized morphology showed composite particles constituied by nodular powder as matrix, with the nucleation of the TiB2 crystalline phases like inserted needles, indicated with the black arrows in **Figure 6d** and **f**. The minimum particle size was 823.07 and 741.4 nm for the M2 and M3 mixtures, respectively. The increase of particle size in these two last mixtures respecting to the M1 mixture was due to the formation of a composite particle conformed by nodules and needles. The average density reported for both mixtures (M2 and M3) was 15.5 g/c.c. similar with the M1 mixture.

**Figure 6.** Characteristic particle shapes after milling (24 h). (a) and (b) Acicular powder particles in the M1 mixture; (c)–(f) nodular powder particles with inserted needles of TiB2 phase in the M2 and M3 mixtures.

**Figure 5.** X-ray diffraction patterns for Ni70Ti30B10 at different milling times (composition in wt.%).

**Figure 4.** X-ray diffraction patterns for Ni70Ti20B10 at different milling times (composition in wt.%).

44 Novel Nanomaterials - Synthesis and Applications

#### **4.2. Heat treatment after milling**

XRD patterns from the M1 mixture after milling and heat treatment (**Figure 7**) show the reflections corresponding with the Ni3 B, TiB2 and TiB phases where the deformation energy stored in the (Ni) phase at the maximum milling time was thermally activated to transform into the Ti and Ni borides. According to the number of reflections, the Ni<sup>3</sup> B phase is more crystalline and stable with respect to the TiB2 and TiB phases. According with the titanium content was expected more nickel boride phases (Ni3 B and Ni2 B) as is indicated in **Table 2**; however the reactions betweeen titanium and boron were more stable to transform in the two crystalline phases (TiB2 and TiB) with a Gibbs free energy more negative with respect the reactions between Ni-B and Ni-Ti [14–16].

For the mixtures M2 and M3 (more titanium content respectivelly) after heat treatment (**Figures 8** and **9**), the phases identified in the milling change to Ni<sup>3</sup> B and TiB2 , moving the equilibrio towards the formation of Ni3 B with less boron content, but with a higher mass relation because the relative intensity of the Ni3 B peaks was taller than the TiB2 peaks, resulted of a more coalescence and growth of Ni3 B with respect the TiB2 phase [17, 18]. Certain consistency was conserved with the equilibium phases indicated in **Table 2**, with the exception that was not identified the eutectic Ƭ phase. The representative reaction in the Ni-Ti-B system synthesized by MA and heat treated is as follows:

$$\text{Ni} + \text{Ti} + \text{B} \rightarrow \text{ (yNi)} + \text{TiB}\_2 \xrightarrow[\text{}] \text{Ni}\_3\text{B} + \text{TiB}\_2 \tag{4}$$

inserted at the matrix. The average particle size like composite was of 639.45 nm with a density powders that increased after heat treatment of 15.63–16.30 g/c.c. This density change was correlated with the more density of the phases Ni3B (orthorhombic), combined with the hexagonal TiB2 phase in the mixtures after heat treatment and with as a result of the nucleation

Synthesis of TiB2-Ni3B Nanocomposite Powders by Mechanical Alloying

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47

**Figure 8.** X-ray diffraction patterns for Ni70Ti20B10 after milling and heat treatment (composition in wt.%).

**Figure 9.** X-ray diffraction patterns for Ni60Ti30B10 after milling and heat treatment (composition in wt.%).

and growth mechanism of both phases.

Powder morphology after heat treatment for each composition is shown in **Figure 10**. Composite particles were characterized in all the mixtures and are formed by particles with a nodular morphology as matrix and particles with thin flakes morphology as second phase

**Figure 7.** X-ray diffraction patterns for Ni80Ti10B10 after milling and heat treatment (composition in wt.%).

**4.2. Heat treatment after milling**

46 Novel Nanomaterials - Synthesis and Applications

tions corresponding with the Ni3

between Ni-B and Ni-Ti [14–16].

line phases (TiB2

line and stable with respect to the TiB2

equilibrio towards the formation of Ni3

tion because the relative intensity of the Ni3

synthesized by MA and heat treated is as follows:

*Ni* + *Ti* + *B* → (*Ni*) + *TiB*<sup>2</sup>

of a more coalescence and growth of Ni3

was expected more nickel boride phases (Ni3

XRD patterns from the M1 mixture after milling and heat treatment (**Figure 7**) show the reflec-

in the (Ni) phase at the maximum milling time was thermally activated to transform into the

the reactions betweeen titanium and boron were more stable to transform in the two crystal-

For the mixtures M2 and M3 (more titanium content respectivelly) after heat treatment

sistency was conserved with the equilibium phases indicated in **Table 2**, with the exception that was not identified the eutectic Ƭ phase. The representative reaction in the Ni-Ti-B system

Powder morphology after heat treatment for each composition is shown in **Figure 10**. Composite particles were characterized in all the mixtures and are formed by particles with a nodular morphology as matrix and particles with thin flakes morphology as second phase

**Figure 7.** X-ray diffraction patterns for Ni80Ti10B10 after milling and heat treatment (composition in wt.%).

B and Ni2

and TiB) with a Gibbs free energy more negative with respect the reactions

B with respect the TiB2

\_\_\_ *Δ*

and TiB phases where the deformation energy stored

and TiB phases. According with the titanium content

B with less boron content, but with a higher mass rela-

B peaks was taller than the TiB2

B phase is more crystal-

, moving the

peaks, resulted

B) as is indicated in **Table 2**; however

B and TiB2

<sup>→</sup> *Ni*<sup>3</sup> *B* + *TiB*<sup>2</sup> (4)

phase [17, 18]. Certain con-

B, TiB2

Ti and Ni borides. According to the number of reflections, the Ni<sup>3</sup>

(**Figures 8** and **9**), the phases identified in the milling change to Ni<sup>3</sup>

**Figure 8.** X-ray diffraction patterns for Ni70Ti20B10 after milling and heat treatment (composition in wt.%).

inserted at the matrix. The average particle size like composite was of 639.45 nm with a density powders that increased after heat treatment of 15.63–16.30 g/c.c. This density change was correlated with the more density of the phases Ni3B (orthorhombic), combined with the hexagonal TiB2 phase in the mixtures after heat treatment and with as a result of the nucleation and growth mechanism of both phases.

**Figure 9.** X-ray diffraction patterns for Ni60Ti30B10 after milling and heat treatment (composition in wt.%).

*Ni* + *Ti* + *B* → (*Ni*) + *TiB*<sup>2</sup>

phase in Ni3

was not enough the activation energy for the formation of the TiNi3

of thin flakes from the TiB<sup>2</sup>

phases.

and Ni3

research group.

**Author details**

Jorge Morales Hernández1

Materiales, Estado de México, México

José M. Juárez García<sup>3</sup>

**References**

and wear resistance.

**Acknowledgements**

TiB2

\_\_\_ *Δ*

B matrix after heat treatment.

B intermetallic compounds, have a great potential in the development of coatings

Nanocomposite materials conformed by particles with nodular morphology and inserted needles were characterized after milling. High resolution microscopy showed the nucleation

Titanium and nickel not showed the formation of a solid solution with the milling, so that

for cutting tools [19]; specialty when is necessary high speed of machining with the minimum dimensional variation (tolerances) and associate with their thermal stability, high hardness

This work was supported by the Sectorial Fund CONACYT-SENER (project number 259334) Energy Sustainability and was developed at the facilities of the Center of Research and Technological Development in Electrochemistry (CIDETEQ) with the INNOVA-COATINGS

\*, Verónica N. Martínez Escobedo1

1 Centro de Investigación y Desarrollo Tecnológico en Electroquímica, Querétaro, México 2 Universidad Autónoma del Estado de México, Área de Electroquímica y Corrosión de

[1] Merk N, Morris DG, Stadelmann P. Acta Metallurgica. 1987;**35**(9):2213-2225

[2] Semenova E, Rokhlin L, Dobatkina T, Kolchugina N. Chapter boron-nickel-titanium. In:

and Joel Moreno Palmerin4

3 Centro Nacional de Metrología, El Márquez Querétaro, México

Refractory Metal Systems. Vol. 11E2. 2010. pp. 153-162

\*Address all correspondence to: jmorales@cideteq.mx

4 Universidad de Guanajuato, Guanajuato, México

indicating a Gibbs free energy (ΔG°) more negative for the transformation of TiB2

<sup>→</sup> *Ni*<sup>3</sup> *B* + *TiB*<sup>2</sup>

Synthesis of TiB2-Ni3B Nanocomposite Powders by Mechanical Alloying

http://dx.doi.org/10.5772/intechopen.73593

intermetallic compound;

, Héctor Herrera Hernández2

and Ni3

B

49

,

**Figure 10.** Nodular morphology particles with inserted thin flakes in all the mixtures after heat treatment at 1173 K (900°C). M1 (a and b), M2 (c and d) and M3 (e and f).

### **5. Conclusions**

Ni-Ti-B system synthesized by mechanical alloying reported for the compositions with high content of nickel, a transformation sequence that consisted in the formation of (ɣ Ni) and TiB2 phases after milling. With the increase in the titanium content, the formation of the TiB2 phase was more evident, as was observed in the M2 and M3 mixtures.

Formation of the (ɣ Ni) phase by the accumulation of deformation energy during the milling was the first step for the subsequent transformation of the orthorhombic Ni<sup>3</sup> B (O boride) and TiB2 (Hex) phases thermally activated. The transformation sequence for the synthesis of Ni3 B and TiB2 phases by MA was:

$$\text{Ni} + \text{Ti} + \text{B} \rightarrow \text{(yNi)} + \text{TiB}\_2 \xrightarrow{\Delta} \text{Ni}\_3\text{B} + \text{TiB}\_2$$

Nanocomposite materials conformed by particles with nodular morphology and inserted needles were characterized after milling. High resolution microscopy showed the nucleation of thin flakes from the TiB<sup>2</sup> phase in Ni3 B matrix after heat treatment.

Titanium and nickel not showed the formation of a solid solution with the milling, so that was not enough the activation energy for the formation of the TiNi3 intermetallic compound; indicating a Gibbs free energy (ΔG°) more negative for the transformation of TiB2 and Ni3 B phases.

TiB2 and Ni3 B intermetallic compounds, have a great potential in the development of coatings for cutting tools [19]; specialty when is necessary high speed of machining with the minimum dimensional variation (tolerances) and associate with their thermal stability, high hardness and wear resistance.

#### **Acknowledgements**

This work was supported by the Sectorial Fund CONACYT-SENER (project number 259334) Energy Sustainability and was developed at the facilities of the Center of Research and Technological Development in Electrochemistry (CIDETEQ) with the INNOVA-COATINGS research group.

#### **Author details**

Jorge Morales Hernández1 \*, Verónica N. Martínez Escobedo1 , Héctor Herrera Hernández2 , José M. Juárez García<sup>3</sup> and Joel Moreno Palmerin4

\*Address all correspondence to: jmorales@cideteq.mx

1 Centro de Investigación y Desarrollo Tecnológico en Electroquímica, Querétaro, México

2 Universidad Autónoma del Estado de México, Área de Electroquímica y Corrosión de Materiales, Estado de México, México

3 Centro Nacional de Metrología, El Márquez Querétaro, México

4 Universidad de Guanajuato, Guanajuato, México

#### **References**

phase

B

B (O boride) and

**5. Conclusions**

TiB2

and TiB2

phases by MA was:

(900°C). M1 (a and b), M2 (c and d) and M3 (e and f).

48 Novel Nanomaterials - Synthesis and Applications

Ni-Ti-B system synthesized by mechanical alloying reported for the compositions with high content of nickel, a transformation sequence that consisted in the formation of (ɣ Ni) and TiB2 phases after milling. With the increase in the titanium content, the formation of the TiB2

**Figure 10.** Nodular morphology particles with inserted thin flakes in all the mixtures after heat treatment at 1173 K

Formation of the (ɣ Ni) phase by the accumulation of deformation energy during the milling

(Hex) phases thermally activated. The transformation sequence for the synthesis of Ni3

was more evident, as was observed in the M2 and M3 mixtures.

was the first step for the subsequent transformation of the orthorhombic Ni<sup>3</sup>


[3] Huang L, Wang HY, Li Q, Yin SQ, Jiang QC. Journal of Alloys and Compounds. 2008;**457**:286-291

**Chapter 4**

**Provisional chapter**

**Properties and Catalytic Effects of Nanoparticles**

**Properties and Catalytic Effects of Nanoparticles**

DOI: 10.5772/intechopen.72158

One of the gas-phased methods of the levitational gas condensation (LGC) process was developed to obtain nanopowders with high purity. The instrument designed by unique concept using magnetically levitated melted droplet of metal is easily operated to synthesize nanopowder with highly defected surface. The complex compounds are also easily prepared using micron powder feeding (MPF) system in the instrument. The metals, ceramics, and carbon-coated metal nanoparticles prepared using the LGC are introduced in this chapter. Various applications such as magnetic and catalytic properties are also introduced. Nanoparticles prepared using LGC showed significantly enhanced catalytic activities during chemical reaction due to the high level of defects on their surface structure. The new heterogeneous catalysts of the solid nanoparticles were introduced in this

**Keywords:** levitational gas condensation (LGC), catalytic activities, metals and metal

Among the various methods for preparing nanopowders, almost all of the processes face important challenges, such as poor control of size distribution, surface contamination, the agglomeration of the particles, and so on [1–5]. Many attempts have been made to develop processes and techniques that can synthesize nanoparticles with specific functional properties [4–6]. Dry methods such as the levitational gas condensation (LGC) process have been developed to obtain high-purity nanopowders while suppressing the agglomeration of the produced particles [7–15]. The catalytic effects of nanopowders are influenced not only by the reduced size of the particles but also by their increased surface area [16, 17]. The surface of metal oxides

oxide nanoparticles, carbon encapsulated nanoparticles

© 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 IntechOpen. 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.

**Synthesized by Levitational Gas Condensation**

**Synthesized by Levitational Gas Condensation**

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.72158

Young Rang Uhm

**Abstract**

chapter.

**1. Introduction**

Young Rang Uhm


**Provisional chapter**

### **Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation Synthesized by Levitational Gas Condensation**

**Properties and Catalytic Effects of Nanoparticles**

DOI: 10.5772/intechopen.72158

#### Young Rang Uhm Additional information is available at the end of the chapter

Young Rang Uhm

[3] Huang L, Wang HY, Li Q, Yin SQ, Jiang QC. Journal of Alloys and Compounds.

[4] Yi HC, Woodger TC, Guigné JY, Moore JJ. Metallurgical and Materials Transaction B.

[7] Mohammad Sharifi E, Karimzadeh F, Enayati MH. Advanced Power Technology.

[8] Mohammad Sharifi E, Karimzadeh F, Enayati MH. Journal of Alloys and Compounds.

[9] Wen-Ming T, Zhi-Xiang Z, Yu-chen W, Jian-Min W, Jun L, Jun-Wu L. Transaction of

[11] Battezati L, Antonione C, Baricco M. Journal of Alloy and Compounds. 1997;**247**:164-171 [12] Neikov OD, Naboychenko SS, Dowson G. Handbook of Non-ferrous Metal Powders,

[13] Xi L, Kaban I, Nowak R, Bruzda G, Sobczak N, Stoica M, Eckert J. Journal of Materials

[14] Selva Kumar M, Chandrasekar P, Chandramohan P, Mohanraj M. Materials Character-

[10] Nakama Y, Ohtani H, Hasebe M. Materials Transaction. 2009;**50**(5):984-993

[15] Karolus M, Panek J. Journal of Alloy and Compounds. 2016;**658**:709-715

[16] Calka A, Radlinski AP. Journal of the Less-Common Metals. 1990;**161**:23-26

[17] Sklad PS, Yust CS. Science of Hard Materials. Plenum Press; 1983. pp. 911-912

[18] Calka A, Radlinski AP. Materials Science and Engineering. 1991;**A134**:1350-1353 [19] Bobzin K. CIRP Journal of Manufacturing Science and Technology. 2017;**18**:1-9

[5] Eckert J. Materials Science and Engineering. 1997;**A226-228**:364-373

[6] Suryanarayana C. Progress in Materials Science. 2001;**46**:1-184

Nonferrous Metals Society of China. 2006;**16**:613-617

Technologies and Applications. Elsevier; 2009

Engineering and Performance. 2016:3204-3210

2008;**457**:286-291

50 Novel Nanomaterials - Synthesis and Applications

1998;**29B**:867-875

2011;**22**:526-531

2010:508-512

ization. 2012;**73**:43-51

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/intechopen.72158

#### **Abstract**

One of the gas-phased methods of the levitational gas condensation (LGC) process was developed to obtain nanopowders with high purity. The instrument designed by unique concept using magnetically levitated melted droplet of metal is easily operated to synthesize nanopowder with highly defected surface. The complex compounds are also easily prepared using micron powder feeding (MPF) system in the instrument. The metals, ceramics, and carbon-coated metal nanoparticles prepared using the LGC are introduced in this chapter. Various applications such as magnetic and catalytic properties are also introduced. Nanoparticles prepared using LGC showed significantly enhanced catalytic activities during chemical reaction due to the high level of defects on their surface structure. The new heterogeneous catalysts of the solid nanoparticles were introduced in this chapter.

**Keywords:** levitational gas condensation (LGC), catalytic activities, metals and metal oxide nanoparticles, carbon encapsulated nanoparticles

#### **1. Introduction**

Among the various methods for preparing nanopowders, almost all of the processes face important challenges, such as poor control of size distribution, surface contamination, the agglomeration of the particles, and so on [1–5]. Many attempts have been made to develop processes and techniques that can synthesize nanoparticles with specific functional properties [4–6]. Dry methods such as the levitational gas condensation (LGC) process have been developed to obtain high-purity nanopowders while suppressing the agglomeration of the produced particles [7–15]. The catalytic effects of nanopowders are influenced not only by the reduced size of the particles but also by their increased surface area [16, 17]. The surface of metal oxides

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 IntechOpen. 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

exhibits nonstoichiometry, resulting from oxygen defect structures [17]. Particles prepared by the LGC process show enhanced catalytic activities due to the high level of defects on their surface structure. In this chapter, the synthesis processes and resulting properties of various nanoparticles prepared by LGC are introduced. The sections are focused on three aspects:


## **2. Levitational gas condensation (LGC)**

The LGC method is a kind of gas-phased method in which an electric current is flowed to two inductor coils, which are each wound opposite directions. The electric current following in different direction in each coil induces a magnetic field and creates a magnetic moment, which opposes gravity on the inside lower part of the coil [11, 18]. To synthesize a nanopowder, a melted metal is continuously evaporated and condensed in the levitated condition, suspended in the magnetic field. This method is shown in **Figure 1(a)**. In this study, we modified the inductor with a downward spiral type of coil, so that the levitation region produced by the magnetic field would be more stable for the melted droplet by generating an equivalent magnetic flux density, as shown in **Figure 1(b)**.

The total LGC system is illustrated in **Figure 2**. The nanoparticles formed at the surface of the liquid droplet are then flowed to a filter by the gas stream using a vacuum pump:

#### **2.1. Levitational gas condensation (LGC) starting materials with a melting temperature of over 900°C**

Metallic atoms were evaporated from an overheated surface and condensed by cold inert gas and then collected from the filter. To stabilize the powder surface and prohibit oxidation, the powders were passivated with thin oxide layers. The LGC apparatus consists of a high-frequency induction generator, levitation and evaporation chamber, and oxygen concentration control unit.

The operating values used for the induction generator were 6, 5, 4.5, and 3 kW for the Ni, Fe, Cu, and Ag, respectively [7–15]. These values depended on the melting temperature as well as the magnetic permeability of the metals. The melting temperature of iron (1535°C) is higher than

**Figure 1.** (a) The mechanism of forming nanoparticles and the contour maps of magnetic flux density depending on shape of inductors for a cylinder type and (b) a spiral type and its strength at in-plane (a~b) and vertical (c~d) direction

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

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53

[11].

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation http://dx.doi.org/10.5772/intechopen.72158 53

exhibits nonstoichiometry, resulting from oxygen defect structures [17]. Particles prepared by the LGC process show enhanced catalytic activities due to the high level of defects on their surface structure. In this chapter, the synthesis processes and resulting properties of various nanoparticles prepared by LGC are introduced. The sections are focused on three aspects:

**1.** Levitational gas condensation (LGC): The unique instrument used for the synthesis of the

**2.** Magnetic metal and carbon-encapsulated metal nanoparticles: The produced magnetic metal (Ni and Fe) and carbon-encapsulated metal (Ni@C and Fe@C) nanoparticles showed a noncollinear magnetic structure between the core and surface layer of the particles. The morphologies and the dispersion stability kinetics in the solvents are introduced. Also, the carbon-encapsulated metal nanoparticles were successfully applied as a catalyst for the mul-

**3.** Nonmagnetic metal and metal oxides: Cu oxides, Bi, and NiO nanopowders prepared by LGC are introduced. Cu oxide and NiO alloy nanopowders are widely applied as heterogeneous catalysts in oxidizing processes used in organic synthesis. Nanopowders of bismuth (Bi) with its low melting temperature were applied as a sensor electrode for

The LGC method is a kind of gas-phased method in which an electric current is flowed to two inductor coils, which are each wound opposite directions. The electric current following in different direction in each coil induces a magnetic field and creates a magnetic moment, which opposes gravity on the inside lower part of the coil [11, 18]. To synthesize a nanopowder, a melted metal is continuously evaporated and condensed in the levitated condition, suspended in the magnetic field. This method is shown in **Figure 1(a)**. In this study, we modified the inductor with a downward spiral type of coil, so that the levitation region produced by the magnetic field would be more stable for the melted droplet by generating an equivalent

The total LGC system is illustrated in **Figure 2**. The nanoparticles formed at the surface of the

Metallic atoms were evaporated from an overheated surface and condensed by cold inert gas and then collected from the filter. To stabilize the powder surface and prohibit oxidation, the powders were passivated with thin oxide layers. The LGC apparatus consists of a high-frequency induction generator, levitation and evaporation chamber, and oxygen concentration control unit.

liquid droplet are then flowed to a filter by the gas stream using a vacuum pump:

**2.1. Levitational gas condensation (LGC) starting materials with a melting** 

nanoparticles is introduced in this section [18–20].

52 Novel Nanomaterials - Synthesis and Applications

ticomponent Biginelli reaction [7, 9, 12, 18–23].

detecting heavy metals in water [9, 14, 17, 24, 25].

**2. Levitational gas condensation (LGC)**

magnetic flux density, as shown in **Figure 1(b)**.

**temperature of over 900°C**

**Figure 1.** (a) The mechanism of forming nanoparticles and the contour maps of magnetic flux density depending on shape of inductors for a cylinder type and (b) a spiral type and its strength at in-plane (a~b) and vertical (c~d) direction [11].

The operating values used for the induction generator were 6, 5, 4.5, and 3 kW for the Ni, Fe, Cu, and Ag, respectively [7–15]. These values depended on the melting temperature as well as the magnetic permeability of the metals. The melting temperature of iron (1535°C) is higher than

**Figure 2.** The concept of system for LGC.

that of Ni (1450°C). However, the high magnetic permeability of Ni affects the magnetic force and levitation of the melted droplets. Accordingly, the input power for Ni needs to be increased up to 6 kW, the maximum power for the inductor. Preparing pure Ti nanopowders using the LGC is impossible. To do so, the temperature of the inductor must be increased up to 2000°C; however, the starting materials have to be very thin and strongly passivated by a layer of titanium oxide. The seed materials need to be fully melted before levitation, because only liquid seeds can be suspended in the inductor, due to their lower density. However, it takes too long time to melt the Ti seed to produce liquid droplet for levitation.

methane (CH4

the mixed Ar and CH4

**Figure 3.** TEM images for metals and ceramic nanopowders.

of carbon [26–28].

) concentration. The starting materials were Ni and Fe wires with a diameter of

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

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55

was 10% of the mixture gas. The

4 mm. The Ni and Fe wires were fed into a melted droplet using the wire feeding system at a feeding rate of 2 mm/min. An ingot of 85 mg, which was used as the seed material for the levitation and the evaporation reactions, was melted by using electric induction. The pressure of

inductor was heated up to a temperature of 2000°C, and the metallic atoms were evaporated from the overheated surface of the liquid droplet and condensed by cold inert gas and then collected into the filter. At the same time, the molecular CH introduced into the chamber was converted to atomic C and H with high activity under high temperature. The highly active C atoms react with the Ni and Fe atoms, and the H atoms are converted to H molecules. The newly created H gas is vented out of the reaction chamber by continuous vacuum operation. The results indicated that all of the as-made materials were composed of nanocapsules with uniform particle size at and below 10 nm. The nanocapsules consisted of outer multi-shells

gas in the chamber was 100 Torr. CH4

**2.2. Starting materials with a melting temperature below 900°C: Zn, Sn, Bi, and ZnO**

For most metals, the optimal design of the material heating method allows a metal drop to be heated and kept in a noncontact condition in the evaporation zone by high-frequency magnetic field. However, this method does not ensure the optimal heating of light-volatile metals such as Zn, Sn, and Bi. Therefore, another material evaporation method using a refractory crucible was applied, for heating and evaporation. This method is generally suitable for

We also supplied the starting materials for the melted liquid droplet during synthesis. We utilized a metal wire feed system, which is very convenient for fabricating nanopowder continuously. The amount of material fed over time can be controlled. The average size of the nanopowder is increased with increasing feed speed because of the increased amount of material introduced to the liquid droplet. The optimal feeding speed is between 10 and 30 mm/s during fabrication. If the feeding speed is very slow, below 10 mm/s, the size of the liquid droplets decreased, and they disappear. In contrast, if the feeing speed is increased to over 30 mm/s, the prepared powders have a wire shape because the particles are connected with each other. By adding oxygen to the inert gas, metal oxide particles or metal particles oxidized on the surface may be obtained. **Figure 3** shows metal and metal oxide particles prepared by LGC. The gas condensation (GC) method was used to obtain nanocrystalline powders of pure metal and nano-oxides with different compositions. Oxide nanopowders, such as Fe3 O4 , γ-Fe<sup>2</sup> O3 , and Cu2 O, were produced by the LGC of metal wires at elevated pressure in an (Ar + O2 ) atmosphere [8, 16].

High-purity Ni@C and Fe@C nanopowders were synthesized using the LGC method. The LGC apparatus consisted of a high-frequency induction generator, operating at 6 kW for Ni and 5 kW for Fe, a reaction chamber for the levitated liquid seed, and a unit to control the Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation http://dx.doi.org/10.5772/intechopen.72158 55

**Figure 3.** TEM images for metals and ceramic nanopowders.

that of Ni (1450°C). However, the high magnetic permeability of Ni affects the magnetic force and levitation of the melted droplets. Accordingly, the input power for Ni needs to be increased up to 6 kW, the maximum power for the inductor. Preparing pure Ti nanopowders using the LGC is impossible. To do so, the temperature of the inductor must be increased up to 2000°C; however, the starting materials have to be very thin and strongly passivated by a layer of titanium oxide. The seed materials need to be fully melted before levitation, because only liquid seeds can be suspended in the inductor, due to their lower density. However, it takes too long

We also supplied the starting materials for the melted liquid droplet during synthesis. We utilized a metal wire feed system, which is very convenient for fabricating nanopowder continuously. The amount of material fed over time can be controlled. The average size of the nanopowder is increased with increasing feed speed because of the increased amount of material introduced to the liquid droplet. The optimal feeding speed is between 10 and 30 mm/s during fabrication. If the feeding speed is very slow, below 10 mm/s, the size of the liquid droplets decreased, and they disappear. In contrast, if the feeing speed is increased to over 30 mm/s, the prepared powders have a wire shape because the particles are connected with each other. By adding oxygen to the inert gas, metal oxide particles or metal particles oxidized on the surface may be obtained. **Figure 3** shows metal and metal oxide particles prepared by LGC. The gas condensation (GC) method was used to obtain nanocrystalline powders of pure metal and nano-oxides with different compositions. Oxide nanopowders,

High-purity Ni@C and Fe@C nanopowders were synthesized using the LGC method. The LGC apparatus consisted of a high-frequency induction generator, operating at 6 kW for Ni and 5 kW for Fe, a reaction chamber for the levitated liquid seed, and a unit to control the

O, were produced by the LGC of metal wires at elevated pres-

time to melt the Ti seed to produce liquid droplet for levitation.

, and Cu2

) atmosphere [8, 16].

such as Fe3

O4 , γ-Fe<sup>2</sup> O3

**Figure 2.** The concept of system for LGC.

54 Novel Nanomaterials - Synthesis and Applications

sure in an (Ar + O2

methane (CH4 ) concentration. The starting materials were Ni and Fe wires with a diameter of 4 mm. The Ni and Fe wires were fed into a melted droplet using the wire feeding system at a feeding rate of 2 mm/min. An ingot of 85 mg, which was used as the seed material for the levitation and the evaporation reactions, was melted by using electric induction. The pressure of the mixed Ar and CH4 gas in the chamber was 100 Torr. CH4 was 10% of the mixture gas. The inductor was heated up to a temperature of 2000°C, and the metallic atoms were evaporated from the overheated surface of the liquid droplet and condensed by cold inert gas and then collected into the filter. At the same time, the molecular CH introduced into the chamber was converted to atomic C and H with high activity under high temperature. The highly active C atoms react with the Ni and Fe atoms, and the H atoms are converted to H molecules. The newly created H gas is vented out of the reaction chamber by continuous vacuum operation. The results indicated that all of the as-made materials were composed of nanocapsules with uniform particle size at and below 10 nm. The nanocapsules consisted of outer multi-shells of carbon [26–28].

#### **2.2. Starting materials with a melting temperature below 900°C: Zn, Sn, Bi, and ZnO**

For most metals, the optimal design of the material heating method allows a metal drop to be heated and kept in a noncontact condition in the evaporation zone by high-frequency magnetic field. However, this method does not ensure the optimal heating of light-volatile metals such as Zn, Sn, and Bi. Therefore, another material evaporation method using a refractory crucible was applied, for heating and evaporation. This method is generally suitable for evaporating materials with high vapor pressures at moderate temperature. **Figure 4** shows a simple diagram of the device for obtaining light-volatile metal nanopowders. The apparatus consists of a high-frequency induction generator operating at 2.5 kW, a levitation and evaporation chamber, and an oxygen concentration control unit [11]. The wire feeding velocity (VZn) and mixed Ar and O2 gas pressure in the chamber were 50 mm/min and 100 Torr, respectively. The mechanism of ZnO formation using the LGC method was analyzed. First, a liquid droplet, which is levitated against gravity by the magnetic force due to the coupled induction coils, is heated up to the temperature of 1560°C at 2.5 kW. Then Zn clusters are evaporated from the overheated surface of the liquid droplet and condensed by cold inert gas and collected into the filter. At the same time, molecular O<sup>2</sup> introduced into the chamber is converted to atomic O with high activity under high temperature. The highly active O atoms can diffuse into the Zn clusters and react with the Zn atoms. A large amount of the Zn phase was observed at and below an oxygen flow rate of VO2 = 0.05 ℓ/min, whereas mixtures of ZnO and small amounts of the Zn phase were observed under O2 flow rates in the range from VO2 = 0.11ℓ/min to VO2 = 0.21 ℓ/min. However, at and above 0.21 ℓ/min of O<sup>2</sup> flow rate, levitation was impossible. Some metals, such as Bi and Sn, have insufficient tensile strength to prepare wire. In these cases, the micron-sized powders were used as the parent materials. The powder starting materials were supplied by a powder feeding (PF) system. A detailed explanation of the PF is provided in Section 2.3.

Bi powders were prepared using metal bismuth powder as the starting material [28]. The powder was supplied by the feeding system into a graphite crucible at a rate of 20 mg/min. The crucible was heated by induction currents up to T = 700–900°C. Bi particles entering the crucible were evaporated within 1–2 sec and carried by argon flow from the hot zone. The argon flow rate was varied in the range of 80–170 ℓ/h, at pressures in the range of 70–300 torr. The dependence of the mean sizes of the bismuth particles on gas pressure at a flow rate of 80 ℓ/h was 25, 70, and 120 nm for 70, 150, and 300 torr, respectively. Thus, the optimal conditions for obtaining Bi powder were realized at an argon pressure of 70 torr and a rate of 80 ℓ/h. A simple diagram of the process for forming nanoparticles from light-volatile seed in a crucible to produce volatile nanoparticles is represented in **Figure 5**.

**2.3. Powder starting materials: NiFe2O4 and Ti-Ni**

The wire feeding (WF) system was used for synthesizing metal, ceramic, and carbon-encapsulated materials. This system easily supplies seed parent materials continuously. However, it was impossible to synthesize several complicated metal-doped materials such as ferrites, perovskite, garnet, metal-doped ZnO, Ti-Ni, and Al-Ni-Co using the wire feeder in the LGC system, because the parent materials could not be prepared as wire. A newly modified micron powder feeding (MPF) system overcomes this problem of the LGC system [15, 20]. The MPF system can be used for synthesizing brittle metals, alloys, and complex doped materials. Commercial elemental powders of Ti (99.9 at.%, ~500 μm), Ni (99.9 at.%, ~500 μm), and Fe (99.9 at.%, ~ 500 μm) were used as the starting powders for the synthesis of Ti-Ni alloy and Ni-ferrite nanopowder, using the LGC. The Ti and Ni powders were mixed by pestle and mortar to achieve the desired equi-atomic composition and were then incorporated into the micron powder feeding system, which consisted of a rotating part to supply the Ti and Ni micron powders to the melted droplet and a vibrating part for mixing the powder. The Ti and Ni micron powders were fed into the powder feeding system at a feeding rate of 38 mg/min. An 83 mg Ti-Ni alloy ingot, which was used as the seed material for the levitation and

**Figure 5.** (a) Elementary diagram of the process for forming nanoparticles from light-volatile seed in a crucible and

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(b) TEM images for the Sn, Bi, and ZnO nanoparticles prepared by LGC using light-volatile seed.

**Figure 4.** High-resolution TEM images of carbon encapsulated (a) Ni and (b) Fe.

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation http://dx.doi.org/10.5772/intechopen.72158 57

**Figure 5.** (a) Elementary diagram of the process for forming nanoparticles from light-volatile seed in a crucible and (b) TEM images for the Sn, Bi, and ZnO nanoparticles prepared by LGC using light-volatile seed.

#### **2.3. Powder starting materials: NiFe2O4 and Ti-Ni**

evaporating materials with high vapor pressures at moderate temperature. **Figure 4** shows a simple diagram of the device for obtaining light-volatile metal nanopowders. The apparatus consists of a high-frequency induction generator operating at 2.5 kW, a levitation and evaporation chamber, and an oxygen concentration control unit [11]. The wire feeding veloc-

respectively. The mechanism of ZnO formation using the LGC method was analyzed. First, a liquid droplet, which is levitated against gravity by the magnetic force due to the coupled induction coils, is heated up to the temperature of 1560°C at 2.5 kW. Then Zn clusters are evaporated from the overheated surface of the liquid droplet and condensed by cold inert

is converted to atomic O with high activity under high temperature. The highly active O atoms can diffuse into the Zn clusters and react with the Zn atoms. A large amount of the Zn phase was observed at and below an oxygen flow rate of VO2 = 0.05 ℓ/min, whereas mixtures

levitation was impossible. Some metals, such as Bi and Sn, have insufficient tensile strength to prepare wire. In these cases, the micron-sized powders were used as the parent materials. The powder starting materials were supplied by a powder feeding (PF) system. A detailed

Bi powders were prepared using metal bismuth powder as the starting material [28]. The powder was supplied by the feeding system into a graphite crucible at a rate of 20 mg/min. The crucible was heated by induction currents up to T = 700–900°C. Bi particles entering the crucible were evaporated within 1–2 sec and carried by argon flow from the hot zone. The argon flow rate was varied in the range of 80–170 ℓ/h, at pressures in the range of 70–300 torr. The dependence of the mean sizes of the bismuth particles on gas pressure at a flow rate of 80 ℓ/h was 25, 70, and 120 nm for 70, 150, and 300 torr, respectively. Thus, the optimal conditions for obtaining Bi powder were realized at an argon pressure of 70 torr and a rate of 80 ℓ/h. A simple diagram of the process for forming nanoparticles from light-volatile seed in a crucible to pro-

gas and collected into the filter. At the same time, molecular O<sup>2</sup>

explanation of the PF is provided in Section 2.3.

duce volatile nanoparticles is represented in **Figure 5**.

**Figure 4.** High-resolution TEM images of carbon encapsulated (a) Ni and (b) Fe.

of ZnO and small amounts of the Zn phase were observed under O2

from VO2 = 0.11ℓ/min to VO2 = 0.21 ℓ/min. However, at and above 0.21 ℓ/min of O<sup>2</sup>

gas pressure in the chamber were 50 mm/min and 100 Torr,

introduced into the chamber

flow rates in the range

flow rate,

ity (VZn) and mixed Ar and O2

56 Novel Nanomaterials - Synthesis and Applications

The wire feeding (WF) system was used for synthesizing metal, ceramic, and carbon-encapsulated materials. This system easily supplies seed parent materials continuously. However, it was impossible to synthesize several complicated metal-doped materials such as ferrites, perovskite, garnet, metal-doped ZnO, Ti-Ni, and Al-Ni-Co using the wire feeder in the LGC system, because the parent materials could not be prepared as wire. A newly modified micron powder feeding (MPF) system overcomes this problem of the LGC system [15, 20]. The MPF system can be used for synthesizing brittle metals, alloys, and complex doped materials. Commercial elemental powders of Ti (99.9 at.%, ~500 μm), Ni (99.9 at.%, ~500 μm), and Fe (99.9 at.%, ~ 500 μm) were used as the starting powders for the synthesis of Ti-Ni alloy and Ni-ferrite nanopowder, using the LGC. The Ti and Ni powders were mixed by pestle and mortar to achieve the desired equi-atomic composition and were then incorporated into the micron powder feeding system, which consisted of a rotating part to supply the Ti and Ni micron powders to the melted droplet and a vibrating part for mixing the powder. The Ti and Ni micron powders were fed into the powder feeding system at a feeding rate of 38 mg/min. An 83 mg Ti-Ni alloy ingot, which was used as the seed material for the levitation and

nanoparticles and their dispersion stability in the solvent. It is also worth noting that after encapsulation with a carbon coating layer, these materials are not prone to agglomeration because the coating reduces their magnetic interaction. In addition, the surface diffusion processes can preserve the chemical and structural properties of the nanopowder for a long time in many chemically aggressive conditions. A graphitic carbon shell in particular is regarded as an ideal coating since it is light and shows high stability in both chemical and physical

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

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The contents of the Ni and Ni@C nanoparticles synthesized by LGC using the micron powder feeding system were confirmed by XRD pattern. The XRD results for Ni and Ni@C showed the lattice parameters and the positions of the main peaks of the Ni powders. A small amount of NiO phase and amorphous graphitic layers was found in the XRD patterns and in the TEM images as mentioned in Section 1 [18]. The diffraction peaks at 44.4°, 51.8°, and 76.3° are due to the (1 1 1), (2 0 0), and (2 2 0) planes of fcc-Ni, respectively. The Ni powders synthesized by the LGC method showed low saturation magnetization. These results were attributed to the spin-canting effect and oxide phase on the surface [32]. The magnetic properties would be weak due to the antiferromagnetic NiO phase on the powder surface. The saturation magnetization was Ms. = 42 emu/g, as shown in **Figure 7(a)**. The slightly shifted hysteresis loop for the Ni sample can be explained by exchange bias between the ferromagnetic core of Ni and the antiferromagnetic surface of the NiO. The initial magnetization curve is not explained by the size effect. In previous studies, the virgin magnetization curve slightly spills over the limited hysteresis loop at 655 Oe. We assume that this effect is enhanced when the size of the particles is reduced, as suggested in a previous study. With decreasing particle size, the defects and the different magnetic structure on the surface of the particles are increased. The nature of this irreversibility in high magnetic fields follows a physical model and can be explained by a spin-glass or spin-canting behavior. The hysteresis loop of the as-made M@C materials in magnetic fields up to 2 T reveals their intrinsic magnetic behavior, indicated by the magnetization (M), the remanent magnetization (Mr), and the coercive force (Hc) of the M@C samples. The saturation magnetization demonstrates that

**Figure 7.** M-H loops for (a) Ni and (b) carbon-encapsulated Ni measured at 20°C.

environments [29–31].

**Figure 6.** (a) Micron powder feeding (MPF) system and (b) wire feeding (WF) system in the LGC instrument. TEM images for (c) Ti-Ni alloys and (d) NiFe2 O4 .

evaporation reactions, was melted by an electric induction heating with an applied power of 6 kW at an argon gas pressure of 100 torr. The evaporated powders were filtered and finally passivated by partial oxidation. The starting materials were the mixed micron powders of Ni and Fe, which has a size ranging from 100 to 500 μm. The amount of micron powder fed into the liquid seed droplet was controlled at 80 mg/min. The mixed Ar and O2 gas pressure in the chamber was 100 torr [15, 20] (**Figure 6**).

#### **3. Magnetic metal and carbon-encapsulated metal nanoparticles**

#### **3.1. Magnetic properties of Ni, Fe, Ni@C, and Fe@C**

Magnetic nanoparticles have attracted much attention because of their use in nano-fluids for biomedical application, thermally conductive fluids, various catalysts, etc. However, metallic nano-fluids end to be inherently vulnerable to oxidation, dissolution, and agglomeration during synthesis. In particular, agglomeration of the particles in a solvent is a serious problem when preparing nano-fluids. To overcome these problems, encapsulating the particles in a protective shell has been recommended to improve the chemical stability of the metal nanoparticles and their dispersion stability in the solvent. It is also worth noting that after encapsulation with a carbon coating layer, these materials are not prone to agglomeration because the coating reduces their magnetic interaction. In addition, the surface diffusion processes can preserve the chemical and structural properties of the nanopowder for a long time in many chemically aggressive conditions. A graphitic carbon shell in particular is regarded as an ideal coating since it is light and shows high stability in both chemical and physical environments [29–31].

The contents of the Ni and Ni@C nanoparticles synthesized by LGC using the micron powder feeding system were confirmed by XRD pattern. The XRD results for Ni and Ni@C showed the lattice parameters and the positions of the main peaks of the Ni powders. A small amount of NiO phase and amorphous graphitic layers was found in the XRD patterns and in the TEM images as mentioned in Section 1 [18]. The diffraction peaks at 44.4°, 51.8°, and 76.3° are due to the (1 1 1), (2 0 0), and (2 2 0) planes of fcc-Ni, respectively. The Ni powders synthesized by the LGC method showed low saturation magnetization. These results were attributed to the spin-canting effect and oxide phase on the surface [32]. The magnetic properties would be weak due to the antiferromagnetic NiO phase on the powder surface. The saturation magnetization was Ms. = 42 emu/g, as shown in **Figure 7(a)**. The slightly shifted hysteresis loop for the Ni sample can be explained by exchange bias between the ferromagnetic core of Ni and the antiferromagnetic surface of the NiO. The initial magnetization curve is not explained by the size effect. In previous studies, the virgin magnetization curve slightly spills over the limited hysteresis loop at 655 Oe. We assume that this effect is enhanced when the size of the particles is reduced, as suggested in a previous study. With decreasing particle size, the defects and the different magnetic structure on the surface of the particles are increased. The nature of this irreversibility in high magnetic fields follows a physical model and can be explained by a spin-glass or spin-canting behavior. The hysteresis loop of the as-made M@C materials in magnetic fields up to 2 T reveals their intrinsic magnetic behavior, indicated by the magnetization (M), the remanent magnetization (Mr), and the coercive force (Hc) of the M@C samples. The saturation magnetization demonstrates that

**Figure 7.** M-H loops for (a) Ni and (b) carbon-encapsulated Ni measured at 20°C.

evaporation reactions, was melted by an electric induction heating with an applied power of 6 kW at an argon gas pressure of 100 torr. The evaporated powders were filtered and finally passivated by partial oxidation. The starting materials were the mixed micron powders of Ni and Fe, which has a size ranging from 100 to 500 μm. The amount of micron powder fed into

**Figure 6.** (a) Micron powder feeding (MPF) system and (b) wire feeding (WF) system in the LGC instrument. TEM

Magnetic nanoparticles have attracted much attention because of their use in nano-fluids for biomedical application, thermally conductive fluids, various catalysts, etc. However, metallic nano-fluids end to be inherently vulnerable to oxidation, dissolution, and agglomeration during synthesis. In particular, agglomeration of the particles in a solvent is a serious problem when preparing nano-fluids. To overcome these problems, encapsulating the particles in a protective shell has been recommended to improve the chemical stability of the metal

gas pressure in the

the liquid seed droplet was controlled at 80 mg/min. The mixed Ar and O2

O4 .

**3. Magnetic metal and carbon-encapsulated metal nanoparticles**

chamber was 100 torr [15, 20] (**Figure 6**).

images for (c) Ti-Ni alloys and (d) NiFe2

58 Novel Nanomaterials - Synthesis and Applications

**3.1. Magnetic properties of Ni, Fe, Ni@C, and Fe@C**

the carbon-coated Ni nanocrystallites exhibited a superparamagnetic behavior at room temperature, which is related to the demagnetization effect arising from the additional energy of the magnetic fields outside the graphitic carbon encapsulation as shown in **Figure 7(b)**. The coercive force (Hc) and magnetization (M) were 76.6 Oe and 19.6 emu/g, respectively. The ratio of remanence to the saturation magnetization (Mr/M) was 0.04. The low magnetization compared with the Ni nanoparticles without the carbon shell is due to the coexistence of nonmagnetic carbon and the large percentage of surface spin due to the disordered magnetization orientation of the nanoparticles. The magnetic properties are influenced by both the particle size and the surface properties of the particle [33, 34].

phase (in ethanol and polyethylene glycol) by placing a magnet bar near the glass bottle. The carbon-encapsulated magnetic metals moved under the magnetic force. This suggests that Ni@C and Fe@C materials would be ideal adsorbents and catalyst supports because they are

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

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61

To evaluate the dispersion stability and agglomeration phenomena of the carbon-encapsulated Ni and Fe nanoparticles in solvents of ethanol and ethylene glycol (EG), their time-dependent sedimentation behavior was investigated using transmission profile measurements obtained with a Turbiscan Lab [36–38]. The transmission profiles were taken every 1 h for 60 h when the suspending medium was ethanol. It was found that the transmission intensity decreased at the sample top owing to clarification and increased at the sample bottom due to sedimentation. A very stable Ni@C dispersion was observed without showing any clarification or sedimentation in EG. In contrast, a progressive fall signal was observed as a function of time in the middle region of Ni nanoparticles which had an average particle size of 20 nm. This can be explained by flocculation-induced particle growth. **Figure 9(a)** shows the Turbiscan screen data taken every 1 h for 3 days. The time-dependent transmission rates on the top, middle, and bottom show the same tendencies. The clarification in the top region and the progressive fall in the middle region of the ΔT signal were not observed in all suspensions. These imply that flocculation due to a coalescing reaction between the nanoparticles was insignificant. A very stable Fe@C dispersion, without any clarification on the top layer or sedimentation on the bottom layer, was observed in ethanol and EG. The viscosity of the solvent affected the dispersion stability kinetics. The dispersion stability of the solvents increased in the following order: water, ethanol, and ethylene glycol (or poly ethylene glycol). **Figure 9(b)** shows the effect of the solvent on the dispersion stability, as measured by using Turbiscan Lab, as well as the calculated mean value of the kinetics for each transmission (ΔT) profile as a function of time. The suspensions prepared in water displayed a rapid change in the mean ΔT values. As a result, sedimentation of the Fe@C nanoparticles in suspensions of water commenced as soon as the suspension was prepared. For the suspensions prepared in ethanol and EG, the variation in the mean ΔT was much less. However, this value increased continuously. Visual inspection confirmed that the suspension was stable, but sedimentation slowly occurred. However, coalescence between the Fe@C nanoparticles rarely occurred in the suspension because the carbon shell layer prevented agglomeration of the particles. The variation in the mean ΔT for the suspension prepared in EG was the smallest. The mean value of ΔBS increased when the particles were smaller than the wavelength of the incident light (880 nm). The tendency of ΔBS was similar to those of ΔT. From these results, for three kinds of solvent,

magnetically separable [35].

**3.2. Dispersion stabilities of Ni, Fe, Ni@C, and Fe@C**

we determined EG to be the most suitable solvent [38, 39].

In this study, we introduce the catalytic effects of the Ni and Ni@C nanopowders observed during the synthesis of S-enantiomer from 3,4-dihydropyrimidine (DHPM). The synthesis

**3.3. Catalyst for the multicomponent Biginelli reaction**

A typical hysteresis loop of the Fe nanopowder at room temperature shows a saturation magnetization of Ms = 157 emu/g and coercivity of Hc = 836 Oe as shown in **Figure 8(a)**. An estimated single domain size of 14 nm for spherical iron particles with no shape anisotropy is reported. The size of the iron nanopowder is large enough to show very large value of coercivity. The hysteresis loops of the as-made Fe@Cs in magnetic fields up to 1 T reveal their intrinsic magnetic behavior, as shown in **Figure 8(b)**.

The hysteresis loops indicate that the carbon-coated Fe nanocrystallites exhibit superparamagnetic behavior at 50 and 300 K. The magnetization was not saturated in the applied fields up to 1 T, as shown in **Figure 6(b)**. In the nanoparticles, one can observe superparamagnetic behavior, which is related to the demagnetization effect arising from the additional energy of the magnetic fields outside the graphitic carbon encapsulation. The coercive force (Hc) and the magnetization (M) at 50 K were 130 Oe and 69.6 (emu/g), respectively. In a previous study, the Mössbauer spectrum for Fe@C nanopowder was measured at room temperature. The relative fraction of the α-Fe, Fe3C, and γ-FeC phases was determined to be about 27.6, 26.3, and 46.1%, respectively. The low magnetization compared with metal nanoparticles without a carbon shell was due to the coexistence of nonmagnetic carbon and the large percentage of surface spins due to the disordered magnetization orientation of the nanoparticles. The magnetic performance of the Ni@C and Fe@C samples was demonstrated in a liquid

**Figure 8.** M-H loops for (a) *α*-Fe and (b) carbon-encapsulated Fe measured at 20°C.

phase (in ethanol and polyethylene glycol) by placing a magnet bar near the glass bottle. The carbon-encapsulated magnetic metals moved under the magnetic force. This suggests that Ni@C and Fe@C materials would be ideal adsorbents and catalyst supports because they are magnetically separable [35].

#### **3.2. Dispersion stabilities of Ni, Fe, Ni@C, and Fe@C**

the carbon-coated Ni nanocrystallites exhibited a superparamagnetic behavior at room temperature, which is related to the demagnetization effect arising from the additional energy of the magnetic fields outside the graphitic carbon encapsulation as shown in **Figure 7(b)**. The coercive force (Hc) and magnetization (M) were 76.6 Oe and 19.6 emu/g, respectively. The ratio of remanence to the saturation magnetization (Mr/M) was 0.04. The low magnetization compared with the Ni nanoparticles without the carbon shell is due to the coexistence of nonmagnetic carbon and the large percentage of surface spin due to the disordered magnetization orientation of the nanoparticles. The magnetic properties are influenced by both the

A typical hysteresis loop of the Fe nanopowder at room temperature shows a saturation magnetization of Ms = 157 emu/g and coercivity of Hc = 836 Oe as shown in **Figure 8(a)**. An estimated single domain size of 14 nm for spherical iron particles with no shape anisotropy is reported. The size of the iron nanopowder is large enough to show very large value of coercivity. The hysteresis loops of the as-made Fe@Cs in magnetic fields up to 1 T reveal their

The hysteresis loops indicate that the carbon-coated Fe nanocrystallites exhibit superparamagnetic behavior at 50 and 300 K. The magnetization was not saturated in the applied fields up to 1 T, as shown in **Figure 6(b)**. In the nanoparticles, one can observe superparamagnetic behavior, which is related to the demagnetization effect arising from the additional energy of the magnetic fields outside the graphitic carbon encapsulation. The coercive force (Hc) and the magnetization (M) at 50 K were 130 Oe and 69.6 (emu/g), respectively. In a previous study, the Mössbauer spectrum for Fe@C nanopowder was measured at room temperature. The relative fraction of the α-Fe, Fe3C, and γ-FeC phases was determined to be about 27.6, 26.3, and 46.1%, respectively. The low magnetization compared with metal nanoparticles without a carbon shell was due to the coexistence of nonmagnetic carbon and the large percentage of surface spins due to the disordered magnetization orientation of the nanoparticles. The magnetic performance of the Ni@C and Fe@C samples was demonstrated in a liquid

particle size and the surface properties of the particle [33, 34].

**Figure 8.** M-H loops for (a) *α*-Fe and (b) carbon-encapsulated Fe measured at 20°C.

intrinsic magnetic behavior, as shown in **Figure 8(b)**.

60 Novel Nanomaterials - Synthesis and Applications

To evaluate the dispersion stability and agglomeration phenomena of the carbon-encapsulated Ni and Fe nanoparticles in solvents of ethanol and ethylene glycol (EG), their time-dependent sedimentation behavior was investigated using transmission profile measurements obtained with a Turbiscan Lab [36–38]. The transmission profiles were taken every 1 h for 60 h when the suspending medium was ethanol. It was found that the transmission intensity decreased at the sample top owing to clarification and increased at the sample bottom due to sedimentation. A very stable Ni@C dispersion was observed without showing any clarification or sedimentation in EG. In contrast, a progressive fall signal was observed as a function of time in the middle region of Ni nanoparticles which had an average particle size of 20 nm. This can be explained by flocculation-induced particle growth. **Figure 9(a)** shows the Turbiscan screen data taken every 1 h for 3 days. The time-dependent transmission rates on the top, middle, and bottom show the same tendencies. The clarification in the top region and the progressive fall in the middle region of the ΔT signal were not observed in all suspensions. These imply that flocculation due to a coalescing reaction between the nanoparticles was insignificant. A very stable Fe@C dispersion, without any clarification on the top layer or sedimentation on the bottom layer, was observed in ethanol and EG. The viscosity of the solvent affected the dispersion stability kinetics. The dispersion stability of the solvents increased in the following order: water, ethanol, and ethylene glycol (or poly ethylene glycol). **Figure 9(b)** shows the effect of the solvent on the dispersion stability, as measured by using Turbiscan Lab, as well as the calculated mean value of the kinetics for each transmission (ΔT) profile as a function of time. The suspensions prepared in water displayed a rapid change in the mean ΔT values. As a result, sedimentation of the Fe@C nanoparticles in suspensions of water commenced as soon as the suspension was prepared. For the suspensions prepared in ethanol and EG, the variation in the mean ΔT was much less. However, this value increased continuously. Visual inspection confirmed that the suspension was stable, but sedimentation slowly occurred. However, coalescence between the Fe@C nanoparticles rarely occurred in the suspension because the carbon shell layer prevented agglomeration of the particles. The variation in the mean ΔT for the suspension prepared in EG was the smallest. The mean value of ΔBS increased when the particles were smaller than the wavelength of the incident light (880 nm). The tendency of ΔBS was similar to those of ΔT. From these results, for three kinds of solvent, we determined EG to be the most suitable solvent [38, 39].

#### **3.3. Catalyst for the multicomponent Biginelli reaction**

In this study, we introduce the catalytic effects of the Ni and Ni@C nanopowders observed during the synthesis of S-enantiomer from 3,4-dihydropyrimidine (DHPM). The synthesis

Vigorous agitation appeared to be an extremely important factor influencing stereo selectivity. The results of stereo selectivity are represented in **Table 1**. The simultaneous use of a heterogeneous catalyst along with the chiral modifier allowed the ratio between stereoisomer in the Biginelli reaction to be changed in some experiments in favor of the S-enantiomer, with an excess of about 19.6%. The best results were obtained when using carbon-encapsulated Ni nanoparticles as the catalyst, L-proline as the chiral modifier, and methanol as the solvent. The catalytic reaction with Ni@C showed higher stereo selectivity than with Ni. The carbon

**Table 1.** Synthesis of 3,4-dihydropyrimidine based on the Biginelli reaction using nanosized catalysts of Ni and Ni@C.

**S-enantiomer R-enantiomer**

**HPLC Δ (ee. %)**

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Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

Cu oxides are widely applied in various organic syntheses such as reduction and oxidation processes, various condensation processes, for the syntheses of complex compounds, etc. The surface of the nanocrystalline Cu oxide includes a defect structure, resulting in nonstoichiometry. Such materials in themselves have the advantages of both homogeneous and heterogeneous catalysts. The aim of our investigation was the development of an effective catalytic and reaction systems based on nanocrystalline Cu oxides, with high reactivity at ambient temperature. To test the catalytic reaction, both the reaction of the liquid-phase oxidation of 2,3,5-trimethyl-1,4-hydroquinone (TMHQ) and the catalase activity were chosen. The oxidation of TMHQ is an intermediate stage of the hydroxylation of 2,3,6-trimethyl phenol in the synthesis of tocopherol. The process of TMHQ oxidation was carried out in a thermostatically controlled chamber, under agitation in a mixed water and methanol solution (1:1 in volume) at 50 ± 0.2°C. The rate of air supply was 6.2 ℓ/h. The reaction was carried out using the parent material (0.66 mmol) and the nanopowders (1 mmol). To compare the catalytic properties of the Cu oxides, powders with various sizes were synthesized. The size control of the powder was carried out by altering the feeding velocity of the Cu wire from 20 to 80 mm/min. Phase control of the Cu, Cu2O, and CuO was carried out by controlling the pressure of the inert gas. In **Table 2**, the crystallite conditions of the copper oxides are displayed. The dehydrogenation (oxidation) of the TMHQ in solution was to practically form 2,3,5-trimethyl-1,4-quinone (TMQ) (selectivity on TMQ > 99.5%). The kinetic curves of TMHQ oxidation in the presence of the nanocrystalline Cu oxide particles (Samples 1, 2, and

**4. Nonmagnetic metal and metal oxides: Cu oxide, Bi, and NiO**

Ni 47 53.2 45.8 7.4 Ni@C 70 59.8 40.2 19.6

3) are given in **Figure 10**. Samples containing mainly pure Cu (Cu 78%, Сu<sup>2</sup>

O 22%, sample a)

shell influences the catalytic effect during synthesis [43].

**4.1. Catalytic activities of Cu oxide and ZnO**

**Catalysts Yield (%)**

**Racemic**

**Figure 9.** Variations in the mean ΔT for the (a) Ni and Ni@C suspensions and (b) Fe and Fe@C suspensions prepared in various solvents (ethanol and ethylene glycol).

of 4-Aryl-substituted DHPM compounds by the Biginelli reaction has attracted great attention in synthetic organic chemistry due to their pharmacological and therapeutic properties such as antibacterial and antihypertensive activity as well as their behavior as calcium channel blockers. Given the versatile biological activity of DHPM, development of an alternative synthetic methodology is of paramount importance [40–42]. This has led to the development of several new synthesis strategies involving combinations of Lewis acids and transition metal salts such as mainly homogeneous catalysts, which give high yields. However, in spite of their potential utility, many of these methods involve expensive reagents, long reaction times, high temperatures, and stoichiometric amounts of catalysts and result in unsatisfactory yields. Therefore, discovering a new, inexpensive catalyst for the Biginelli-type reaction under neutral and mild conditions is of prime importance. The starting materials used in this study were ethyl acetoacetate (I) (0.25 mmol), benzaldehyde (II) (0.25 mmol), and urea (III) (0.3 mmol). First, the benzaldehyde (II) (0.25 mmol), urea (III) (0.3 mmol), 0.1 g of catalyst (Ni or Ni@C), and chiral modifier of L-proline (0.025 mmol) were mixed and react in ethanol (50 ml) at 70°C for 2 hours. In the second step, ethyl acetoacetate (I) (0.25 mmol) was added and reacted under microwave for 3h. The ratio of the s-enantiomer in the asprepared sample was characterized by high-performance liquid chromatography (HPLC) with a chiral column (Chiralcel OD-H) [13].


**Table 1.** Synthesis of 3,4-dihydropyrimidine based on the Biginelli reaction using nanosized catalysts of Ni and Ni@C.

Vigorous agitation appeared to be an extremely important factor influencing stereo selectivity. The results of stereo selectivity are represented in **Table 1**. The simultaneous use of a heterogeneous catalyst along with the chiral modifier allowed the ratio between stereoisomer in the Biginelli reaction to be changed in some experiments in favor of the S-enantiomer, with an excess of about 19.6%. The best results were obtained when using carbon-encapsulated Ni nanoparticles as the catalyst, L-proline as the chiral modifier, and methanol as the solvent. The catalytic reaction with Ni@C showed higher stereo selectivity than with Ni. The carbon shell influences the catalytic effect during synthesis [43].

#### **4. Nonmagnetic metal and metal oxides: Cu oxide, Bi, and NiO**

#### **4.1. Catalytic activities of Cu oxide and ZnO**

of 4-Aryl-substituted DHPM compounds by the Biginelli reaction has attracted great attention in synthetic organic chemistry due to their pharmacological and therapeutic properties such as antibacterial and antihypertensive activity as well as their behavior as calcium channel blockers. Given the versatile biological activity of DHPM, development of an alternative synthetic methodology is of paramount importance [40–42]. This has led to the development of several new synthesis strategies involving combinations of Lewis acids and transition metal salts such as mainly homogeneous catalysts, which give high yields. However, in spite of their potential utility, many of these methods involve expensive reagents, long reaction times, high temperatures, and stoichiometric amounts of catalysts and result in unsatisfactory yields. Therefore, discovering a new, inexpensive catalyst for the Biginelli-type reaction under neutral and mild conditions is of prime importance. The starting materials used in this study were ethyl acetoacetate (I) (0.25 mmol), benzaldehyde (II) (0.25 mmol), and urea (III) (0.3 mmol). First, the benzaldehyde (II) (0.25 mmol), urea (III) (0.3 mmol), 0.1 g of catalyst (Ni or Ni@C), and chiral modifier of L-proline (0.025 mmol) were mixed and react in ethanol (50 ml) at 70°C for 2 hours. In the second step, ethyl acetoacetate (I) (0.25 mmol) was added and reacted under microwave for 3h. The ratio of the s-enantiomer in the asprepared sample was characterized by high-performance liquid chromatography (HPLC)

**Figure 9.** Variations in the mean ΔT for the (a) Ni and Ni@C suspensions and (b) Fe and Fe@C suspensions prepared in

with a chiral column (Chiralcel OD-H) [13].

various solvents (ethanol and ethylene glycol).

62 Novel Nanomaterials - Synthesis and Applications

Cu oxides are widely applied in various organic syntheses such as reduction and oxidation processes, various condensation processes, for the syntheses of complex compounds, etc. The surface of the nanocrystalline Cu oxide includes a defect structure, resulting in nonstoichiometry. Such materials in themselves have the advantages of both homogeneous and heterogeneous catalysts. The aim of our investigation was the development of an effective catalytic and reaction systems based on nanocrystalline Cu oxides, with high reactivity at ambient temperature. To test the catalytic reaction, both the reaction of the liquid-phase oxidation of 2,3,5-trimethyl-1,4-hydroquinone (TMHQ) and the catalase activity were chosen. The oxidation of TMHQ is an intermediate stage of the hydroxylation of 2,3,6-trimethyl phenol in the synthesis of tocopherol. The process of TMHQ oxidation was carried out in a thermostatically controlled chamber, under agitation in a mixed water and methanol solution (1:1 in volume) at 50 ± 0.2°C. The rate of air supply was 6.2 ℓ/h. The reaction was carried out using the parent material (0.66 mmol) and the nanopowders (1 mmol). To compare the catalytic properties of the Cu oxides, powders with various sizes were synthesized. The size control of the powder was carried out by altering the feeding velocity of the Cu wire from 20 to 80 mm/min. Phase control of the Cu, Cu2O, and CuO was carried out by controlling the pressure of the inert gas. In **Table 2**, the crystallite conditions of the copper oxides are displayed. The dehydrogenation (oxidation) of the TMHQ in solution was to practically form 2,3,5-trimethyl-1,4-quinone (TMQ) (selectivity on TMQ > 99.5%). The kinetic curves of TMHQ oxidation in the presence of the nanocrystalline Cu oxide particles (Samples 1, 2, and 3) are given in **Figure 10**. Samples containing mainly pure Cu (Cu 78%, Сu<sup>2</sup> O 22%, sample a)


TMHQ oxidation depended on the particle size and amount of the Сu<sup>2</sup>

which was activated in the matrix of nanopowders [44–46].

Sample 2 with a size of 35 nm, containing mainly Cu2

the total organic carbon (TOC) value was reduced to 60% [48–51].

**4.2. Electrochemical analyses using Bi nanopowder**

Samples 1, 2, and 3 were 17, 35, and 38 m2

high specific surface area [47].

oxidation of the parent material substantially occurred under the action of the fixed oxygen,

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

The catalase activity is an informative parameter about the catalytic properties of the materials in the redox process. It was simulated by measuring the ability of the catalase in the decomposition of hydrogen peroxide to isolate molecular oxygen. Decomposition of the hydrogen peroxide was carried out in a thermostatically isolated chemical reactor (10 ml). A mixture of water and methanol (1:1 in volume) was agitated with a stirring rod at 50 ± 0.2°C. The reaction was carried out with hydrogen peroxide (1.7 mmol) and nanopowder (2 mg). The catalytic activity of the Co, Mn, Fe, and Cu hydroxides was also estimated by catalase activity. The aim of this work was to solve a scientific problem related to the chemical intoxication mechanisms of water phenol solutions and its derivatives during their cleaning. The data in **Table 2** show the results PF the catalase activities for the same nano-Cu oxide samples. The reaction of

Sample 1 with a size of 90 nm and the same oxide phase. The size of the particles was the most significant factor. The particle size affects the surface state. The specific surface areas of

the particular manufacturing method. The LGC method produces nano-scaled powders with

The photocatalytic activity of the ZnO was evaluated based on the photodegradation of phenol aqueous solutions under different irradiation conditions. For experiments under UV-visible light, 100 mL of 50 ppm phenol in aqueous solution with 0.5 g catalytic powders was loaded in a glass container and stirred with a magnetic stirrer a under irradiation form, a Hg-Xe lamp. Total organic carbon (TOC) values as a function of time were measured after filtration under reduced pressure. **Figure 11** shows the photo mineralization of phenol with UV-visible light (solar simulator) in the presence of ZnO. Obviously, when ZnO is added to the phenol,

The anodic stripping voltammetry (ASV) method is a powerful electrochemical technique for trace metal analysis. The traditional electrodes for ASV measurements are mercury-drop electrode and a mercury-film electrode, due to high sensitivity of the mercury [52–56]. However, mercury is very toxic. The toxicity of the mercury has led its usage to be completely banned in some countries. In this study, we focused on searching for alternative environment-friendly electrode materials. The Bi-film electrode has been considered replacing the mercury-film electrode due to its nontoxicity. The properties of Bi materials show not only excellent resolution of neighboring peaks but also insensitivity to the dissolved oxygen in the solution. However, there are still some problems to use the electrode such as a low detection limit comparing to mercury electrode and complication of preparing electrode including processes of additional washing or polishing of the carbon surface and dissolving Bi ions into the solution for the pre-deposition of the Bi on the film electrode. In order to overcome the above weaknesses of the Bi-film electrode, a Bi nanopowder-labeled electrode with

O phase. Obviously,

65

http://dx.doi.org/10.5772/intechopen.72158

O, showed much higher activity than

/g, respectively. The specific surface area is related to

**Table 2.** The concentration of Cu oxide nanopowders, the reaction yields for dehydrogenation of TMHQ, and hydrogen peroxide (catalase activity).

**Figure 10.** Kinetic curves for the dehydrogenation of TMHQ using Cu oxide catalysts with average particle size of 35 and 90 nm.

in the structure showed rare catalytic reaction. Catalytic activity depended on both the average particle size and the oxide phase such as the concentration of Cu2 O in the nanopowders. Sample 1 with mainly Cu2 O phase and an average size of 90 nm performed as catalyst with theoretical yield. Samples 2 and 3 with an average size of 35 nm showed significantly active effect. These samples contained a small quantity of CuO, not exceeding 0.15 mol fraction. The results of the oxidation of TMHQ are represented in **Table 2**. The catalytic yield of Sample 1 compared with Samples 2 and 3 was relatively very low. The Samples 2 and 3 with the same range of particle sizes included different ratios of Cu and Cu oxide phase. The yield of TMHQ oxidation depended on the particle size and amount of the Сu<sup>2</sup> O phase. Obviously, oxidation of the parent material substantially occurred under the action of the fixed oxygen, which was activated in the matrix of nanopowders [44–46].

The catalase activity is an informative parameter about the catalytic properties of the materials in the redox process. It was simulated by measuring the ability of the catalase in the decomposition of hydrogen peroxide to isolate molecular oxygen. Decomposition of the hydrogen peroxide was carried out in a thermostatically isolated chemical reactor (10 ml). A mixture of water and methanol (1:1 in volume) was agitated with a stirring rod at 50 ± 0.2°C. The reaction was carried out with hydrogen peroxide (1.7 mmol) and nanopowder (2 mg). The catalytic activity of the Co, Mn, Fe, and Cu hydroxides was also estimated by catalase activity. The aim of this work was to solve a scientific problem related to the chemical intoxication mechanisms of water phenol solutions and its derivatives during their cleaning. The data in **Table 2** show the results PF the catalase activities for the same nano-Cu oxide samples. The reaction of Sample 2 with a size of 35 nm, containing mainly Cu2 O, showed much higher activity than Sample 1 with a size of 90 nm and the same oxide phase. The size of the particles was the most significant factor. The particle size affects the surface state. The specific surface areas of Samples 1, 2, and 3 were 17, 35, and 38 m2 /g, respectively. The specific surface area is related to the particular manufacturing method. The LGC method produces nano-scaled powders with high specific surface area [47].

The photocatalytic activity of the ZnO was evaluated based on the photodegradation of phenol aqueous solutions under different irradiation conditions. For experiments under UV-visible light, 100 mL of 50 ppm phenol in aqueous solution with 0.5 g catalytic powders was loaded in a glass container and stirred with a magnetic stirrer a under irradiation form, a Hg-Xe lamp. Total organic carbon (TOC) values as a function of time were measured after filtration under reduced pressure. **Figure 11** shows the photo mineralization of phenol with UV-visible light (solar simulator) in the presence of ZnO. Obviously, when ZnO is added to the phenol, the total organic carbon (TOC) value was reduced to 60% [48–51].

#### **4.2. Electrochemical analyses using Bi nanopowder**

in the structure showed rare catalytic reaction. Catalytic activity depended on both the aver-

**Figure 10.** Kinetic curves for the dehydrogenation of TMHQ using Cu oxide catalysts with average particle size of 35

**Table 2.** The concentration of Cu oxide nanopowders, the reaction yields for dehydrogenation of TMHQ, and hydrogen

theoretical yield. Samples 2 and 3 with an average size of 35 nm showed significantly active effect. These samples contained a small quantity of CuO, not exceeding 0.15 mol fraction. The results of the oxidation of TMHQ are represented in **Table 2**. The catalytic yield of Sample 1 compared with Samples 2 and 3 was relatively very low. The Samples 2 and 3 with the same range of particle sizes included different ratios of Cu and Cu oxide phase. The yield of

O phase and an average size of 90 nm performed as catalyst with

**Phase composition, wt. % Oxidized yield** 

4 96 — Initial 60.1

— 100 — Active 99.8

10 85 5 Active 82.3

**O CuO**

78 22 — No

**of TMHQ**

reaction

**conversion Cu Cu2**

**H2 O2**

10.2

O in the nanopowders.

age particle size and the oxide phase such as the concentration of Cu2

Sample 1 with mainly Cu2

**Average particle size (nm)**

peroxide (catalase activity).

**Conditions for synthesis: feeding speed, draft velocity (***l***/min), and** 

**pressure in chamber**

0.0 ≤ VO2≤ 0.05, 80 torr

VO2 = 0.2, 120 torr

VO2 = 0.2, 100 torr

0.1 ≤ VO2≤ 0.15, 100 torr

a 20 Slow feeding (20~30 mm/s)

64 Novel Nanomaterials - Synthesis and Applications

1 90 Fast feeding (60 mm/min)

2 35 Slow feeding(20~30 mm/min)

3 35 Slow feeding (20~30 mm/min)

and 90 nm.

The anodic stripping voltammetry (ASV) method is a powerful electrochemical technique for trace metal analysis. The traditional electrodes for ASV measurements are mercury-drop electrode and a mercury-film electrode, due to high sensitivity of the mercury [52–56]. However, mercury is very toxic. The toxicity of the mercury has led its usage to be completely banned in some countries. In this study, we focused on searching for alternative environment-friendly electrode materials. The Bi-film electrode has been considered replacing the mercury-film electrode due to its nontoxicity. The properties of Bi materials show not only excellent resolution of neighboring peaks but also insensitivity to the dissolved oxygen in the solution. However, there are still some problems to use the electrode such as a low detection limit comparing to mercury electrode and complication of preparing electrode including processes of additional washing or polishing of the carbon surface and dissolving Bi ions into the solution for the pre-deposition of the Bi on the film electrode. In order to overcome the above weaknesses of the Bi-film electrode, a Bi nanopowder-labeled electrode with

partially covered by an insulating layer. Bi nanopowders were well dispersed into 20 ml of distilled water using an ultrasonic treatment. A Nafion solution (Fluka) was added in to the Bi-dispersed suspension for strong chemical bonding between nanopowder and the carbon paste. Finally, the Bi nanopowder-dispersed suspension was dropped onto the working area and dried in the air at room temperature. As the concentration of Nafion in suspension was

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

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67

When the Bi nanoparticles were dispersed in distilled water without Nafion, the zeta potential showed a positive value [14]. However, as Nafion was added in the suspension, the zeta potential changed to a negative value. The amount of Nafion should be optimized to be 200 μℓ for dispersion stability and the phase stability of Bi nanoparticles. The sensor electrodes were prepared using the screen-printed carbon surface with the Bi nanoparticles strongly attached by Nafion. A platinum wire and a saturated calomel electrode (SCE) were used as a counter electrode and a reference electrode, respectively. The supporting electrolyte was a 0.1 M NaAc and 0.025 M HCl solution of pH 5.0. The prepared nanoparticles are confirmed by XRD as shown in **Figure 13(a)**. Also, the screen-printed Bi nanoparticles dispersed in Nafion

**Figure 14** shows results of the anodic stripping voltammograms (ASV) using the Bi nanopowder-attached electrode for measuring various concentrations of Cd and Pb ions in solution. The ASV showed well-defined peaks at −0.85 V and −0.65 V corresponding to the oxidation of Cd and Pb, respectively. **Figure 15** demonstrates the dependence of the stripping peak current density Ip on the Cd and Pb concentrations over a range of 3~30 ppb (deposition potential = −1.35 V and deposition time = 3 min). From the linearity between the metal concentration and the peak current, the values of the sensitivity of the nano-Bi-fixed electrodes

The estimated detection limits of the nano-Bi-fixed electrode were 0.31 and 0.42 ppb for Cd and Pb, respectively, on the basis of the signal-to-noise characteristics (S/N = 3) under a 10 min accumulation. These values are much lower than the domestic and the international content limits of Cd and Pb ions in drinking water, which are listed in **Table 3**, indicating the excellent

for Cd and Pb, respectively.

increased, the value of pH was decreased due to the strong acidity of the Nafion.

on the electrode could be observed by TEM, as shown in **Figure 13(b)**.

were determined to be 9.01 ± 0.012 and 7.15 ± 0.007 μA/ppb·cm2

**Figure 13.** XRD pattern for (a) Bi nanopowders and (b) SEM image for screen-printed Bi.

**Figure 11.** Photo mineralization of phenol with sunlight (TOC: total organic carbon content at times) in the presence of ZnO (Hg-Xe lamp with a wavelength of 200 ∼ 2500 nm and 1 kW of power).

a larger electrochemical active surface area was fabricated [57–59]. In this study, the nano-Bi-fixed electrode sensor and a nanosized Bi-binding technology were developed to improve the electrochemical characteristics of Bi for detecting heavy metals. For this purpose, the Bi nanopowder was synthesized using the LGC method and was then coated on a conductive carbon layer using a Nafion solution. **Figure 12** illustrates the attached Bi working electrode and the analysis system setup for measuring Zn, Cd, Pb, and Ta [14, 24, 58–62].

The working electrode was prepared using conductive carbon ink (DongYoung Chemical Co., LTD, in South Korea) painted flexible polyester film by a semiautomatic screen printing instrument. Then the prepared carbon ink with a thickness of 80 μm on painted thick film was

**Figure 12.** Illustration for working electrode and total system for electrochemical analyses.

partially covered by an insulating layer. Bi nanopowders were well dispersed into 20 ml of distilled water using an ultrasonic treatment. A Nafion solution (Fluka) was added in to the Bi-dispersed suspension for strong chemical bonding between nanopowder and the carbon paste. Finally, the Bi nanopowder-dispersed suspension was dropped onto the working area and dried in the air at room temperature. As the concentration of Nafion in suspension was increased, the value of pH was decreased due to the strong acidity of the Nafion.

When the Bi nanoparticles were dispersed in distilled water without Nafion, the zeta potential showed a positive value [14]. However, as Nafion was added in the suspension, the zeta potential changed to a negative value. The amount of Nafion should be optimized to be 200 μℓ for dispersion stability and the phase stability of Bi nanoparticles. The sensor electrodes were prepared using the screen-printed carbon surface with the Bi nanoparticles strongly attached by Nafion. A platinum wire and a saturated calomel electrode (SCE) were used as a counter electrode and a reference electrode, respectively. The supporting electrolyte was a 0.1 M NaAc and 0.025 M HCl solution of pH 5.0. The prepared nanoparticles are confirmed by XRD as shown in **Figure 13(a)**. Also, the screen-printed Bi nanoparticles dispersed in Nafion on the electrode could be observed by TEM, as shown in **Figure 13(b)**.

**Figure 14** shows results of the anodic stripping voltammograms (ASV) using the Bi nanopowder-attached electrode for measuring various concentrations of Cd and Pb ions in solution. The ASV showed well-defined peaks at −0.85 V and −0.65 V corresponding to the oxidation of Cd and Pb, respectively. **Figure 15** demonstrates the dependence of the stripping peak current density Ip on the Cd and Pb concentrations over a range of 3~30 ppb (deposition potential = −1.35 V and deposition time = 3 min). From the linearity between the metal concentration and the peak current, the values of the sensitivity of the nano-Bi-fixed electrodes were determined to be 9.01 ± 0.012 and 7.15 ± 0.007 μA/ppb·cm2 for Cd and Pb, respectively. The estimated detection limits of the nano-Bi-fixed electrode were 0.31 and 0.42 ppb for Cd and Pb, respectively, on the basis of the signal-to-noise characteristics (S/N = 3) under a 10 min accumulation. These values are much lower than the domestic and the international content limits of Cd and Pb ions in drinking water, which are listed in **Table 3**, indicating the excellent

a larger electrochemical active surface area was fabricated [57–59]. In this study, the nano-Bi-fixed electrode sensor and a nanosized Bi-binding technology were developed to improve the electrochemical characteristics of Bi for detecting heavy metals. For this purpose, the Bi nanopowder was synthesized using the LGC method and was then coated on a conductive carbon layer using a Nafion solution. **Figure 12** illustrates the attached Bi working electrode

**Figure 11.** Photo mineralization of phenol with sunlight (TOC: total organic carbon content at times) in the presence of

The working electrode was prepared using conductive carbon ink (DongYoung Chemical Co., LTD, in South Korea) painted flexible polyester film by a semiautomatic screen printing instrument. Then the prepared carbon ink with a thickness of 80 μm on painted thick film was

and the analysis system setup for measuring Zn, Cd, Pb, and Ta [14, 24, 58–62].

ZnO (Hg-Xe lamp with a wavelength of 200 ∼ 2500 nm and 1 kW of power).

66 Novel Nanomaterials - Synthesis and Applications

**Figure 12.** Illustration for working electrode and total system for electrochemical analyses.

**Figure 13.** XRD pattern for (a) Bi nanopowders and (b) SEM image for screen-printed Bi.

by the XRD method appears to result from the incomplete oxidation of some particles, and it is located in the center of the particles, while the outer layers must certainly be in an oxidized

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

The NiO nanoparticles were applied for the synthesis of both dihydropyridine (DHP) and dihydropyrimidine (DHPM) as mentioned in Section 3.3. The most plausible pathway for DHP prepared by the Hantzsch reaction has been shown to involve the interaction of benzaldehyde with one molecule of β-dicarbonyl compound 1 to give chalcone 3, while another molecule of β-dicarbonyl compound 1 is transformed into enamine 2. In route A, enamine 2 is condensed with an aldehyde and ethyl acetoacetate 1 in the reflux in a suitable solvent (methanol or ethanol) [63–65]. Route B involves the reaction of chalcone 3 with enamine 2, and it seems to give better yields of products and easier purification. In the presence of aqueous ammonia, compound 3 undergoes a partial decomposition into benzaldehyde and diketone 1, thus giving a rise to the formation of symmetrical analogues nitrendipine 16a, b. When the Hantzsch reaction is carried out at 22–25°C in the presence of L-proline and nanosized NiO, the ratio of enantiomers of nitrendipine is changed in favor of the

The Biginelli reaction for synthesizing DHPM was carried out in the presence of L-proline and nanosized NiO (obtained by the Institute of Metal Physics) to change the ratio of enantiomers of 3S in Section 2.3 is changed in favor of S-enantiomer by 15.4%. Our future plans involve studying the factors that affect enantiofacial discrimination for the Hantzsch and Biginelli reaction, such as the nature of nanosized metal oxides or chiral modifiers, reaction time, temperature, and solvent. We also plan to synthesize new nanosized metals and their oxides, as

**N H**

**EtO CO <sup>2</sup>Me 2C**

**\***

**NO <sup>2</sup>**

5 5 5 5

**N H**

**<sup>16</sup> 16a 16b**

**MeO CO <sup>2</sup>Me 2C**

**+ +**

**NO2**

http://dx.doi.org/10.5772/intechopen.72158

**N H**

**EtO 2C CO 2Et**

**NO <sup>2</sup>**

69

3 10

**Heavy metal unit Korea IBWA FDA WHO/FAO**

5 50

**Table 3.** Domestic and international content limits of cd and Pb ions in drinking water.

ppb ppb

state (**Figures 16** and **17**).

S-enantiomer by 3.4%.

**O**

**NO2**

**+ COH**

**NO <sup>2</sup>**

**+**

**CO <sup>2</sup>Me**

**route A**

**route B**

**O NH <sup>2</sup>**

**OEt**

**2**

well as chiral modifiers.

**3**

**MeO <sup>2</sup> C**

**2**

**O**

**O NH <sup>2</sup>**

**OEt**

Cd Pb

**Figure 14.** Square wave anodic stripping voltammograms experimentally measured on the nano-Bi-fixed electrode for various concentrations of Zn, Cd, and Pb ions.

**Figure 15.** Dependence of the stripping peak current density Ip on the Cd and Pb concentrations over the range of 3 ~ 30 ppb (deposition potential = −1.35 V; deposition time = 3 min).

detection of the Bi nanopowder-fixed electrode. Consequently, the low toxicity of the Bi nanopowder-fixed electrode with high sensitivity about heavy metals promises the development of an attractive sensor for monitoring toxic chemical species in environmental matrices with a clean methodology [62].

#### **4.3. Catalytic effect of NiO in the Biginelli and Hantzsch reaction**

Nickel oxide powders were obtained by the gas condensation method in an argon-oxygen mixture flow. The argon flow rate was 130 ℓ/h, oxygen concentration 8.5 vol %, pressure equal to 90 torr, and Ni feed rate 1.8 g/h. XRD analysis showed the following phase composition: NiO 90.7%, Ni 9.3%, and mean size of nickel oxide particles 13 nm. According to the electron microscopy, the particles proved to have a nearly uniaxial shape. The nickel phase detected by the XRD method appears to result from the incomplete oxidation of some particles, and it is located in the center of the particles, while the outer layers must certainly be in an oxidized state (**Figures 16** and **17**).

The NiO nanoparticles were applied for the synthesis of both dihydropyridine (DHP) and dihydropyrimidine (DHPM) as mentioned in Section 3.3. The most plausible pathway for DHP prepared by the Hantzsch reaction has been shown to involve the interaction of benzaldehyde with one molecule of β-dicarbonyl compound 1 to give chalcone 3, while another molecule of β-dicarbonyl compound 1 is transformed into enamine 2. In route A, enamine 2 is condensed with an aldehyde and ethyl acetoacetate 1 in the reflux in a suitable solvent (methanol or ethanol) [63–65]. Route B involves the reaction of chalcone 3 with enamine 2, and it seems to give better yields of products and easier purification. In the presence of aqueous ammonia, compound 3 undergoes a partial decomposition into benzaldehyde and diketone 1, thus giving a rise to the formation of symmetrical analogues nitrendipine 16a, b. When the Hantzsch reaction is carried out at 22–25°C in the presence of L-proline and nanosized NiO, the ratio of enantiomers of nitrendipine is changed in favor of the S-enantiomer by 3.4%.

The Biginelli reaction for synthesizing DHPM was carried out in the presence of L-proline and nanosized NiO (obtained by the Institute of Metal Physics) to change the ratio of enantiomers of 3S in Section 2.3 is changed in favor of S-enantiomer by 15.4%. Our future plans involve studying the factors that affect enantiofacial discrimination for the Hantzsch and Biginelli reaction, such as the nature of nanosized metal oxides or chiral modifiers, reaction time, temperature, and solvent. We also plan to synthesize new nanosized metals and their oxides, as well as chiral modifiers.


**Table 3.** Domestic and international content limits of cd and Pb ions in drinking water.

detection of the Bi nanopowder-fixed electrode. Consequently, the low toxicity of the Bi nanopowder-fixed electrode with high sensitivity about heavy metals promises the development of an attractive sensor for monitoring toxic chemical species in environmental matrices

**Figure 15.** Dependence of the stripping peak current density Ip on the Cd and Pb concentrations over the range of

**Figure 14.** Square wave anodic stripping voltammograms experimentally measured on the nano-Bi-fixed electrode for

Nickel oxide powders were obtained by the gas condensation method in an argon-oxygen mixture flow. The argon flow rate was 130 ℓ/h, oxygen concentration 8.5 vol %, pressure equal to 90 torr, and Ni feed rate 1.8 g/h. XRD analysis showed the following phase composition: NiO 90.7%, Ni 9.3%, and mean size of nickel oxide particles 13 nm. According to the electron microscopy, the particles proved to have a nearly uniaxial shape. The nickel phase detected

**4.3. Catalytic effect of NiO in the Biginelli and Hantzsch reaction**

3 ~ 30 ppb (deposition potential = −1.35 V; deposition time = 3 min).

with a clean methodology [62].

various concentrations of Zn, Cd, and Pb ions.

68 Novel Nanomaterials - Synthesis and Applications

been characterized using a vibrating sample magnetometer (VSM). The size and shape of the nanopowders were investigated by transmission electron microscopy (TEM). The surface effect influenced the magnetic behaviors of nanopowders. Bi metals were dispersed in Nafion. The Bi particles could be applied as sensor electrode instead of mercury-based electrolyte. The particle size of carbon-coated metal with diameters in the range of up to 10 m was smaller than those of metals without a carbon shell. The dispersion stability kinetics of carbon-coated nanopowders showed good dispersion. The best results were obtained when using carbon-encapsulated Ni nanoparticles as a catalyst, L-proline as a chiral modifier, and methanol as a solvent. The catalytic reaction of Ni@C showed enhanced stereoselectivity. Also, the simultaneous use of the heterogeneous catalyst and chiral modifier may lead to an increase in the selectivity of the Biginelli reaction. Nanoparticles prepared using LGC showed significantly enhanced catalytic activities during chemical reaction due to the high

Properties and Catalytic Effects of Nanoparticles Synthesized by Levitational Gas Condensation

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71

Radioisotope Research Division, Korea Atomic Energy Research Institute (KAERI), Daejeon,

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a carbon arc. Physical Review B. 1995;**52**(17):12564-12571

level of defects on their surface structure.

Address all correspondence to: uyrang@kaeri.re.kr

**Author details**

Young Rang Uhm

Republic of Korea

2006;**44**(14):2943-2949

**References**

**Figure 16.** TEM image of NiO nanopowder.

**Figure 17.** Chromatogram of a racemic mixture of (a) DHPM, enriched with S-enantiomer by 15.4%. The product was obtained in the presence of L-proline and NiO nanopowder mic mixture of (b) DHP, enriched with S-enantiomer by 3.4%.

#### **5. Conclusion**

Ceramics, such as NiO, ZnO, and Cu2 O, magnetic nanoparticles, including γ-Fe<sup>2</sup> O3 , Fe3 O4 , and NiFe2 O4, and metals, such as Cu, Ni, Zn, Sn, Ag, Au, Bi, and carbon-encapsulated metals (Ni and Fe), were synthesized by levitational gas condensation (LGC) method using wire feeding (WF) and micron powder feeding (MPF) systems. The magnetic properties have been characterized using a vibrating sample magnetometer (VSM). The size and shape of the nanopowders were investigated by transmission electron microscopy (TEM). The surface effect influenced the magnetic behaviors of nanopowders. Bi metals were dispersed in Nafion. The Bi particles could be applied as sensor electrode instead of mercury-based electrolyte. The particle size of carbon-coated metal with diameters in the range of up to 10 m was smaller than those of metals without a carbon shell. The dispersion stability kinetics of carbon-coated nanopowders showed good dispersion. The best results were obtained when using carbon-encapsulated Ni nanoparticles as a catalyst, L-proline as a chiral modifier, and methanol as a solvent. The catalytic reaction of Ni@C showed enhanced stereoselectivity. Also, the simultaneous use of the heterogeneous catalyst and chiral modifier may lead to an increase in the selectivity of the Biginelli reaction. Nanoparticles prepared using LGC showed significantly enhanced catalytic activities during chemical reaction due to the high level of defects on their surface structure.

### **Author details**

#### Young Rang Uhm

Address all correspondence to: uyrang@kaeri.re.kr

Radioisotope Research Division, Korea Atomic Energy Research Institute (KAERI), Daejeon, Republic of Korea

#### **References**

**5. Conclusion**

and NiFe2

3.4%.

Ceramics, such as NiO, ZnO, and Cu2

**Figure 16.** TEM image of NiO nanopowder.

70 Novel Nanomaterials - Synthesis and Applications

O, magnetic nanoparticles, including γ-Fe<sup>2</sup>

O4, and metals, such as Cu, Ni, Zn, Sn, Ag, Au, Bi, and carbon-encapsulated metals

(Ni and Fe), were synthesized by levitational gas condensation (LGC) method using wire feeding (WF) and micron powder feeding (MPF) systems. The magnetic properties have

**Figure 17.** Chromatogram of a racemic mixture of (a) DHPM, enriched with S-enantiomer by 15.4%. The product was obtained in the presence of L-proline and NiO nanopowder mic mixture of (b) DHP, enriched with S-enantiomer by

> O3 , Fe3 O4 ,


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[63] Marinkovic V, Agbaba D, Vladimirov S, Stankovic S. Simultaneous HPLC determination of nitrendipine and impurities of the process of synthesis. Journal of Pharmaceutical and Biomedical Analysis. 2001;**24**:993-998

**Chapter 5**

Provisional chapter

**Vibrational Behavior of Single-Walled Carbon**

Vibrational Behavior of Single-Walled Carbon

Muzamal Hussain and Muhammad Nawaz Naeem

Muzamal Hussain and Muhammad Nawaz Naeem

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

**Propagation Approach**

Propagation Approach

http://dx.doi.org/10.5772/intechopen.73503

somewhat in agreement.

1. Introduction

single-walled carbon nanotube

Abstract

**Nanotubes Based on Donnell Shell Theory Using Wave**

DOI: 10.5772/intechopen.73503

This chapter is concerned with the vibration analysis of single-walled carbon nanotubes (SWCNTs). This analysis is based on the Donnell thin shell theory. The wave propagation approach in standard eigenvalue form has been employed in order to derive the characteristic frequency equation describing the natural frequencies of vibration in SWCNTs. The axial modal dependence is measured by the complex exponential functions implicating the axial modal numbers. Vibration frequency spectra are gained and evaluated for physical parameter like length-to-diameter ratios. The dimensionless frequency is also investigated in armchair and zigzag SWCNTs with in-plane rigidity and mass density per unit lateral area for armchair and zigzag SWCNTs. These frequencies of the SWCNTs are computed with the aid of the computer software MATLAB. These results are compared with those obtained using molecular dynamics (MD) simulation and the results are

Keywords: vibration analysis, wave propagation approach, Donnell thin shell theory,

Iijima [1] discovered the carbon nanotubes (CNTs) in 1991 and the uses of carbon nanotubes (CNTs) have been originate in various areas such as electronics, optical, medicine, charge detectors, sensors, field emission devices, aerospace, defense, construction and even fashion. To study their remarkable properties, a bulk of research work was performed for their high springiness and characteristic ratio [2], a very effective Young modulus and tensile potency [3], well-bonding

> © 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 IntechOpen. 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.

Nanotubes Based on Donnell Shell Theory Using Wave


#### **Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation Approach** Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation Approach

DOI: 10.5772/intechopen.73503

#### Muzamal Hussain and Muhammad Nawaz Naeem Muzamal Hussain and Muhammad Nawaz Naeem

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.73503

#### Abstract

[63] Marinkovic V, Agbaba D, Vladimirov S, Stankovic S. Simultaneous HPLC determination of nitrendipine and impurities of the process of synthesis. Journal of Pharmaceutical and

[64] Oliver Kappe C. Recent advances in the Biginelli Dihydropyrimidine synthesis new

[65] Atwal KS, Rovnyak, O'Reilly GC, Schwartz. Synthesis of selectivity functionalized 2-Hetero-1,4-dihydropyrimidines. Journal of Organic Chemistry. 1989;**54**:5898-5907

tricks from an old dog. Accounts of Chemical Research. 2000;**33**(12):879-888

Biomedical Analysis. 2001;**24**:993-998

76 Novel Nanomaterials - Synthesis and Applications

This chapter is concerned with the vibration analysis of single-walled carbon nanotubes (SWCNTs). This analysis is based on the Donnell thin shell theory. The wave propagation approach in standard eigenvalue form has been employed in order to derive the characteristic frequency equation describing the natural frequencies of vibration in SWCNTs. The axial modal dependence is measured by the complex exponential functions implicating the axial modal numbers. Vibration frequency spectra are gained and evaluated for physical parameter like length-to-diameter ratios. The dimensionless frequency is also investigated in armchair and zigzag SWCNTs with in-plane rigidity and mass density per unit lateral area for armchair and zigzag SWCNTs. These frequencies of the SWCNTs are computed with the aid of the computer software MATLAB. These results are compared with those obtained using molecular dynamics (MD) simulation and the results are somewhat in agreement.

Keywords: vibration analysis, wave propagation approach, Donnell thin shell theory, single-walled carbon nanotube

#### 1. Introduction

Iijima [1] discovered the carbon nanotubes (CNTs) in 1991 and the uses of carbon nanotubes (CNTs) have been originate in various areas such as electronics, optical, medicine, charge detectors, sensors, field emission devices, aerospace, defense, construction and even fashion. To study their remarkable properties, a bulk of research work was performed for their high springiness and characteristic ratio [2], a very effective Young modulus and tensile potency [3], well-bonding

> © 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 IntechOpen. 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.

strength and superconductivity between carbon atoms [4]. Study of free vibrations of these tube have has been examined with regard to their properties and material behavior. For their useful applications, it needs more explorations to examine vibration characteristics of SWCNTs.

Here an analytical investigation of single-walled carbon nanotubes is conducted for extracting their vibration characteristics. The study of free vibration of SWCNTs is done based on cylindrical shell model. This analysis based on the Donnell thin shell theory. These shell dynamics equations are solved by wave propagation approach. The Donnell shell theory based on WPA is, therefore, another choice of powerful research technique of CNTs whose results are applicable in the limit of acceptable statistical errors than the earlier used BM and other approaches [3–6, 22, 23]. The shell frequency equation is formulated in the eigenvalue form. To provide the complete characteristic of vibrational behavior of SWCNTs by using wave propagation approach is studied in the present chapter. Results are obtained for various material parameters. The dimensionless frequency is also investigated in armchair and zigzag SWCNTs with in-plane rigidity. Now the gape is that there is no research to find directly the dimensionless frequencies of SWCNTs based on cylindrical shell model by using wave propagation approach. However, to the best of authors' knowledge, to find the frequency of SWCNTs, there is no research works on the vibration analyses of zigzag, armchair SWCNTs based on cylindrical shell using wave propagation approach. These frequencies of the SWCNTs are computed with the aid of the computer software MATLAB and these results are compared against MD simulation results in order to assess the accuracy and validity of the cylindrical shell model for predicting the vibration frequencies of

Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation…

http://dx.doi.org/10.5772/intechopen.73503

79

Carbon nanotubes have two kinds, which are single-walled carbon nanotubes and multiwalled carbon nanotubes. Actually multi-walled carbon nanotubes are singled walled carbon nanotubes that are coaxially interposed with different radii. When a graphene sheet rolled up into one time, then it becomes a SWCNTs to produce a hollow cylinder but with end caps. A

Armchair and zigzag nanotubes are made when chiral angle is equal to 0 and 30 respectively and both are the limiting cases with (m, m) and (m, 0). The structure of single-walled carbon nanotubes is similar to the circular cylinders with regard to geometrical shapes as shown in Figure 2. So, the motion equations for cylindrical shells are utilized for studying the free vibrations of SWCNTs. According to the Donnell thin shell theory (He et al. [24]), the governing equation of motion for free vibration of a CNTs is used. Where v1, v2, and v<sup>3</sup> are the longitudinal, circumferential, and radial displacements of the shell, R is the radius of the shell, Eh is the inplane rigidity, ρh is the mass density per unit lateral area, t is the time and ν is the Poisson ratio. It is assumed that for the representation of the modal deformation displacement functions in the axial, circumferential and radial directions are v1ð Þ x; θ; t , v2ð Þ x; θ; t and v3ð Þ x; θ; t correspondingly. The three unknown displacement functions for SWCNTs executing vibra-

schema of graphene sheet and single-walled carbon nanotube are shown in Figure 1.

SWCNTs.

2. Theoretical formation

tion, a system of PDE is given as:

2.1. Cylindrical shell model for the vibration of SWCNT

Poncharal et al. [5] and Treacy et al. [6] conducted the experiments, the resonance frequency of multi-walled carbon nanotubes for clamped-free excited by electrical loads or thermal were detected in a transmission electron microscope (TEM). Thermal vibrations of SWCNTs have been performed for natural frequencies by Molecular dynamics (MD) to predict the Young's modulus by Zhao et al. [7].

The behaviors and material properties of CNTs using either or continuum mechanics modeling or atomistic modeling have been conducted in Wang et al. [8]. For the analysis of CNTs, when compared to continuum mechanics modeling, atomistic modeling is an easy approach and relatively inexpensive. Consequently, the development of continuum mechanics model has attracted much attention of researchers; especially after Yakobson et al. [9] showed that the results obtained using continuum mechanics modeling and molecular dynamics (MD) simulations are in good agreement.

A comprehensive molecular dynamics (MD) study for the contraction and thermal expansion behaviors on different mode of vibration analysis carried by Cao et al. [10]. The fundamental frequency for deformed clamped-clamped shift of SWCNTs under torsion, bending and axial loadings investigated by them. Lordi and Yao [11] performed molecular dynamic (MD) simulations to determine the Young's modulus and thermal vibration frequencies of SWCNTs using the universal force field with various clamped-free conditions based on the Euler beam theory. Carbon nanotubes model of chiral SWCNTs for analyzing their resonant frequency are developed by Hsu et al. [12] and these tubes were observed under a thermal vibration. The model used for implicating the shear deformation and rotatory inertia was Timoshenko beam model. Chawis et al. [13] and Bocko et al. [14] used nonlocal theory of elasticity for the vibration analysis of SWCNTs. An analysis of vibration characteristics of SWCNTs was examined by Yang et al. [15] and initiated this analysis is based on Timoshenko beam model for nonlocal theory. A number of end conditions have interpreted by Azrar et al. [16–17] for the vibrations of these tubes. Recently, vibration behaviors of SWCNTs have been investigated by some researchers [18, 19, 20].

To examine the feasibility of SWCNTs as a nano-resonator, the molecular structural mechanics method was employed by Li and Chou [21] .The predicted fundamental frequencies were perceptive to dimensions such as diameter, length along with boundary conditions clampedfree or clamped SWNTs, but the frequencies are correlatively imperceptive to chirality of the tubes. The vibration and buckling aspects of carbon nanotubes using nonlocal Donnell shell theory was examined by Ansari et al. [22, 23].

Vibration analysis of SWCNTs is examined by using the present approach with clampedclamped and clamped-free vibration. Single-walled carbon nanotubes (SWCNTs) have three distinctive structures as: (i). armchair (ii). zigzag (iii) chiral. These structures have different properties but their vibrational behavior is less clear according to the regarding situation. Vibration analysis of armchair and zigzag type of carbon nanotubes is executed for following boundary conditions: clamped-clamped (C-C), and clamped-free (C-F). Variations of dimensionless frequencies are attained for length-to-diameter ratio.

Here an analytical investigation of single-walled carbon nanotubes is conducted for extracting their vibration characteristics. The study of free vibration of SWCNTs is done based on cylindrical shell model. This analysis based on the Donnell thin shell theory. These shell dynamics equations are solved by wave propagation approach. The Donnell shell theory based on WPA is, therefore, another choice of powerful research technique of CNTs whose results are applicable in the limit of acceptable statistical errors than the earlier used BM and other approaches [3–6, 22, 23]. The shell frequency equation is formulated in the eigenvalue form. To provide the complete characteristic of vibrational behavior of SWCNTs by using wave propagation approach is studied in the present chapter. Results are obtained for various material parameters. The dimensionless frequency is also investigated in armchair and zigzag SWCNTs with in-plane rigidity. Now the gape is that there is no research to find directly the dimensionless frequencies of SWCNTs based on cylindrical shell model by using wave propagation approach. However, to the best of authors' knowledge, to find the frequency of SWCNTs, there is no research works on the vibration analyses of zigzag, armchair SWCNTs based on cylindrical shell using wave propagation approach. These frequencies of the SWCNTs are computed with the aid of the computer software MATLAB and these results are compared against MD simulation results in order to assess the accuracy and validity of the cylindrical shell model for predicting the vibration frequencies of SWCNTs.

### 2. Theoretical formation

strength and superconductivity between carbon atoms [4]. Study of free vibrations of these tube have has been examined with regard to their properties and material behavior. For their useful applications, it needs more explorations to examine vibration characteristics of SWCNTs.

Poncharal et al. [5] and Treacy et al. [6] conducted the experiments, the resonance frequency of multi-walled carbon nanotubes for clamped-free excited by electrical loads or thermal were detected in a transmission electron microscope (TEM). Thermal vibrations of SWCNTs have been performed for natural frequencies by Molecular dynamics (MD) to predict the Young's

The behaviors and material properties of CNTs using either or continuum mechanics modeling or atomistic modeling have been conducted in Wang et al. [8]. For the analysis of CNTs, when compared to continuum mechanics modeling, atomistic modeling is an easy approach and relatively inexpensive. Consequently, the development of continuum mechanics model has attracted much attention of researchers; especially after Yakobson et al. [9] showed that the results obtained using continuum mechanics modeling and molecular dynamics (MD) simula-

A comprehensive molecular dynamics (MD) study for the contraction and thermal expansion behaviors on different mode of vibration analysis carried by Cao et al. [10]. The fundamental frequency for deformed clamped-clamped shift of SWCNTs under torsion, bending and axial loadings investigated by them. Lordi and Yao [11] performed molecular dynamic (MD) simulations to determine the Young's modulus and thermal vibration frequencies of SWCNTs using the universal force field with various clamped-free conditions based on the Euler beam theory. Carbon nanotubes model of chiral SWCNTs for analyzing their resonant frequency are developed by Hsu et al. [12] and these tubes were observed under a thermal vibration. The model used for implicating the shear deformation and rotatory inertia was Timoshenko beam model. Chawis et al. [13] and Bocko et al. [14] used nonlocal theory of elasticity for the vibration analysis of SWCNTs. An analysis of vibration characteristics of SWCNTs was examined by Yang et al. [15] and initiated this analysis is based on Timoshenko beam model for nonlocal theory. A number of end conditions have interpreted by Azrar et al. [16–17] for the vibrations of these tubes. Recently,

vibration behaviors of SWCNTs have been investigated by some researchers [18, 19, 20].

To examine the feasibility of SWCNTs as a nano-resonator, the molecular structural mechanics method was employed by Li and Chou [21] .The predicted fundamental frequencies were perceptive to dimensions such as diameter, length along with boundary conditions clampedfree or clamped SWNTs, but the frequencies are correlatively imperceptive to chirality of the tubes. The vibration and buckling aspects of carbon nanotubes using nonlocal Donnell shell

Vibration analysis of SWCNTs is examined by using the present approach with clampedclamped and clamped-free vibration. Single-walled carbon nanotubes (SWCNTs) have three distinctive structures as: (i). armchair (ii). zigzag (iii) chiral. These structures have different properties but their vibrational behavior is less clear according to the regarding situation. Vibration analysis of armchair and zigzag type of carbon nanotubes is executed for following boundary conditions: clamped-clamped (C-C), and clamped-free (C-F). Variations of dimen-

modulus by Zhao et al. [7].

78 Novel Nanomaterials - Synthesis and Applications

tions are in good agreement.

theory was examined by Ansari et al. [22, 23].

sionless frequencies are attained for length-to-diameter ratio.

#### 2.1. Cylindrical shell model for the vibration of SWCNT

Carbon nanotubes have two kinds, which are single-walled carbon nanotubes and multiwalled carbon nanotubes. Actually multi-walled carbon nanotubes are singled walled carbon nanotubes that are coaxially interposed with different radii. When a graphene sheet rolled up into one time, then it becomes a SWCNTs to produce a hollow cylinder but with end caps. A schema of graphene sheet and single-walled carbon nanotube are shown in Figure 1.

Armchair and zigzag nanotubes are made when chiral angle is equal to 0 and 30 respectively and both are the limiting cases with (m, m) and (m, 0). The structure of single-walled carbon nanotubes is similar to the circular cylinders with regard to geometrical shapes as shown in Figure 2. So, the motion equations for cylindrical shells are utilized for studying the free vibrations of SWCNTs. According to the Donnell thin shell theory (He et al. [24]), the governing equation of motion for free vibration of a CNTs is used. Where v1, v2, and v<sup>3</sup> are the longitudinal, circumferential, and radial displacements of the shell, R is the radius of the shell, Eh is the inplane rigidity, ρh is the mass density per unit lateral area, t is the time and ν is the Poisson ratio.

It is assumed that for the representation of the modal deformation displacement functions in the axial, circumferential and radial directions are v1ð Þ x; θ; t , v2ð Þ x; θ; t and v3ð Þ x; θ; t correspondingly. The three unknown displacement functions for SWCNTs executing vibration, a system of PDE is given as:

$$\frac{\partial^2 v\_1}{\partial \mathbf{x}^2} + \frac{1 - \nu}{2R^2} \frac{\partial^2 v\_1}{\partial \theta^2} + \frac{1 + \nu}{2R} \frac{\partial^2 v\_2}{\partial \mathbf{x} \partial \theta} - \frac{\nu}{R} \frac{\partial v\_3}{\partial \mathbf{x}} = \frac{\left(1 - \nu^2\right) \rho h}{Eh} \frac{\partial^2 v\_1}{\partial t^2} \tag{1}$$

2.2. Applications of the wave propagation approach

are written in the assumed expression as:

SWCNTs, an eigenvalue problem is formed:

The form of non-zero solution of pm; qm;rm

m, n in order to obtain the frequency of vibration.

2 6 4

3. Result and discussion

L<sup>11</sup> L<sup>12</sup> L<sup>13</sup> L<sup>21</sup> L<sup>22</sup> L<sup>23</sup> L<sup>31</sup> L<sup>32</sup> L<sup>33</sup>

An efficient and a simple technique which corporate as wave propagation approach is employed for the solution of CNT problem in the form of differential equation. Before this, present method has been successively used for the study of shell vibrations [25–27]. The axial coordinate and time variable are denoted by x, t correspondingly and the circumferential coordinate signifies by θ. The functions v1ð Þ x; θ; t , v2ð Þ x; θ; t and v3ð Þ x; θ; t are used to designate their respective displacement deformation function. So for modal deformation displacements

Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation…

�ikmx cos ð Þ <sup>n</sup><sup>θ</sup> <sup>e</sup>

�ikmx sin ð Þ <sup>n</sup><sup>θ</sup> <sup>e</sup>

�ikmx cos ð Þ <sup>n</sup><sup>θ</sup> <sup>e</sup>

where pm, qm and rm stand for three vibration amplitude coefficients in the axial, circumferential and radial directions. The axial half and the circumferential wave numbers are denoted by m and n respectively and angular frequency is designated by ω. The formula for fundamental frequency f which is written as: f ¼ ω=2π: Where km is the axial wave number related with an end conditions. Using the expressions for v1ð Þ x; θ; t , v2ð Þ x; θ; t , v3ð Þ x; θ; t and their partial derivatives in applying the product method by substituting the modal displacement functions

After putting Eqs. (4)–(6), into Eqs. (1)–(3), the above equations is transmuted in matrix representation after the arrangement of terms, and to designate the vibration frequency equation for

CA <sup>¼</sup> <sup>1</sup> � <sup>ν</sup><sup>2</sup> � �ρ<sup>h</sup>

modes for SWCNTs. The expressions for the terms Lij 's are given in the Appendix-I. Where the roots of the equation furnish the frequencies. The lowest root corresponds to the frequency of vibration. It is clear that the frequency should be minimized with respect to the wave numbers

The vibration frequency spectra for SWCNTs are evaluated by Eq. (3) based on Donnell thin cylindrical shell theory. Variations of the frequencies are obtained with regard to the material properties and tube thickness. Keeping in view of this aspect, the natural frequencies of the

Eh <sup>ω</sup><sup>2</sup>

100 010 001

� � yields the vibration frequency and associated

3 7 5

0 B@

pm qm rm

1

CA (7)

2 6 4 <sup>ω</sup><sup>t</sup> (4)

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81

<sup>ω</sup><sup>t</sup> (5)

<sup>ω</sup><sup>t</sup> (6)

v1ð Þ¼ x; θ; t pme

v2ð Þ¼ x; θ; t qme

v3ð Þ¼ x; θ; t rme

for partial differential equations, the space and time variable are split.

3 7 5

0 B@

pm qm rm

1

$$\frac{1+\nu}{2R}\frac{\partial^2 v\_1}{\partial x \partial \theta} + \frac{1-\nu}{2}\frac{\partial^2 v\_2}{\partial x^2} + \frac{1}{R^2}\frac{\partial^2 v\_2}{\partial \theta^2} - \frac{1}{R^2}\frac{\partial v\_3}{\partial \theta} = \frac{(1-\nu^2)\rho h}{Eh}\frac{\partial^2 v\_2}{\partial t^2} \tag{2}$$

$$\frac{\nu}{R}\frac{\partial v\_1}{\partial x} + \frac{1}{R^2}\frac{\partial v\_2}{\partial \theta} - \left(\frac{1}{R^2} + \frac{\left(1 - \nu^2\right)}{Eh}D\left(\frac{\partial^4 v\_3}{\partial x^4} + 2.\frac{1}{R^2}\frac{\partial^4 v\_3}{\partial x^2 \partial \theta^2} + \frac{1}{R^4}\frac{\partial^4 v\_3}{\partial \theta^4}\right)\right) = \frac{\left(1 - \nu^2\right)\rho h}{Eh}\frac{\partial^2 v\_3}{\partial t^2} \tag{3}$$

where <sup>D</sup> <sup>¼</sup> Eh<sup>3</sup> 12 1�v<sup>2</sup> ð Þ denotes the effective bending stiffness.

Figure 1. Hexagonal lattice (a) graphene sheet (b) single-walled carbon nanotube.

Figure 2. Geometry of SWCNTs.

#### 2.2. Applications of the wave propagation approach

∂2 v1 ∂x<sup>2</sup> þ

80 Novel Nanomaterials - Synthesis and Applications

1 þ ν 2R

ν R ∂v<sup>1</sup> ∂x þ 1 R2 ∂v<sup>2</sup> <sup>∂</sup><sup>θ</sup> � <sup>1</sup>

where <sup>D</sup> <sup>¼</sup> Eh<sup>3</sup>

Figure 2. Geometry of SWCNTs.

∂2 v1 ∂x∂θ þ

<sup>R</sup><sup>2</sup> <sup>þ</sup>

1 � ν 2R<sup>2</sup>

∂2 v1 <sup>∂</sup>θ<sup>2</sup> <sup>þ</sup>

1 � ν 2

12 1�v<sup>2</sup> ð Þ denotes the effective bending stiffness.

Figure 1. Hexagonal lattice (a) graphene sheet (b) single-walled carbon nanotube.

<sup>1</sup> � <sup>ν</sup><sup>2</sup> Eh :<sup>D</sup> <sup>∂</sup><sup>4</sup>

∂2 v2 ∂x<sup>2</sup> þ

1 þ ν 2R

∂2 v2 <sup>∂</sup>x∂<sup>θ</sup> � <sup>ν</sup> R ∂v<sup>3</sup>

1 R2 ∂2 v2 <sup>∂</sup>θ<sup>2</sup> � <sup>1</sup> R2 ∂v<sup>3</sup>

∂4 v3 <sup>∂</sup>x<sup>2</sup>∂θ<sup>2</sup> <sup>þ</sup>

v3 <sup>∂</sup>x<sup>4</sup> <sup>þ</sup> <sup>2</sup>: <sup>1</sup> R2 <sup>∂</sup><sup>x</sup> <sup>¼</sup> <sup>1</sup> � <sup>ν</sup><sup>2</sup> <sup>ρ</sup><sup>h</sup> Eh

> <sup>∂</sup><sup>θ</sup> <sup>¼</sup> <sup>1</sup> � <sup>ν</sup><sup>2</sup> <sup>ρ</sup><sup>h</sup> Eh

> > 1 R4 ∂4 v3 ∂θ<sup>4</sup>

∂2 v1 ∂t

> ∂2 v2 ∂t

<sup>¼</sup> <sup>1</sup> � <sup>ν</sup><sup>2</sup> <sup>ρ</sup><sup>h</sup> Eh

<sup>2</sup> (1)

<sup>2</sup> (2)

∂2 v3 ∂t <sup>2</sup> (3) An efficient and a simple technique which corporate as wave propagation approach is employed for the solution of CNT problem in the form of differential equation. Before this, present method has been successively used for the study of shell vibrations [25–27]. The axial coordinate and time variable are denoted by x, t correspondingly and the circumferential coordinate signifies by θ. The functions v1ð Þ x; θ; t , v2ð Þ x; θ; t and v3ð Þ x; θ; t are used to designate their respective displacement deformation function. So for modal deformation displacements are written in the assumed expression as:

$$\left. \psi\_1(\mathbf{x}, \theta, t) = p\_m e^{-ik\_m x} \cos \left( n\theta \right) e^{\omega t} \right. \tag{4}$$

$$\upsilon\_2(\mathbf{x}, \theta, t) = q\_m e^{-i\mathbf{k}\_m \mathbf{x}} \sin \left( n\theta \right) e^{\omega t} \tag{5}$$

$$\sigma\_3(\mathbf{x}, \theta, t) = r\_m e^{-ik\_m \mathbf{x}} \cos \left( n\theta \right) e^{\omega t} \tag{6}$$

where pm, qm and rm stand for three vibration amplitude coefficients in the axial, circumferential and radial directions. The axial half and the circumferential wave numbers are denoted by m and n respectively and angular frequency is designated by ω. The formula for fundamental frequency f which is written as: f ¼ ω=2π: Where km is the axial wave number related with an end conditions. Using the expressions for v1ð Þ x; θ; t , v2ð Þ x; θ; t , v3ð Þ x; θ; t and their partial derivatives in applying the product method by substituting the modal displacement functions for partial differential equations, the space and time variable are split.

After putting Eqs. (4)–(6), into Eqs. (1)–(3), the above equations is transmuted in matrix representation after the arrangement of terms, and to designate the vibration frequency equation for SWCNTs, an eigenvalue problem is formed:

$$
\begin{bmatrix} L\_{11} & L\_{12} & L\_{13} \\ L\_{21} & L\_{22} & L\_{23} \\ L\_{31} & L\_{32} & L\_{33} \end{bmatrix} \begin{pmatrix} p\_m \\ q\_m \\ r\_m \end{pmatrix} = \frac{(1-\nu^2)\rho h}{E\hbar} \omega^2 \begin{bmatrix} 1 & 0 & 0 \\ 0 & 1 & 0 \\ 0 & 0 & 1 \end{bmatrix} \begin{pmatrix} p\_m \\ q\_m \\ r\_m \end{pmatrix} \tag{7}
$$

The form of non-zero solution of pm; qm;rm � � yields the vibration frequency and associated modes for SWCNTs. The expressions for the terms Lij 's are given in the Appendix-I. Where the roots of the equation furnish the frequencies. The lowest root corresponds to the frequency of vibration. It is clear that the frequency should be minimized with respect to the wave numbers m, n in order to obtain the frequency of vibration.

#### 3. Result and discussion

The vibration frequency spectra for SWCNTs are evaluated by Eq. (3) based on Donnell thin cylindrical shell theory. Variations of the frequencies are obtained with regard to the material properties and tube thickness. Keeping in view of this aspect, the natural frequencies of the longitudinal clamped-free vibration of SWCNTs with a length 6.92 nm are first determined by the MD simulation. The E=ρ ratio as measured from molecular dynamic simulation is 3.6481 <sup>10</sup><sup>8</sup> m<sup>2</sup>=s2. By applying this ratio on the longer tube having length 14.4 nm is simulated by MD simulation. The results were found to be same, demonstrating the unconventionality of the ratio on the length. We shall adopt the material properties and tube thickness as suggested by Zhang et al. [18], i.e. the in-plane rigidity Eh =278.25 Gpanm, <sup>E</sup>=ρ=3.6481 <sup>10</sup><sup>8</sup> m<sup>2</sup>=s2, Poisson's ratio ν = 0.2. The in-plane stiffness or rigidity is computed as Eh = 278.25 GPa.nm which is based on the E=ρ ratio, to calculate the natural and dimensional frequencies of vibration using wave propagation approach, the ratio <sup>E</sup>=ρ= 3.6481 <sup>10</sup><sup>8</sup> m<sup>2</sup>=s<sup>2</sup> is used throughout this study. For example, the range of reported thickness is from 0.0612 to 0.69 nm in [9, 28–29] and v varies from 0.14 to 0.34. Considering a diameter <sup>d</sup> = 6.86645 <sup>10</sup><sup>10</sup> m, the vibration frequencies for Single-walled carbon nanotubes of various length-to-diameter ratios are calculated using Eq. (3). In present model, the effects of different length-to-diameter ratio for clamped-clamped and clamped-free boundary condition have been considered and matched quantitatively with MD results as well as for the validity and to assure the accuracy. In this study, all frequency results are presented in THz unless otherwise stated. In present study, the frequencies of SWCNTs are obtained by using the some parameters which are compared with MD simulation and continuum shell. However the MD results were obtained for a clamped-clamped SWCNT. Two sets of result are compared as shown in the Tables 1 and 2.

compared with molecular dynamic simulations and Timoshenko beam model with thicknesses h = 0.34 nm. The results are in good agreement with the MD and Timoshenko beam model

4.67 0.17074 0.23193 �26.38 6.47 0.09048 0.12872 �29.70 7.55 0.06678 0.1000 �31.61 8.28 0.05566 0.07935 �29.85 10.07 0.03777 0.05493 �31.23

Table 2. Comparison of frequencies of clamed-free SWCNT for the first vibration mode.

Present MD Percentage error

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83

Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation…

The first and third mode natural frequencies accessed by present model and compared with MD simulations are depicted graphically in Figure 3 respectively. It can be observed from Tables 1 and 2 that the length-to-diameter ratio of the set of C-C SWCNTs is somewhat dissimilar from that of C-F SWCNTs. For the prediction of mechanical characters [7, 19–21] from atomistic studies and the experimental studies [5, 6, 33–38] are often used for the clamped-free carbon nanotubes. The frequencies for the first and third modes obtained from present model which is compared with molecular dynamic simulation and Timoshenko beam model are shown in Figure 3. It can be readily seen that higher frequencies are produced by higher modes and when length-to-diameter ratio rises at each mode then frequency falls down smoothly as shown in Figure 3. The relationship between the length-to-diameter ratios and

Figure 3. Comparison of numerically obtained results for clamped-clamped and clamped-free frequencies of SWCNTs for first and third mode versus length-to-diameter ratio L/d with MD by Cao et al. [32] and Timoshenko beam model [18].

results showing same trend in the open literature.

L/d Frequencies (THz)

It can be observed that the results which are obtained from present model, the values are nearer to the molecular dynamics results when the length-to-diameter ratio is greater than 10.26. From Table 1, one can notice that the average percentage error between MD results is approximately 3.6%. This fact shows that the results obtained by present method and earlier MD simulation model are in good agreement.

#### 3.1. Vibration of clamped-clamped and clamped-free SWCNTs

In the application of micro-oscillators and micro or nano-strain sensors, the carbon nanotube sensor is generally clamped with both ends [30]. The clamped-clamped and clamped-free single-walled carbon nanotubes have been performed by atomistic simulations [20–21, 31]. In this section, the distinctive first and third mode frequencies for the set of clamped-clamped single-walled carbon nanotubes is given by present models for their vibration frequencies and


Table 1. Comparison of frequencies of C-C SWCNT with MD simulation for the first vibration mode.

Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation… http://dx.doi.org/10.5772/intechopen.73503 83


Table 2. Comparison of frequencies of clamed-free SWCNT for the first vibration mode.

longitudinal clamped-free vibration of SWCNTs with a length 6.92 nm are first determined by the MD simulation. The E=ρ ratio as measured from molecular dynamic simulation is

simulated by MD simulation. The results were found to be same, demonstrating the unconventionality of the ratio on the length. We shall adopt the material properties and tube thickness as suggested by Zhang et al. [18], i.e. the in-plane rigidity Eh =278.25 Gpanm,

as Eh = 278.25 GPa.nm which is based on the E=ρ ratio, to calculate the natural and dimensional frequencies of vibration using wave propagation approach, the ratio <sup>E</sup>=ρ= 3.6481 <sup>10</sup><sup>8</sup> m<sup>2</sup>=s<sup>2</sup> is used throughout this study. For example, the range of reported thickness is from 0.0612 to 0.69 nm in [9, 28–29] and v varies from 0.14 to 0.34. Considering a diameter <sup>d</sup> = 6.86645 <sup>10</sup><sup>10</sup> m, the vibration frequencies for Single-walled carbon nanotubes of various length-to-diameter ratios are calculated using Eq. (3). In present model, the effects of different length-to-diameter ratio for clamped-clamped and clamped-free boundary condition have been considered and matched quantitatively with MD results as well as for the validity and to assure the accuracy. In this study, all frequency results are presented in THz unless otherwise stated. In present study, the frequencies of SWCNTs are obtained by using the some parameters which are compared with MD simulation and continuum shell. However the MD results were obtained for a clamped-clamped SWCNT. Two sets of result are

It can be observed that the results which are obtained from present model, the values are nearer to the molecular dynamics results when the length-to-diameter ratio is greater than 10.26. From Table 1, one can notice that the average percentage error between MD results is approximately 3.6%. This fact shows that the results obtained by present method and earlier

In the application of micro-oscillators and micro or nano-strain sensors, the carbon nanotube sensor is generally clamped with both ends [30]. The clamped-clamped and clamped-free single-walled carbon nanotubes have been performed by atomistic simulations [20–21, 31]. In this section, the distinctive first and third mode frequencies for the set of clamped-clamped single-walled carbon nanotubes is given by present models for their vibration frequencies and

Present MD Percentage error

m<sup>2</sup>=s2. By applying this ratio on the longer tube having length 14.4 nm is

m<sup>2</sup>=s2, Poisson's ratio ν = 0.2. The in-plane stiffness or rigidity is computed

3.6481 <sup>10</sup><sup>8</sup>

82 Novel Nanomaterials - Synthesis and Applications

<sup>E</sup>=ρ=3.6481 <sup>10</sup><sup>8</sup>

compared as shown in the Tables 1 and 2.

MD simulation model are in good agreement.

L/d Frequencies (THz)

3.1. Vibration of clamped-clamped and clamped-free SWCNTs

6.67 0.67832 0.64697 4.85 8.47 0.44146 0.43335 1.87 10.26 0.30922 0.30518 1.32 13.89 0.17360 0.18311 5.19

Table 1. Comparison of frequencies of C-C SWCNT with MD simulation for the first vibration mode.

compared with molecular dynamic simulations and Timoshenko beam model with thicknesses h = 0.34 nm. The results are in good agreement with the MD and Timoshenko beam model results showing same trend in the open literature.

The first and third mode natural frequencies accessed by present model and compared with MD simulations are depicted graphically in Figure 3 respectively. It can be observed from Tables 1 and 2 that the length-to-diameter ratio of the set of C-C SWCNTs is somewhat dissimilar from that of C-F SWCNTs. For the prediction of mechanical characters [7, 19–21] from atomistic studies and the experimental studies [5, 6, 33–38] are often used for the clamped-free carbon nanotubes. The frequencies for the first and third modes obtained from present model which is compared with molecular dynamic simulation and Timoshenko beam model are shown in Figure 3. It can be readily seen that higher frequencies are produced by higher modes and when length-to-diameter ratio rises at each mode then frequency falls down smoothly as shown in Figure 3. The relationship between the length-to-diameter ratios and

Figure 3. Comparison of numerically obtained results for clamped-clamped and clamped-free frequencies of SWCNTs for first and third mode versus length-to-diameter ratio L/d with MD by Cao et al. [32] and Timoshenko beam model [18].

natural frequencies is inversely proportional indicates that the vibrations are very sensitive due to long tube and since the SWCNTs are of almost the same diameter. The results for SWCNTs given by MD are little bit higher than the frequencies investigated by the present model. In MD simulation, the frequencies of length-to-radius ratio are 8.28 is 0.0793 and at 20.89 is 0.0138. But in the present model, the frequencies at 8.28 are 0.05566 and at 20.89 is 0.00883, when compared to the MD results.

zigzag CNT, changing the length-to-diameter ratio from 4.86 to 35.53, the dimensionless frequency changes from 0.0303 to 0.0049 THz in case of clamped-clamped boundary condition.

Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation…

For the results generated so far, the nanotube in-plane rigidity has been taken to be Eh =278.25 Gpa�nm. However, there exist some inconsistencies concerning this quantity in the literature. The reported CNT in-plane stiffness is largely scattered, ranging from Eh =300 Gpa�nm to Eh = 400 Gpa�nm [42]. Both set of Figure 5 is presented to investigate the influence of the in-plane rigidity Eh variation on the dimensionless frequency of a (7, 7) armchair and (9, 0) Zigzag SWCNT with different boundary conditions likewise as clamped-clamped and clamped-free boundary conditions. These figures shows that for all the selected boundary conditions, dimensionless frequency calculated via shell model are sensitive to the nanotube in-plane rigidity Eh and also the larger the in-plane rigidity in-plane rigidity Eh, the higher the dimen-

Previous study reveals that the bending rigidity of SWCNTs should be considered as an independent material parameter not linked to the representative thickness by the classic

is lesser than its classical counterpart [43, 44]. For shorter length-to-diameter ratio, the value of dimensionless frequencies for clamped-clamped at Eh = 300 Gpa�nm, Eh = 400 Gpa�nm is 0.04863, 0.05788, respectively which shows that a slight increase in frequency due to increase of in-plane rigidity Eh. Same trend is observed for dimensionless frequency. For the present shell model with in-plane rigidity Eh, the values of the C-F single-walled carbon nanotubes respectively, which are a little lower than those of corresponding CC SWCNTs with bending

Figure 5. Variations of dimensionless frequencies Ω of CC and CF armchair and zigzag SWCNTs when Eh = 300 GPa.Nm

<sup>=</sup>12 1 � <sup>ν</sup><sup>2</sup> and the actual bending rigidity of SWCNTs

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85

Likewise, in clamped-free condition it varies from 1.16605 to 0.4609 THz.

sionless frequency. The difference is more considerable for shorter length CNTs.

3.3. Vibration of SWCNTs with in-plane rigidity

bending rigidity formula, i.e., <sup>D</sup> <sup>¼</sup> Eh<sup>3</sup>

rigidity are plotted in Figure 5.

and Eh = 425Gpa.Nm.

#### 3.2. Vibration of SWCNTs with dimensionless frequency

Furthermore, the parametric study for the vibrational behavior of SWCNTs with dimensionless is carried out and presented in Figure 4. Alibeigloo et al. [39, 40] and Soldatos et al. [41] used the dimensionless frequency for multi-walled carbon nanotubes and for thin cylindrical shell with respect to length-to-radius ratio respectively. This frequency is associated with frequency Ω through the following formula: Ω ¼ ωR ffiffiffiffi ρ E : q A variation of non-dimensional frequency versus length-to-diameter ratio is presented in Figure 4. This figure shows that, increasing the value of length-to-diameter ratio as well as there is a decrease in dimensionless frequency. From the physical point of view, it is noted that when the length of SWCNTs becomes small, the effect of the atomic interactions among a reference point and all other atoms becomes significant.

Figure 4 shows an armchair and zigzag CNT, vary the length-to-diameter ratio from 8.3 to 20.9 will change the dimensionless frequency from 0.0385 to 0.0063 THz in case of clampedclamped boundary condition. Likewise, in clamped-free condition it changes from 1.2827 to 0.5094 THz in armchair case. It may be seen from the above Figure 4 that the resulting value of dimensionless frequency decreases with the increase in length-to-diameter ratio. Now in

Figure 4. Variations of dimensionless frequencies of CC and CF armchair and zigzag SWCNTs.

zigzag CNT, changing the length-to-diameter ratio from 4.86 to 35.53, the dimensionless frequency changes from 0.0303 to 0.0049 THz in case of clamped-clamped boundary condition. Likewise, in clamped-free condition it varies from 1.16605 to 0.4609 THz.

#### 3.3. Vibration of SWCNTs with in-plane rigidity

natural frequencies is inversely proportional indicates that the vibrations are very sensitive due to long tube and since the SWCNTs are of almost the same diameter. The results for SWCNTs given by MD are little bit higher than the frequencies investigated by the present model. In MD simulation, the frequencies of length-to-radius ratio are 8.28 is 0.0793 and at 20.89 is 0.0138. But in the present model, the frequencies at 8.28 are 0.05566 and at 20.89 is

Furthermore, the parametric study for the vibrational behavior of SWCNTs with dimensionless is carried out and presented in Figure 4. Alibeigloo et al. [39, 40] and Soldatos et al. [41] used the dimensionless frequency for multi-walled carbon nanotubes and for thin cylindrical shell with respect to length-to-radius ratio respectively. This frequency is associated with

frequency versus length-to-diameter ratio is presented in Figure 4. This figure shows that, increasing the value of length-to-diameter ratio as well as there is a decrease in dimensionless frequency. From the physical point of view, it is noted that when the length of SWCNTs becomes small, the effect of the atomic interactions among a reference point and all other

Figure 4 shows an armchair and zigzag CNT, vary the length-to-diameter ratio from 8.3 to 20.9 will change the dimensionless frequency from 0.0385 to 0.0063 THz in case of clampedclamped boundary condition. Likewise, in clamped-free condition it changes from 1.2827 to 0.5094 THz in armchair case. It may be seen from the above Figure 4 that the resulting value of dimensionless frequency decreases with the increase in length-to-diameter ratio. Now in

Figure 4. Variations of dimensionless frequencies of CC and CF armchair and zigzag SWCNTs.

ffiffiffiffi ρ E : q

A variation of non-dimensional

0.00883, when compared to the MD results.

84 Novel Nanomaterials - Synthesis and Applications

atoms becomes significant.

3.2. Vibration of SWCNTs with dimensionless frequency

frequency Ω through the following formula: Ω ¼ ωR

For the results generated so far, the nanotube in-plane rigidity has been taken to be Eh =278.25 Gpa�nm. However, there exist some inconsistencies concerning this quantity in the literature. The reported CNT in-plane stiffness is largely scattered, ranging from Eh =300 Gpa�nm to Eh = 400 Gpa�nm [42]. Both set of Figure 5 is presented to investigate the influence of the in-plane rigidity Eh variation on the dimensionless frequency of a (7, 7) armchair and (9, 0) Zigzag SWCNT with different boundary conditions likewise as clamped-clamped and clamped-free boundary conditions. These figures shows that for all the selected boundary conditions, dimensionless frequency calculated via shell model are sensitive to the nanotube in-plane rigidity Eh and also the larger the in-plane rigidity in-plane rigidity Eh, the higher the dimensionless frequency. The difference is more considerable for shorter length CNTs.

Previous study reveals that the bending rigidity of SWCNTs should be considered as an independent material parameter not linked to the representative thickness by the classic bending rigidity formula, i.e., <sup>D</sup> <sup>¼</sup> Eh<sup>3</sup> <sup>=</sup>12 1 � <sup>ν</sup><sup>2</sup> and the actual bending rigidity of SWCNTs is lesser than its classical counterpart [43, 44]. For shorter length-to-diameter ratio, the value of dimensionless frequencies for clamped-clamped at Eh = 300 Gpa�nm, Eh = 400 Gpa�nm is 0.04863, 0.05788, respectively which shows that a slight increase in frequency due to increase of in-plane rigidity Eh. Same trend is observed for dimensionless frequency. For the present shell model with in-plane rigidity Eh, the values of the C-F single-walled carbon nanotubes respectively, which are a little lower than those of corresponding CC SWCNTs with bending rigidity are plotted in Figure 5.

Figure 5. Variations of dimensionless frequencies Ω of CC and CF armchair and zigzag SWCNTs when Eh = 300 GPa.Nm and Eh = 425Gpa.Nm.

#### 3.4. Vibration with mass density per unit lateral area

Figure 6 is presented to investigate the influence of the mass density per unit lateral area variation on the dimensionless frequency of a (12, 12) armchair and (14, 0) zigzag SWCNT with boundary conditions: clamped-clamped and clamped-free. These figures shows that for all the selected boundary conditions, the frequency calculated via shell model are sensitive to the nanotube mass density and also the larger the mass density per unit lateral area ρh, lower the frequency. It is observed that applying the mass density per unit lateral area ρh to the present shell model, yields the slight decrease of the frequency. For shorter length-to-diameter ratio, the value of dimensionless frequencies at ρh =740.52 nm, 800.64 nm and 820.80 nm is 0.3667, 0.03342, 0.03128 respectively for clamped-clamped and 1.74285, 1.71298, 1.6001 respectively for clamped-free, which shows that decreases in frequency. For the present shell model, the values of length-to-diameter ratio for C-C SWCNTs, which are a little higher than those corresponding C-F SWCNTs values with mass density per unit lateral area are plotted in Figure 6.

4. Conclusions

Appendix 1

<sup>R</sup>2, <sup>L</sup><sup>31</sup> ¼ � <sup>v</sup>

Author details

References

<sup>2</sup>R<sup>2</sup> <sup>n</sup>2, <sup>L</sup><sup>12</sup> <sup>¼</sup> ikm <sup>1</sup>þ<sup>v</sup>

<sup>R</sup> ikm, <sup>L</sup><sup>32</sup> <sup>¼</sup> <sup>n</sup>

Muzamal Hussain\* and Muhammad Nawaz Naeem

\*Address all correspondence to: muzamal45@gmail.com

L<sup>11</sup> ¼ k 2 <sup>m</sup> <sup>þ</sup> <sup>1</sup>�<sup>v</sup>

<sup>L</sup><sup>23</sup> <sup>¼</sup> <sup>n</sup>

The vibration behavior of CF and CC SWCNTs are extensively investigated by present model compared with MD simulation. With properly chosen parameters, the present models can reproduce satisfactory frequencies that are in reasonable agreement with those results obtained by MD simulations and Timoshenko beam model. The effects of the length-to-diameter ratio for armchair and zigzag CNTs with in-plane rigidity, mass density per unit lateral area on the dimensionless frequencies are also examined with present models. It is found that the frequencies decreases smoothly when length-to-diameter ratio would increases and higher mode of vibration occurred when the frequencies are higher. For a clamped-free SWCNT, their exist an inverse proportionality which is observed between the resulting frequency and lengthto-diameter ratio. For clamped-clamped SWCNTs, the results took a similar trend but in this case frequency values are much higher. The results are obtained numerically for different boundary conditions and plotted in graphical forms. In the field of CNTs vibrations, wave propagation approach presents a good application. A better cylindrical shell model is needed to furnish more accurate prediction of the vibration frequencies of SWCNTs, such as the

Vibrational Behavior of Single-Walled Carbon Nanotubes Based on Donnell Shell Theory Using Wave Propagation…

nonlocal shell theory that incorporates the effect of small length scale effect.

<sup>2</sup><sup>R</sup> <sup>n</sup>, <sup>L</sup><sup>13</sup> <sup>¼</sup> ikm <sup>v</sup>

<sup>R</sup><sup>2</sup> þ

[1] Iijima S. Helical microtubules of graphitic carbon. Nature. 1991;345:56-58

International Journal of Solids and Structures. 2003;40:2487-2499

<sup>1</sup>�ν<sup>2</sup> ð Þ Eh D k<sup>4</sup>

Department of Mathematics, Government College University Faisalabad (GCUF), Pakistan

[2] Falvo MR, Clary GJ, Taylor RM II, Chi V, Brooks FP Jr, Washburn S, Superfine R. Bending and buckling of carbon nanotubes under large strain. Nature. 1997;389:532-534

[3] Li CY, Chou T. A structural mechanics approach for the analysis of carbon nanotubes.

<sup>R</sup>2, <sup>L</sup><sup>33</sup> <sup>¼</sup> <sup>1</sup>

<sup>R</sup>, <sup>L</sup><sup>21</sup> ¼ �<sup>n</sup> <sup>1</sup>þ<sup>v</sup>

<sup>m</sup> <sup>þ</sup> <sup>2</sup> <sup>1</sup>

<sup>2</sup><sup>R</sup> ikm, <sup>L</sup><sup>22</sup> <sup>¼</sup> <sup>1</sup>�<sup>v</sup>

<sup>R</sup><sup>2</sup> km<sup>2</sup>

.

<sup>2</sup> k 2 <sup>m</sup> <sup>þ</sup> <sup>n</sup><sup>2</sup> R2.

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87

<sup>n</sup><sup>2</sup> <sup>þ</sup> <sup>n</sup><sup>4</sup> R4

Figure 6. Variations of dimensionless frequencies of CC and CF armchair and zigzag SWCNTs when ρh ¼ 740:52nm, ρh ¼ 800:64nm and ρh ¼ 820:80nm.

#### 4. Conclusions

3.4. Vibration with mass density per unit lateral area

86 Novel Nanomaterials - Synthesis and Applications

Figure 6.

ρh ¼ 800:64nm and ρh ¼ 820:80nm.

Figure 6 is presented to investigate the influence of the mass density per unit lateral area variation on the dimensionless frequency of a (12, 12) armchair and (14, 0) zigzag SWCNT with boundary conditions: clamped-clamped and clamped-free. These figures shows that for all the selected boundary conditions, the frequency calculated via shell model are sensitive to the nanotube mass density and also the larger the mass density per unit lateral area ρh, lower the frequency. It is observed that applying the mass density per unit lateral area ρh to the present shell model, yields the slight decrease of the frequency. For shorter length-to-diameter ratio, the value of dimensionless frequencies at ρh =740.52 nm, 800.64 nm and 820.80 nm is 0.3667, 0.03342, 0.03128 respectively for clamped-clamped and 1.74285, 1.71298, 1.6001 respectively for clamped-free, which shows that decreases in frequency. For the present shell model, the values of length-to-diameter ratio for C-C SWCNTs, which are a little higher than those corresponding C-F SWCNTs values with mass density per unit lateral area are plotted in

Figure 6. Variations of dimensionless frequencies of CC and CF armchair and zigzag SWCNTs when ρh ¼ 740:52nm,

The vibration behavior of CF and CC SWCNTs are extensively investigated by present model compared with MD simulation. With properly chosen parameters, the present models can reproduce satisfactory frequencies that are in reasonable agreement with those results obtained by MD simulations and Timoshenko beam model. The effects of the length-to-diameter ratio for armchair and zigzag CNTs with in-plane rigidity, mass density per unit lateral area on the dimensionless frequencies are also examined with present models. It is found that the frequencies decreases smoothly when length-to-diameter ratio would increases and higher mode of vibration occurred when the frequencies are higher. For a clamped-free SWCNT, their exist an inverse proportionality which is observed between the resulting frequency and lengthto-diameter ratio. For clamped-clamped SWCNTs, the results took a similar trend but in this case frequency values are much higher. The results are obtained numerically for different boundary conditions and plotted in graphical forms. In the field of CNTs vibrations, wave propagation approach presents a good application. A better cylindrical shell model is needed to furnish more accurate prediction of the vibration frequencies of SWCNTs, such as the nonlocal shell theory that incorporates the effect of small length scale effect.

#### Appendix 1

 $L\_{11} = k^2$  $\eta\_m + \frac{1-v}{2\mathbb{R}^2}$  $n^2$ ,  $L\_{12} = \text{ik}\_m$  $\frac{1+v}{2\mathbb{R}}$  $\eta\_r$   $L\_{13} = \text{ik}\_m$  $\frac{v}{\mathbb{R}}$   $L\_{21} = -n\frac{1+v}{2\mathbb{R}}$  $\text{ik}\_m$   $L\_{22} = \frac{1-v}{2}$  $\frac{k}{\mathbb{R}}$   $\frac{1}{m} + \frac{n^2}{\mathbb{R}^2}$   $L\_{23} = \frac{v}{\mathbb{R}^2}$  $\text{ik}\_m$   $L\_{32} = \frac{1}{\mathbb{R}^2} + \frac{(1-v^2)}{\mathbb{R}^2}$  $D\left(k\_m + 2\frac{1}{\mathbb{R}^2}k\_m^2 n^2 + \frac{n^4}{\mathbb{R}^4}\right)$ .

#### Author details

Muzamal Hussain\* and Muhammad Nawaz Naeem

\*Address all correspondence to: muzamal45@gmail.com

Department of Mathematics, Government College University Faisalabad (GCUF), Pakistan

#### References


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**Chapter 6**

Provisional chapter

**Nanosynthesis Techniques of Silica-Coated**

DOI: 10.5772/intechopen.74097

Nanosynthesis Techniques of Silica-Coated

Core-shell nanomaterials are fast-emerging hybrid nanocomposites in area of nanotechnology, materials science and biochemistry, which are fast attracting research attention. Nanostructured nanomaterials are utilized in wide industry fields such as electronics, biopharmaceutical, biomedicine, optics and biocatalysis. Owing to the additional exterior shell-coating material, the primary core material's functionality, biocompatibility, chemical stability and colloidal dispersibility can be greatly enhanced. Silica, in particular, is found to be an excellent exterior shell-coating material, has been widely researched for the synthesis of core-shell nanocomposite materials. So far, there have been numerous publications devoted to silica-coating techniques using hydrophobic silanes, such as tetraethylorthosilicate (TEOS) or tetramethyl orthosilicate (TMOS), via the classic Stober method. Recently, there has been strong interest in the use of water-soluble silanes such as MPTMS (3-(mercaptopropyl)-trimethoxysilane), MPTES (3-(mercaptopropyl-triethoxysilane), MTMS (3-(methyltrimethoxysilane)) and sodium silicate for water-based silica-coating techniques have gained much attention, due to the fastgrowing need to focus on process simplicity, large-scale fabrication and environmentalfriendly synthesis techniques of silica-based core-shell nanomaterials. Hence, this chapter focuses on the recent development on silica-coating techniques for colloidal nanoparticles, particularly on water-based techniques and morphologies. In summary, we emphasize the importance of advanced nanomaterials in today's world and envisage there will be more breakthrough research on aqueous silica-coating techniques for silica-encapsulated core-shell

Keywords: silica-shell coatings, core-shell nanomaterials, aqueous one-pot synthesis

© 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 IntechOpen. 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

http://dx.doi.org/10.5772/intechopen.74097

**Nanostructures**

Nanostructures

Kwok Wei Shah

Kwok Wei Shah

Abstract

nanomaterials.


Provisional chapter

#### **Nanosynthesis Techniques of Silica-Coated Nanostructures** Nanosynthesis Techniques of Silica-Coated Nanostructures

DOI: 10.5772/intechopen.74097

#### Kwok Wei Shah Kwok Wei Shah

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Physical Review B. 2006;74:245413

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.74097

#### Abstract

Core-shell nanomaterials are fast-emerging hybrid nanocomposites in area of nanotechnology, materials science and biochemistry, which are fast attracting research attention. Nanostructured nanomaterials are utilized in wide industry fields such as electronics, biopharmaceutical, biomedicine, optics and biocatalysis. Owing to the additional exterior shell-coating material, the primary core material's functionality, biocompatibility, chemical stability and colloidal dispersibility can be greatly enhanced. Silica, in particular, is found to be an excellent exterior shell-coating material, has been widely researched for the synthesis of core-shell nanocomposite materials. So far, there have been numerous publications devoted to silica-coating techniques using hydrophobic silanes, such as tetraethylorthosilicate (TEOS) or tetramethyl orthosilicate (TMOS), via the classic Stober method. Recently, there has been strong interest in the use of water-soluble silanes such as MPTMS (3-(mercaptopropyl)-trimethoxysilane), MPTES (3-(mercaptopropyl-triethoxysilane), MTMS (3-(methyltrimethoxysilane)) and sodium silicate for water-based silica-coating techniques have gained much attention, due to the fastgrowing need to focus on process simplicity, large-scale fabrication and environmentalfriendly synthesis techniques of silica-based core-shell nanomaterials. Hence, this chapter focuses on the recent development on silica-coating techniques for colloidal nanoparticles, particularly on water-based techniques and morphologies. In summary, we emphasize the importance of advanced nanomaterials in today's world and envisage there will be more breakthrough research on aqueous silica-coating techniques for silica-encapsulated core-shell nanomaterials.

Keywords: silica-shell coatings, core-shell nanomaterials, aqueous one-pot synthesis

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons © 2018 The Author(s). Licensee IntechOpen. 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.

#### 1. Introduction

Core-shell nanomaterials are considered as highly functional hybrid nanomaterials with dual properties, originating from either core or shell materials. Core-shell nanomaterials have been widely researched due to its extraordinary ability to exhibit distinctive properties of both core and shell materials combined to deliver a wide spectrum of industry applications and requirements. Core-shell nanomaterials are commonly applied in different industries such as biomedical, bio-electronics, pharmaceutical applications, bio-catalysis, photoluminescence imaging, creating photonic crystals, etc. Especially for bioscience and medical research, the core-shell particles are mainly used for cell detection, bioimaging, targeted drug delivery, controlled drug release and bioengineering applications [1].

coupling agents, surfactants or polymers. The purpose of surface priming is to increase the compatibility of the core surface with silica, while providing particles with sufficient colloidal stability, so that colloids can be transferred into alcohol-based medium and classical Stober method [5] (using hydrophobic TEOS or TMOS in ethanol-water medium under alkaline conditions) is applicable for growing uniform silica shells on metal cores. A seminal publication on silica-coating gold nanoparticles was published by Liz-Marzán et al. [6]. Liz-Marzán et al. proposed silica-coating gold colloids in ethanol medium using a modified Stöber process. To achieve this, the authors surface-primed the nanoparticles with ATPMS before stabilizing the gold colloids by silica coating with a thin silica layer using sodium silicate in water medium, so that they can be stably transferred into ethanol medium. Since Liz-Marzán's pioneering work in 1996, many research studies have been carried out to simplify and shorten

Nanosynthesis Techniques of Silica-Coated Nanostructures

http://dx.doi.org/10.5772/intechopen.74097

93

Recently, there has been an increasing number of studies on aqueous synthesis techniques of silica-coating nanoparticles due to its process simplicity, large scalability and environmental friendliness. The key advantage of synthesizing silica in aqueous solution will be that the problem of poor colloidal stability of core nanoparticles in alcohol mediums, due to reduced surface charges, can be resolved and, thus, obviates any need to perform tedious surface stabilization procedures [7]. Essentially, any aqueous synthetic route will allow users to overcome a key challenge of colloidal instability faced by existing silica-coating techniques, which occurs when metal colloids are transferred into organic solvents such as ethanol or isopropanol

Water-based techniques enable scale-up production of silica-coated nanomaterials by harnessing the excellent colloidal stability of water-soluble nanoparticles in aqueous medium. The water solubility of silane precursors presents a highly attractive property to directly coat nanoparticles with silica in few simpler steps. Another advantage of an aqueous synthesis is that it promotes green chemistry in material synthesis through the use of inexpensive, nontoxic and ecologically friendly solvents such as water, in place of toxic solvents such as ethanol or isopropanol. By adopting a water-based silica-coating strategy, users can stably produce large quantities of core-shell nanostructures via simpler one-pot synthetic route. Essentially, by using water-soluble silane precursors, a water-based technique will allow fast and extensive silica coating of nanoparticle cores within a shorter time; without any need for transfer into alcoholic medium, dialysis or a change in reaction mixture, which are necessary for existing techniques such as Stöber process or reverse microemulsion techniques. A general synthetic route promotes green chemistry through sustainable material synthesis through the use of an inexpensive, non-toxic and ecologically friendly water medium, in place of organic toxic solvents such as ethanol or isopropanol. Furthermore, another important advantage of aqueous silica coating is that it is relatively neutral pH and lack of alcohol, thus being more suitable for encapsulating sensitive biomaterials such as enzymes and cells [8]. Aqueous silica-coating techniques are relatively straightforward processes that are less toxic, more cost-efficient and more environment-friendly. Hence, aqueous silica-coating techniques are not only less complex but also highly upscalable for commercial/industrial productions. These are the key reasons that motivated this chapter into summarizing the latest developments of new water-

the tedious multiple steps procedure [7].

for silica coating via classic Stober synthesis techniques.

based techniques for direct silica coating of core nanoparticles.

The benefits of an exterior shell-coating on a core particle are that surface reactivity and thermal stability can be greatly enhanced, while colloidal stability and dispersibility are vastly improved. Core-shell nanostructure itself can allow controlled release of center material, reduced usage of precious materials, hollow-core templating, etc. Hollow particles can be prepared using as a template by utilizing core-shell nanoparticles and after removing the core either by dissolution or calcination. These hollow-core nanoparticles are useful for various applications such as biocatalytic agent, biosupports, bioadsorbents, super light-weight structure, miniature vessels and thermal/electric insulative materials [1].

Silica is one of the most widely researched as exterior shell-coating material for encapsulating all sorts of core nanomaterials. Silica coating of nano-sized metals, semiconductors, magnetic and ceramic nanomaterials can enhance its large surface area, quantum confinement and photocatalytic, optical and magnetic properties. Silica is often chosen as the ideal shell material to encapsulate core nanomaterials in order to modify its exterior surface properties. Silica is also a favorite choice for core-shell nanostructures to enhance colloidal stability. The main advantages of silica as an excellent candidate for shell-encapsulation material predominantly lies in its exceptionally good colloidal stability, especially in waterbased media, easy controllability of the synthesis, chemical inertness, controlled porosity, high processability and optical transparency. Silica possesses other advantages such as (i) lower Van der Waals interactions compared to bare core particles (Hamaker constant is smaller for core-shell) and (ii) charged molecules that can be grafted onto silica shell at silica/water interfaces under alkaline conditions. Thus, silica shell can provide effective steric-hindrance and electrostatic protection on its core material, as well as function as excellent dispersant to provide electrostatically stable colloids. Furthermore, silica shell allows core nanoparticle to be biocompatible and easily bioconjugated with functional groups, which is necessary for biomedical, diagnostic and therapeutic applications (magnetic resonance imaging, MRI) [2].

However, the simple large-scale production and environmental-friendly fabrication of silicacoated core-shell nanomaterials currently remain a great challenge/bottleneck for practical commercialization [3]. The effective growth of homogenous silica shells on core nanoparticles faces problems of low chemical compatibility between both components [4]. Most existing silica-coating methods are dependent upon prior priming the surface of core material with coupling agents, surfactants or polymers. The purpose of surface priming is to increase the compatibility of the core surface with silica, while providing particles with sufficient colloidal stability, so that colloids can be transferred into alcohol-based medium and classical Stober method [5] (using hydrophobic TEOS or TMOS in ethanol-water medium under alkaline conditions) is applicable for growing uniform silica shells on metal cores. A seminal publication on silica-coating gold nanoparticles was published by Liz-Marzán et al. [6]. Liz-Marzán et al. proposed silica-coating gold colloids in ethanol medium using a modified Stöber process. To achieve this, the authors surface-primed the nanoparticles with ATPMS before stabilizing the gold colloids by silica coating with a thin silica layer using sodium silicate in water medium, so that they can be stably transferred into ethanol medium. Since Liz-Marzán's pioneering work in 1996, many research studies have been carried out to simplify and shorten the tedious multiple steps procedure [7].

1. Introduction

92 Novel Nanomaterials - Synthesis and Applications

drug release and bioengineering applications [1].

netic resonance imaging, MRI) [2].

ture, miniature vessels and thermal/electric insulative materials [1].

Core-shell nanomaterials are considered as highly functional hybrid nanomaterials with dual properties, originating from either core or shell materials. Core-shell nanomaterials have been widely researched due to its extraordinary ability to exhibit distinctive properties of both core and shell materials combined to deliver a wide spectrum of industry applications and requirements. Core-shell nanomaterials are commonly applied in different industries such as biomedical, bio-electronics, pharmaceutical applications, bio-catalysis, photoluminescence imaging, creating photonic crystals, etc. Especially for bioscience and medical research, the core-shell particles are mainly used for cell detection, bioimaging, targeted drug delivery, controlled

The benefits of an exterior shell-coating on a core particle are that surface reactivity and thermal stability can be greatly enhanced, while colloidal stability and dispersibility are vastly improved. Core-shell nanostructure itself can allow controlled release of center material, reduced usage of precious materials, hollow-core templating, etc. Hollow particles can be prepared using as a template by utilizing core-shell nanoparticles and after removing the core either by dissolution or calcination. These hollow-core nanoparticles are useful for various applications such as biocatalytic agent, biosupports, bioadsorbents, super light-weight struc-

Silica is one of the most widely researched as exterior shell-coating material for encapsulating all sorts of core nanomaterials. Silica coating of nano-sized metals, semiconductors, magnetic and ceramic nanomaterials can enhance its large surface area, quantum confinement and photocatalytic, optical and magnetic properties. Silica is often chosen as the ideal shell material to encapsulate core nanomaterials in order to modify its exterior surface properties. Silica is also a favorite choice for core-shell nanostructures to enhance colloidal stability. The main advantages of silica as an excellent candidate for shell-encapsulation material predominantly lies in its exceptionally good colloidal stability, especially in waterbased media, easy controllability of the synthesis, chemical inertness, controlled porosity, high processability and optical transparency. Silica possesses other advantages such as (i) lower Van der Waals interactions compared to bare core particles (Hamaker constant is smaller for core-shell) and (ii) charged molecules that can be grafted onto silica shell at silica/water interfaces under alkaline conditions. Thus, silica shell can provide effective steric-hindrance and electrostatic protection on its core material, as well as function as excellent dispersant to provide electrostatically stable colloids. Furthermore, silica shell allows core nanoparticle to be biocompatible and easily bioconjugated with functional groups, which is necessary for biomedical, diagnostic and therapeutic applications (mag-

However, the simple large-scale production and environmental-friendly fabrication of silicacoated core-shell nanomaterials currently remain a great challenge/bottleneck for practical commercialization [3]. The effective growth of homogenous silica shells on core nanoparticles faces problems of low chemical compatibility between both components [4]. Most existing silica-coating methods are dependent upon prior priming the surface of core material with Recently, there has been an increasing number of studies on aqueous synthesis techniques of silica-coating nanoparticles due to its process simplicity, large scalability and environmental friendliness. The key advantage of synthesizing silica in aqueous solution will be that the problem of poor colloidal stability of core nanoparticles in alcohol mediums, due to reduced surface charges, can be resolved and, thus, obviates any need to perform tedious surface stabilization procedures [7]. Essentially, any aqueous synthetic route will allow users to overcome a key challenge of colloidal instability faced by existing silica-coating techniques, which occurs when metal colloids are transferred into organic solvents such as ethanol or isopropanol for silica coating via classic Stober synthesis techniques.

Water-based techniques enable scale-up production of silica-coated nanomaterials by harnessing the excellent colloidal stability of water-soluble nanoparticles in aqueous medium. The water solubility of silane precursors presents a highly attractive property to directly coat nanoparticles with silica in few simpler steps. Another advantage of an aqueous synthesis is that it promotes green chemistry in material synthesis through the use of inexpensive, nontoxic and ecologically friendly solvents such as water, in place of toxic solvents such as ethanol or isopropanol. By adopting a water-based silica-coating strategy, users can stably produce large quantities of core-shell nanostructures via simpler one-pot synthetic route. Essentially, by using water-soluble silane precursors, a water-based technique will allow fast and extensive silica coating of nanoparticle cores within a shorter time; without any need for transfer into alcoholic medium, dialysis or a change in reaction mixture, which are necessary for existing techniques such as Stöber process or reverse microemulsion techniques. A general synthetic route promotes green chemistry through sustainable material synthesis through the use of an inexpensive, non-toxic and ecologically friendly water medium, in place of organic toxic solvents such as ethanol or isopropanol. Furthermore, another important advantage of aqueous silica coating is that it is relatively neutral pH and lack of alcohol, thus being more suitable for encapsulating sensitive biomaterials such as enzymes and cells [8]. Aqueous silica-coating techniques are relatively straightforward processes that are less toxic, more cost-efficient and more environment-friendly. Hence, aqueous silica-coating techniques are not only less complex but also highly upscalable for commercial/industrial productions. These are the key reasons that motivated this chapter into summarizing the latest developments of new waterbased techniques for direct silica coating of core nanoparticles.


decomposition. For example, synthesizing silica coatings encapsulating gold nanoparticles with Raman-active dye molecules was proven to be useful as a surface-enhanced Raman scattering (SERS) nanoprobes. Inert and biocompatible silica-shell coatings are useful for bioconjugation of quantum dots (QDs). Silica core-shell inhibits agglomeration of nanoparticles, hinders foreign biospecies from contaminating the nanoparticles' surface and aids to enhance the high photoluminescence for biosensing. Surface functionalisation of the silica outer shell with any various types of functional groups, such as carboxylate COOH, amine NH2, phosphate, thiol SH and poly(ethyleneglycol) (PEG) groups, enables enhanced control in

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A primary work on silica-coating gold nanoparticles was published by Liz-Marzán et al. [6] in 1996. The authors proposed silica-coating gold colloids in ethanol-based medium using a modified Stober process [5]. First, the gold colloidal surface must be activated with a surfacecoupling agent, 3-aminopropyltrimethoxysilane (APTMS), to render the gold nanoparticle surface vitreophillic (that is receptive toward silica monomers or oligomers) before proceeding to coat them with water-soluble sodium silicate due to the low affinity between gold nanoparticles and sodium silicate. After APTMS treatment, the vitreophilic Au nanoparticles are coated with a thin inhomogenous silica shell (thickness < 2 nm) using sodium silicate in water at a low pH to stabilize the gold colloids, so that they can be stably transferred into alcoholic medium (Figure 1). However, the critical step of silanization with sodium silicate is poorly reproducible, which requires an extended time duration, typically from a few days to weeks, before a sufficient layer silica shell can be formed to stabilize the nanoparticles in the alcohol solution [10]. Using Stober process as the last step, Liz-Marzán coated a thicker homogenous silica layer (thickness > 5 nm) in ethanol-water mixture using silane precursor

Zhou et al. in 2005 [11] synthesized, based on a simple chemical precipitation, nano-sized CdSe quantum dots at normal temperature in slightly alkaline conditions. Also, silica encapsulation onto these QDs via a slow polymerization of SiO2 using sodium silicate is done by adding ethanol. Cadmium selenide (CdSe) quantum dots were thinly encapsulated with a 1–2 nm thin silica in an aqueous sodium silicate at pH 10.5 over 5 days of vigorous stirring (Figure 2). However, despite the absence of aggregates proving improvement to dispersity, the coated particles were hardly monodispersed and without a distinctive core-shell

Figure 1. (Left) Schematic of aqueous synthesis procedure used by Liz-Marzán to coat thin silica shell on Au nanoparticles. (Right) TEM images of 15 nm-sized gold nanoparticles coated with 2 nm silica shells after addition of sodium silicate [6].

TEOS and the silanol groups provided by sodium silicate as anchor points.

bioconjugation techniques [9].

Table 1. Pros & cons of aqueous and non-aqueous silica-coating routes.

Since water-based silica-coating of core-shell nanoparticles is a fast broadening field that is presently seeing remarkable developments, we summarized a collection of up-to-date work related to a variety of water-based silica core-shell nanomaterials and their synthesis techniques. In this chapter, four main types of aqueous synthesis routes related to water-soluble silane precursors have been widely used in recent publications, namely sodium silicate, MPTMS (3-(mercaptopropyl)-trimethoxysilane), MPTES (3-(mercaptopropyl-triethoxysilane) and MTMS (3-(methyltrimethoxysilane)). These silane precursors possess high water solubility. For example, the alkoxy group of MPTMS has been replaced with a sufficiently polar group. The substitution of the methoxy group in water-insoluble tetramethyl orthosilicate (TMOS) with a mercaptopropyl group ('MP') results in MPTMS that can be hydrolyzed in water prior to condensation to form core-shell structures at the nanoscale.

Table 1 summarizes the pros and cons of aqueous silica coating (using water-soluble silanes, namely sodium silicate, MPTMS, MPTES and MTMS) and solvent-based silica coating (using water-insoluble silanes, such as TEOS), which will be discussed in subsequent sections.

#### 2. Sodium silicate-assisted silica coating

Silanization of diverse metal-core and semiconductor-core nanomaterial into core-shell systems has gathered much research attention due to enhancement of surface properties. Silicashell encapsulation was useful in inhibiting particle growth, agglomeration and photo-induced decomposition. For example, synthesizing silica coatings encapsulating gold nanoparticles with Raman-active dye molecules was proven to be useful as a surface-enhanced Raman scattering (SERS) nanoprobes. Inert and biocompatible silica-shell coatings are useful for bioconjugation of quantum dots (QDs). Silica core-shell inhibits agglomeration of nanoparticles, hinders foreign biospecies from contaminating the nanoparticles' surface and aids to enhance the high photoluminescence for biosensing. Surface functionalisation of the silica outer shell with any various types of functional groups, such as carboxylate COOH, amine NH2, phosphate, thiol SH and poly(ethyleneglycol) (PEG) groups, enables enhanced control in bioconjugation techniques [9].

A primary work on silica-coating gold nanoparticles was published by Liz-Marzán et al. [6] in 1996. The authors proposed silica-coating gold colloids in ethanol-based medium using a modified Stober process [5]. First, the gold colloidal surface must be activated with a surfacecoupling agent, 3-aminopropyltrimethoxysilane (APTMS), to render the gold nanoparticle surface vitreophillic (that is receptive toward silica monomers or oligomers) before proceeding to coat them with water-soluble sodium silicate due to the low affinity between gold nanoparticles and sodium silicate. After APTMS treatment, the vitreophilic Au nanoparticles are coated with a thin inhomogenous silica shell (thickness < 2 nm) using sodium silicate in water at a low pH to stabilize the gold colloids, so that they can be stably transferred into alcoholic medium (Figure 1). However, the critical step of silanization with sodium silicate is poorly reproducible, which requires an extended time duration, typically from a few days to weeks, before a sufficient layer silica shell can be formed to stabilize the nanoparticles in the alcohol solution [10]. Using Stober process as the last step, Liz-Marzán coated a thicker homogenous silica layer (thickness > 5 nm) in ethanol-water mixture using silane precursor TEOS and the silanol groups provided by sodium silicate as anchor points.

Since water-based silica-coating of core-shell nanoparticles is a fast broadening field that is presently seeing remarkable developments, we summarized a collection of up-to-date work related to a variety of water-based silica core-shell nanomaterials and their synthesis techniques. In this chapter, four main types of aqueous synthesis routes related to water-soluble silane precursors have been widely used in recent publications, namely sodium silicate, MPTMS (3-(mercaptopropyl)-trimethoxysilane), MPTES (3-(mercaptopropyl-triethoxysilane) and MTMS (3-(methyltrimethoxysilane)). These silane precursors possess high water solubility. For example, the alkoxy group of MPTMS has been replaced with a sufficiently polar group. The substitution of the methoxy group in water-insoluble tetramethyl orthosilicate (TMOS) with a mercaptopropyl group ('MP') results in MPTMS that can be hydrolyzed in

1. May still require a small amount of acidic and/or basic

1. Complex reactions, multiple precursors, solvents, controls 2. Multiple reactants increase costs, time and environmental

3. Alcoholic and extreme pH conditions are cytotoxic 4. Processes to remove alcohol require evaporation and buffer solutions, increasing time, cost and complexity

2. May require new surface functional groups. 3. Relatively new field, hence limited literature available. Generally applicable to hydrophilic core-materials only

catalysts

pollution, limited scale

5. Need surface functionalisation

Table 1 summarizes the pros and cons of aqueous silica coating (using water-soluble silanes, namely sodium silicate, MPTMS, MPTES and MTMS) and solvent-based silica coating (using water-insoluble silanes, such as TEOS), which will be discussed in subsequent sections.

Silanization of diverse metal-core and semiconductor-core nanomaterial into core-shell systems has gathered much research attention due to enhancement of surface properties. Silicashell encapsulation was useful in inhibiting particle growth, agglomeration and photo-induced

water prior to condensation to form core-shell structures at the nanoscale.

Pros Cons

1. Facile one-step synthesis 2. Precisely controllable shell

3. Shell growth is generally faster than non-aqueous routes 4. Functional groups may be intrinsically incorporated

environmentally friendly 6. Most cost-effective and upscalable technique

Table 1. Pros & cons of aqueous and non-aqueous silica-coating routes.

1. Shell thickness adjustable through reactant conditions 2. Extensively researched and mature technology 3. Mild, ambient temperature, low-pressure conditions 4. Nanometric resolution control of core/shell morphology 5. Structures and morphologies easily reproducible and tunable 6. Generally fast hydrolysis and condensation of silica

thickness

94 Novel Nanomaterials - Synthesis and Applications

5. Mild and less toxic,

"Aqueousbased" silica-coating techniques

"Solventbased" silica-coating techniques

2. Sodium silicate-assisted silica coating

Zhou et al. in 2005 [11] synthesized, based on a simple chemical precipitation, nano-sized CdSe quantum dots at normal temperature in slightly alkaline conditions. Also, silica encapsulation onto these QDs via a slow polymerization of SiO2 using sodium silicate is done by adding ethanol. Cadmium selenide (CdSe) quantum dots were thinly encapsulated with a 1–2 nm thin silica in an aqueous sodium silicate at pH 10.5 over 5 days of vigorous stirring (Figure 2). However, despite the absence of aggregates proving improvement to dispersity, the coated particles were hardly monodispersed and without a distinctive core-shell

Figure 1. (Left) Schematic of aqueous synthesis procedure used by Liz-Marzán to coat thin silica shell on Au nanoparticles. (Right) TEM images of 15 nm-sized gold nanoparticles coated with 2 nm silica shells after addition of sodium silicate [6].

Kobayashi et al. in 2007 [13] carried out silica encapsulation of nano-sized copper prepared from water-based solutions of copper salts. The silica shells chemically stabilized Cu colloids. Hydrazine and citric acid were used to reduce Cu salts and CTAB surfactants to stabilize 50 nm Cu metallic core particles. Then, silane-coupling agent, ATPMS, is added for surface functionalisation, and then Cu nanoparticles were coated with sodium silicate. Cu@SiO2 coreshell nanoparticles are synthesized with ~10 nm silica layer. The mixed solution was left for 15 min for APTMS to attach on the surface of the core particles. Then, sodium silicate was added to the colloid, followed by tuning to pH 10 using CTAB resin for cation exchange. The synthesis time for the silica shell was 24 h. The key problem for using Cu-based particles is their fast oxidation in air environment. This arises from their instability toward oxygen to form CuO. To improve stability and reduce oxidation of copper nanoparticles, surfactant capping surrounding the copper particles inhibits oxygen molecules found inside the water-based medium from reacting with the copper. Previously, Kobayashi performed chemical stabilization by silica protection of cobalt nanoparticles prepared through reducing of Co salts in waterbased medium. The silica shell forms a homogenous protection layer habiting O2 molecules from reacting with the metallic cobalt nanoparticles. The nanoparticles were stable and chemically unchanged under high temperature in O2 environment. This method has benefits such as (i) extraordinary colloidal stable dispersions in water, (ii) facile surface functionalisation to synthesize colloids for non-aqueous mediums and (iii) facile control of particle-to-particle

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Hu et al. in 2009 [14] synthesized Fe3O4@SiO2 magnetic composite nanospheres for fast and recyclable removal of lead and mercury from water (Figure 4). The Fe3O4@SiO2 nanospheres were synthesized with the dissolution of sodium silicate in deionized water, into which Fe3O4 was added. HCL was added as a catalyst to adjust pH to 6.0 and the mixture stirred for 3 h at 80C. A thin layer of silica coating visibly encapsulates the iron oxide spheres; the deposition

Wang et al. in 2010 [15] demonstrated a slightly different route starting with a suspension of Fe3O4 in deionized water at 80C under a nitrogen flow, the pH was adjusted to 6.0 with

Figure 4. (Top row) Schematic of aqueous synthesis procedure used by Hu et al. to coat thin silica shell on Fe3O4 nanoparticles using sodium silicate. (Bottom row) SEM images of bare Fe3O4 microspheres (A, B) and (C) TEM images

interactions.

of silica was also confirmed through XRD.

of Fe3O4@SiO2 composite microparticles [14].

Figure 2. (Left) Schematic of aqueous synthesis procedure used by Zhou et al. to coat thin silica shell on CdSe nanoparticles using sodium silicate. (Right) TEM of CdSe NPs before (A) and after (B) encapsulation [11].

Figure 3. (Left) Schematic of 3-mercaptopropyltrimethoxysilane is binded via the sulfur to the nanowire surface with Ag, leaving three free –OCH2 methoxy groups facing the water medium. Addition of sodium silicate produces silica, SiO2. (Right) TEM images of Ag nanowires and silica-encapsulted Ag nanowires. Scale bar is 500 nm (g) and 100 nm (h). For thick layers of silica shells, the coating becomes inconsistent (h) [12].

structure. The authors attributed the success of silica-coating to the mercapto-groups present in the MPTMS, which are used to bridge the low affinity between CdSe core and silica shell. This principle would be repeated by Wolcott et al. [9] in the following year in their reported attempt to coat cadmium telluride (CdTe) quantum dots as they primed the CdTe nanoparticles with MPTMS prior to gradual silica growth under stirring over 72 h in aqueous sodium silicate solution.

Hunyadi et al. in 2006 [12] synthesized silver nanowires coated with a silica shell. Simona uses an indirect coating method that involved the use of a bifunctional linker, (HS(CH2)3Si(OCH3)3, 3-mercaptopropyltrimethoxysilane), so that the thiol SH groups bind Ag to the core surface [12]. Simona's approach is essentially seedless, surfactantless approach. Different volumes of sodium silicate Na2O(SiO2)3 were added in alkaline pH and left overnight to allow formation of well-controlled layers of silica. The silica thickness was precisely controlled from 10 to 150 nm. The uniform silica coating over the whole surface of the nanowire is subjected to initial sodium silicate concentrations. Simona also found that MPTMS is critical to make a uniform layer of silver. When the MPTMS dosage exceeded 8.08 mM, the silica shell became undulating and inhomogenous (Figure 3). The reference experiment to synthesize SiO2 on silver nanowires without the silane primer proved that the maximum shell thickness achieved was 5–10 nm only, regardless of the starting dosage of the silicate precursor. This outcome proves that silane performs a critical "surface-primer" function at interface of core-shell materials in the encapsulation process as a molecular binder.

Kobayashi et al. in 2007 [13] carried out silica encapsulation of nano-sized copper prepared from water-based solutions of copper salts. The silica shells chemically stabilized Cu colloids. Hydrazine and citric acid were used to reduce Cu salts and CTAB surfactants to stabilize 50 nm Cu metallic core particles. Then, silane-coupling agent, ATPMS, is added for surface functionalisation, and then Cu nanoparticles were coated with sodium silicate. Cu@SiO2 coreshell nanoparticles are synthesized with ~10 nm silica layer. The mixed solution was left for 15 min for APTMS to attach on the surface of the core particles. Then, sodium silicate was added to the colloid, followed by tuning to pH 10 using CTAB resin for cation exchange. The synthesis time for the silica shell was 24 h. The key problem for using Cu-based particles is their fast oxidation in air environment. This arises from their instability toward oxygen to form CuO. To improve stability and reduce oxidation of copper nanoparticles, surfactant capping surrounding the copper particles inhibits oxygen molecules found inside the water-based medium from reacting with the copper. Previously, Kobayashi performed chemical stabilization by silica protection of cobalt nanoparticles prepared through reducing of Co salts in waterbased medium. The silica shell forms a homogenous protection layer habiting O2 molecules from reacting with the metallic cobalt nanoparticles. The nanoparticles were stable and chemically unchanged under high temperature in O2 environment. This method has benefits such as (i) extraordinary colloidal stable dispersions in water, (ii) facile surface functionalisation to synthesize colloids for non-aqueous mediums and (iii) facile control of particle-to-particle interactions.

Hu et al. in 2009 [14] synthesized Fe3O4@SiO2 magnetic composite nanospheres for fast and recyclable removal of lead and mercury from water (Figure 4). The Fe3O4@SiO2 nanospheres were synthesized with the dissolution of sodium silicate in deionized water, into which Fe3O4 was added. HCL was added as a catalyst to adjust pH to 6.0 and the mixture stirred for 3 h at 80C. A thin layer of silica coating visibly encapsulates the iron oxide spheres; the deposition of silica was also confirmed through XRD.

structure. The authors attributed the success of silica-coating to the mercapto-groups present in the MPTMS, which are used to bridge the low affinity between CdSe core and silica shell. This principle would be repeated by Wolcott et al. [9] in the following year in their reported attempt to coat cadmium telluride (CdTe) quantum dots as they primed the CdTe nanoparticles with MPTMS prior to gradual silica growth under stirring over 72 h in aque-

Figure 3. (Left) Schematic of 3-mercaptopropyltrimethoxysilane is binded via the sulfur to the nanowire surface with Ag, leaving three free –OCH2 methoxy groups facing the water medium. Addition of sodium silicate produces silica, SiO2. (Right) TEM images of Ag nanowires and silica-encapsulted Ag nanowires. Scale bar is 500 nm (g) and 100 nm (h). For

Figure 2. (Left) Schematic of aqueous synthesis procedure used by Zhou et al. to coat thin silica shell on CdSe

nanoparticles using sodium silicate. (Right) TEM of CdSe NPs before (A) and after (B) encapsulation [11].

Hunyadi et al. in 2006 [12] synthesized silver nanowires coated with a silica shell. Simona uses an indirect coating method that involved the use of a bifunctional linker, (HS(CH2)3Si(OCH3)3, 3-mercaptopropyltrimethoxysilane), so that the thiol SH groups bind Ag to the core surface [12]. Simona's approach is essentially seedless, surfactantless approach. Different volumes of sodium silicate Na2O(SiO2)3 were added in alkaline pH and left overnight to allow formation of well-controlled layers of silica. The silica thickness was precisely controlled from 10 to 150 nm. The uniform silica coating over the whole surface of the nanowire is subjected to initial sodium silicate concentrations. Simona also found that MPTMS is critical to make a uniform layer of silver. When the MPTMS dosage exceeded 8.08 mM, the silica shell became undulating and inhomogenous (Figure 3). The reference experiment to synthesize SiO2 on silver nanowires without the silane primer proved that the maximum shell thickness achieved was 5–10 nm only, regardless of the starting dosage of the silicate precursor. This outcome proves that silane performs a critical "surface-primer" function at interface of core-shell mate-

ous sodium silicate solution.

96 Novel Nanomaterials - Synthesis and Applications

rials in the encapsulation process as a molecular binder.

thick layers of silica shells, the coating becomes inconsistent (h) [12].

Wang et al. in 2010 [15] demonstrated a slightly different route starting with a suspension of Fe3O4 in deionized water at 80C under a nitrogen flow, the pH was adjusted to 6.0 with

Figure 4. (Top row) Schematic of aqueous synthesis procedure used by Hu et al. to coat thin silica shell on Fe3O4 nanoparticles using sodium silicate. (Bottom row) SEM images of bare Fe3O4 microspheres (A, B) and (C) TEM images of Fe3O4@SiO2 composite microparticles [14].

hydrochloric acid. Sodium silicate precursor was then added dropwise under strong stirring. The coating process was carried out for 5 h and the Fe3O4@SiO2 nano particles collected via magnetic separation (Figure 5). In both Hu's and Wang's experiments, the silica shell was thin (<2 nm) and only just barely discernible under high-magnification TEM.

the authors noted that excessively high pH conditions (>11) resulted in a thin silica layer with pinholes and low conditions (pH < 8) resulted in thicker silica shells. The shell will be dissolved by sodium hydroxide interfering with the silica growth if the pH is too high, e.g. pH greater than 11. This results in holes of nano to micron sizes within the thin silica shells. The silica shell grows too fast and thickens quickly for a low pH of 8 for the sodium silicate medium (Figure 6). Therefore, plasmonic field enhancement from the metal surface cannot extend beyond the shell,

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Shin et al. 2015 [17] demonstrated that a single layer of self-assembled core metal particles with silica shells can produce 1~2 nm empty nano-gaps by assembling silica-gold nanoparticles (Au@SiO2) and chemical etching on different substrates. Colloidal silica shell gold-core nanoparticles were prepared by adding 5% v/v (3-aminopropyl)trimethoxysilane to gold nanoparticle seed solution. The gold colloid reacted with aqueous APTMS solution under vigorous stirring for 30 min at room temperature. After mixing, sodium silicate aqueous solution was added. A 5 nm silica layer is grown by placing in 90C hot bath for 1 h under fast stirring and allowing the solution to cool to normal temperature. By reacting for 0.5 h in hot bath, the silica layer is grown

Silanization methods have also been developed for surface stabilization of core nanoparticles using sodium silicate, attaching artificial resin polymers (e.g., polyvinylpyrrolidone) or substituting the depleted sodium citrate on the surface of core nanoparticle adding extra fresh sodium citrate to metal cores. After surface stabilization steps, traditional Stober method is used for thick silica coating of the core. This requires conventional silane precursors (TMOS, TEOS) to be dissolved in ethanol-water mixture, which facilitates its hydrolysis and condensation. As these hydrophobic precursors are insoluble in aqueous medium and cannot undergo hydrolysis, large amount of alchohol is used to dissolve hydrophobic silanes prior to silica condensation process. Besides using such hydrophobic silane precursors, there arises strong interest to discover watersoluble silane precursors for silica encapsulation in aqueous-based medium, which allows silicacoating process to be much simpler, cost-efficient and environmentally friendly for various

Recently, several research groups have reported the use of 3-mercaptopropyl-trimethoxysilane (MPTMS) as a water-soluble silane precursor to synthesize thick silica shells in fully aqueous medium. Previously, Niu [19], Nakamura [20], Lee [21], Shah [22] and Shang [23] successfully silica-coated metal nanoparticles, quantum dots and iron-oxide nanoparticles using waterbased MPTMS in aqueous solution under mild conditions. Using MPTMS as a water-soluble silane precursor, direct alcohol-free silica coating in water-based medium becomes possible. This straightforward, one-pot method uses MPTMS as (i) surface primer, which renders particle-surface vitreophilic, (ii) a silane-precursor, which provides a thick silica shell layer >10 nm for particle stabilzation, (iii) surface modifier, which provides Thiol SH groups and (iv) a negative surface charge for dispersibility. Furthermore, speed, thickness and uniformity

and no signal is acquired.

to thickness of 2.5 nm only.

applications [18].

3. MPTMS-assisted silica coating

Li et al. in 2013 [16] proposed a shell-isolated nanoparticle-enhanced Raman spectroscopy (SHINERS) technique using core-shell structured gold@silica nanoparticles in aqueous solution with sodium silicate. The authors synthesized shell-isolated nanoparticles (SHINs) of sphere (55 and 120 nm diameter) finding that a 90C reaction temperature accelerates the formation of silica shell on the gold nanoparticles. Gold nanoparticles were surface-primed using APTMS and coated with a 4 and 1 nm layer of silica using sodium silicate. Additionally,

Figure 5. (Top row) Schematic of aqueous synthesis procedure used by Wang et al. to coat thin silica shell on Fe3O4 nanoparticles using sodium silicate. (Bottom row) TEM image showing (A) Fe3O4, (B) Fe3O4@SiO2 particles and (C) Fe3O4@SiO2–NH2 [15].

Figure 6. (Left) Schematic of aqueous synthesis procedure used by Li et al. to coat thin silica shell on Au nanoparticles using sodium silicate. (Right) TEM photos of gold@SiO2 nanoparticles. (a–d) Various pH levels are used to grow silica shells and using APTMS as surface primer. High pH of 11 and low pH of 8 conditions with 1 h. heating using sodium silicate solution were used to prepare the SHINs (a and b). (c and d) gold@SiO2 NPs were synthesized of ~10.2 for 0.5 and 1 h, respectively, at pH 10 [16].

the authors noted that excessively high pH conditions (>11) resulted in a thin silica layer with pinholes and low conditions (pH < 8) resulted in thicker silica shells. The shell will be dissolved by sodium hydroxide interfering with the silica growth if the pH is too high, e.g. pH greater than 11. This results in holes of nano to micron sizes within the thin silica shells. The silica shell grows too fast and thickens quickly for a low pH of 8 for the sodium silicate medium (Figure 6). Therefore, plasmonic field enhancement from the metal surface cannot extend beyond the shell, and no signal is acquired.

Shin et al. 2015 [17] demonstrated that a single layer of self-assembled core metal particles with silica shells can produce 1~2 nm empty nano-gaps by assembling silica-gold nanoparticles (Au@SiO2) and chemical etching on different substrates. Colloidal silica shell gold-core nanoparticles were prepared by adding 5% v/v (3-aminopropyl)trimethoxysilane to gold nanoparticle seed solution. The gold colloid reacted with aqueous APTMS solution under vigorous stirring for 30 min at room temperature. After mixing, sodium silicate aqueous solution was added. A 5 nm silica layer is grown by placing in 90C hot bath for 1 h under fast stirring and allowing the solution to cool to normal temperature. By reacting for 0.5 h in hot bath, the silica layer is grown to thickness of 2.5 nm only.

#### 3. MPTMS-assisted silica coating

hydrochloric acid. Sodium silicate precursor was then added dropwise under strong stirring. The coating process was carried out for 5 h and the Fe3O4@SiO2 nano particles collected via magnetic separation (Figure 5). In both Hu's and Wang's experiments, the silica shell was thin

Li et al. in 2013 [16] proposed a shell-isolated nanoparticle-enhanced Raman spectroscopy (SHINERS) technique using core-shell structured gold@silica nanoparticles in aqueous solution with sodium silicate. The authors synthesized shell-isolated nanoparticles (SHINs) of sphere (55 and 120 nm diameter) finding that a 90C reaction temperature accelerates the formation of silica shell on the gold nanoparticles. Gold nanoparticles were surface-primed using APTMS and coated with a 4 and 1 nm layer of silica using sodium silicate. Additionally,

Figure 5. (Top row) Schematic of aqueous synthesis procedure used by Wang et al. to coat thin silica shell on Fe3O4 nanoparticles using sodium silicate. (Bottom row) TEM image showing (A) Fe3O4, (B) Fe3O4@SiO2 particles and

Figure 6. (Left) Schematic of aqueous synthesis procedure used by Li et al. to coat thin silica shell on Au nanoparticles using sodium silicate. (Right) TEM photos of gold@SiO2 nanoparticles. (a–d) Various pH levels are used to grow silica shells and using APTMS as surface primer. High pH of 11 and low pH of 8 conditions with 1 h. heating using sodium silicate solution were used to prepare the SHINs (a and b). (c and d) gold@SiO2 NPs were synthesized of ~10.2 for 0.5 and

(C) Fe3O4@SiO2–NH2 [15].

1 h, respectively, at pH 10 [16].

(<2 nm) and only just barely discernible under high-magnification TEM.

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Silanization methods have also been developed for surface stabilization of core nanoparticles using sodium silicate, attaching artificial resin polymers (e.g., polyvinylpyrrolidone) or substituting the depleted sodium citrate on the surface of core nanoparticle adding extra fresh sodium citrate to metal cores. After surface stabilization steps, traditional Stober method is used for thick silica coating of the core. This requires conventional silane precursors (TMOS, TEOS) to be dissolved in ethanol-water mixture, which facilitates its hydrolysis and condensation. As these hydrophobic precursors are insoluble in aqueous medium and cannot undergo hydrolysis, large amount of alchohol is used to dissolve hydrophobic silanes prior to silica condensation process. Besides using such hydrophobic silane precursors, there arises strong interest to discover watersoluble silane precursors for silica encapsulation in aqueous-based medium, which allows silicacoating process to be much simpler, cost-efficient and environmentally friendly for various applications [18].

Recently, several research groups have reported the use of 3-mercaptopropyl-trimethoxysilane (MPTMS) as a water-soluble silane precursor to synthesize thick silica shells in fully aqueous medium. Previously, Niu [19], Nakamura [20], Lee [21], Shah [22] and Shang [23] successfully silica-coated metal nanoparticles, quantum dots and iron-oxide nanoparticles using waterbased MPTMS in aqueous solution under mild conditions. Using MPTMS as a water-soluble silane precursor, direct alcohol-free silica coating in water-based medium becomes possible. This straightforward, one-pot method uses MPTMS as (i) surface primer, which renders particle-surface vitreophilic, (ii) a silane-precursor, which provides a thick silica shell layer >10 nm for particle stabilzation, (iii) surface modifier, which provides Thiol SH groups and (iv) a negative surface charge for dispersibility. Furthermore, speed, thickness and uniformity can be easily controlled through the concentration of precursor, reaction times and reaction temperatures, making aqueous routes far less complex than classic Stober method.

rhodamine B were added. After reaction, the solution was centrifuged to discard unreacted contents. The aqueous silica encapsulation using Nakamura's ultra-fast technique took only 3 h, as opposed to using conventional reverse emulsion technique, which requires 1 day or using water-based sodium silicate, which requires 3 days. In addition, the presence of intrinsic thiol SH groups allows conjugation with thiol-reactive dyes or biomolecules, without the need

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Lee et al. in 2011 [21] reported that for in vivo applications, such as biological imaging and medicine, nanoparticles must be safe in cell toxicity. Gold nanorods are often surface-capped and stabilized with CTAB cetyltrimethylammonium bromide, which will disrupt cell biology and cell growth as a cytotoxic surfactant. Silica coating and surface encapsulation of single nanoparticles can provide a safer substitute to surface CTAB for both decreasing cytotoxicity and improving colloidal stability. Lee synthesized silica-coated gold nanorods in aqueous medium. In his synthesis technique, Au nanorods were mixed to a solution of (3-aminopropyldimethylethoxysilane) and MPTMS for 12 h under stirring after which NH4OH was added to form silica shell thickness < 2 nm. The authors found that the amount of amino-silane

Shah et al. in 2014 [22] reported a novel direct silica coating in water-based silica-coating routes, which is alcohol-free one-step method using "pre-hydrolysed" MPTMS making silica coating far less complex than classic Stober method. This technique presents a facile, scalable, environmentally friendly process for thick coating of metal particles with a thick (5–30 nm) silica shell. The success of the process is using "pre-hydrolysed" MPTMS as silane precursor. Silica coating could be directly attributed to the strong and direct metal-thiol bond. Briefly, metal nanoparticles were washed and re-dispersed in deionized water before being combined into a pre-hydrolysed MPTMS in DI water. Ammonia was added as catalyst. While growth was easily scalable, the growth rate of the silica shell is controllable to 25, 11.7 and 10 nm/h. during condensation. The process slows down owing to eventual consumption of hydrolysed MPTMS. Different metal NPs (i.e., gold, silver, platinum) are successfully encapsulated with a silica layer under a completely alcohol-free environment, whose shell thickness can be easily and precisely synthesized by changing the growth duration (Figure 9). This novel water-based procedure is used for the fabrication of SERS-enhanced silver@SiO2 nanoparticles without the

Figure 8. (Left) (a) Schematic of aqueous synthesis procedure used by Lee et al. to coat thin silica shell on Au nanorods using sodium silicate. (Right) Images of AuNRs sized 50 nm, (b) SEM image of core-shell Au nanorods and (c) TEM image

for additional surface functionalization steps.

of silica shell (<2 nm) AuNRs [21].

affected the growth of silica layer surrounding gold-NRs (Figure 8).

Niu et al. in 2010 [19] reported a facile route to fabricate homogenous, monodisperse, accurately size-tunable, SH-functionalized and magnetic silica-shell nanocomposite nanospheres (SHSSCNs) with sizes less than 100 nm. These nanocomposite nanospheres are synthesized based on the self-assembly of magnetic nanoparticles and a surface-priming copolymer polystyrene-block-poly(acrylic acid) (PS-b-PAA) in aqueous medium, followed by an easy silanization reaction with a silane precursor MPTMS to grow a silica shell on the interface of magnetite micelles at ambient temperature. (Step 1) Hydrophobic magnetite (Fe3O4) nanocores with a dimension of 5 nm were added to form magnetite-based micelles through their selfassembly in water with an amphiphilic block copolymer, PS-b-PAA. After filtration with deionized water, (Step 2) MPTMS was used to grow a hybrid silica layer by hydrolyzing and condensing the silane onto surface of the magnetic micelles under alkaline conditions. At the same time, SH groups were functionalised onto the surface of the nanocomposite particle in situ by using MPTMS silane as precursor, which is obviously distinct from traditional silicaencapsulation techniques reported to date. This is possible due to the fact that the core particles are encapsulated by a layer of polymeric shell, which allows the MPTMS silane to attach easily and directly onto the polymer shell's abundant surface hydroxyl groups. During the silicaencapsulation procedure, MPTMS performs as the silane precursor, thiol-functional linker and surface-capping stabilizers for the magnetite micelles and also for SH-modification of the nanoparticles in the silanization step.

Nakamura et al. in 2011 [20] used MPTMS for surface modification of quantum dots and synthesis of the silica shell to grow a layer of organosilica dense shell. MPTMS, quantum dot QD605 and NH4OH were added and then heated for 3 h to grow the silica layer at 100C. Extra steps like filtration or exchanging the reacting medium or adding alcohol were required. The growth of the silica shell was easily done in 3 h (Figure 7). The procedure is fast and easy, and the thiol groups inherent in silica shell enable easy surface bioconjugations to produce functionalised multiluminescent bioimaging markers. MPTMS and ammonia NH4OH were mixed and then incubated to grow the silica layer. Fluorescent markers such as QD 605 and

Figure 7. (Left) (a) Synthesis of dual fluorescent thiol-OS-QDs by Nakamura et al. to coat thin silica shell on Au nanorods using sodium silicate. Synthesis scheme for a single-step fabrication of double fluorescent thiol-OS-QDs. (Right) TEM images of thiol-OS-QDs with different silica shell thickness regulated by MPTMS dosage (b and c) [20].

rhodamine B were added. After reaction, the solution was centrifuged to discard unreacted contents. The aqueous silica encapsulation using Nakamura's ultra-fast technique took only 3 h, as opposed to using conventional reverse emulsion technique, which requires 1 day or using water-based sodium silicate, which requires 3 days. In addition, the presence of intrinsic thiol SH groups allows conjugation with thiol-reactive dyes or biomolecules, without the need for additional surface functionalization steps.

can be easily controlled through the concentration of precursor, reaction times and reaction

Niu et al. in 2010 [19] reported a facile route to fabricate homogenous, monodisperse, accurately size-tunable, SH-functionalized and magnetic silica-shell nanocomposite nanospheres (SHSSCNs) with sizes less than 100 nm. These nanocomposite nanospheres are synthesized based on the self-assembly of magnetic nanoparticles and a surface-priming copolymer polystyrene-block-poly(acrylic acid) (PS-b-PAA) in aqueous medium, followed by an easy silanization reaction with a silane precursor MPTMS to grow a silica shell on the interface of magnetite micelles at ambient temperature. (Step 1) Hydrophobic magnetite (Fe3O4) nanocores with a dimension of 5 nm were added to form magnetite-based micelles through their selfassembly in water with an amphiphilic block copolymer, PS-b-PAA. After filtration with deionized water, (Step 2) MPTMS was used to grow a hybrid silica layer by hydrolyzing and condensing the silane onto surface of the magnetic micelles under alkaline conditions. At the same time, SH groups were functionalised onto the surface of the nanocomposite particle in situ by using MPTMS silane as precursor, which is obviously distinct from traditional silicaencapsulation techniques reported to date. This is possible due to the fact that the core particles are encapsulated by a layer of polymeric shell, which allows the MPTMS silane to attach easily and directly onto the polymer shell's abundant surface hydroxyl groups. During the silicaencapsulation procedure, MPTMS performs as the silane precursor, thiol-functional linker and surface-capping stabilizers for the magnetite micelles and also for SH-modification of the

Nakamura et al. in 2011 [20] used MPTMS for surface modification of quantum dots and synthesis of the silica shell to grow a layer of organosilica dense shell. MPTMS, quantum dot QD605 and NH4OH were added and then heated for 3 h to grow the silica layer at 100C. Extra steps like filtration or exchanging the reacting medium or adding alcohol were required. The growth of the silica shell was easily done in 3 h (Figure 7). The procedure is fast and easy, and the thiol groups inherent in silica shell enable easy surface bioconjugations to produce functionalised multiluminescent bioimaging markers. MPTMS and ammonia NH4OH were mixed and then incubated to grow the silica layer. Fluorescent markers such as QD 605 and

Figure 7. (Left) (a) Synthesis of dual fluorescent thiol-OS-QDs by Nakamura et al. to coat thin silica shell on Au nanorods using sodium silicate. Synthesis scheme for a single-step fabrication of double fluorescent thiol-OS-QDs. (Right) TEM

images of thiol-OS-QDs with different silica shell thickness regulated by MPTMS dosage (b and c) [20].

temperatures, making aqueous routes far less complex than classic Stober method.

nanoparticles in the silanization step.

100 Novel Nanomaterials - Synthesis and Applications

Lee et al. in 2011 [21] reported that for in vivo applications, such as biological imaging and medicine, nanoparticles must be safe in cell toxicity. Gold nanorods are often surface-capped and stabilized with CTAB cetyltrimethylammonium bromide, which will disrupt cell biology and cell growth as a cytotoxic surfactant. Silica coating and surface encapsulation of single nanoparticles can provide a safer substitute to surface CTAB for both decreasing cytotoxicity and improving colloidal stability. Lee synthesized silica-coated gold nanorods in aqueous medium. In his synthesis technique, Au nanorods were mixed to a solution of (3-aminopropyldimethylethoxysilane) and MPTMS for 12 h under stirring after which NH4OH was added to form silica shell thickness < 2 nm. The authors found that the amount of amino-silane affected the growth of silica layer surrounding gold-NRs (Figure 8).

Shah et al. in 2014 [22] reported a novel direct silica coating in water-based silica-coating routes, which is alcohol-free one-step method using "pre-hydrolysed" MPTMS making silica coating far less complex than classic Stober method. This technique presents a facile, scalable, environmentally friendly process for thick coating of metal particles with a thick (5–30 nm) silica shell. The success of the process is using "pre-hydrolysed" MPTMS as silane precursor. Silica coating could be directly attributed to the strong and direct metal-thiol bond. Briefly, metal nanoparticles were washed and re-dispersed in deionized water before being combined into a pre-hydrolysed MPTMS in DI water. Ammonia was added as catalyst. While growth was easily scalable, the growth rate of the silica shell is controllable to 25, 11.7 and 10 nm/h. during condensation. The process slows down owing to eventual consumption of hydrolysed MPTMS. Different metal NPs (i.e., gold, silver, platinum) are successfully encapsulated with a silica layer under a completely alcohol-free environment, whose shell thickness can be easily and precisely synthesized by changing the growth duration (Figure 9). This novel water-based procedure is used for the fabrication of SERS-enhanced silver@SiO2 nanoparticles without the

Figure 8. (Left) (a) Schematic of aqueous synthesis procedure used by Lee et al. to coat thin silica shell on Au nanorods using sodium silicate. (Right) Images of AuNRs sized 50 nm, (b) SEM image of core-shell Au nanorods and (c) TEM image of silica shell (<2 nm) AuNRs [21].

fluorescence background, which is used for bioimaging as SERS markers. The facile silicaencapsulation procedure developed here presents a highly potential encapsulating method and protection for the metallic core nanoparticles. The resulting highly negatively charged and SERS-enhanced metal@SiO2 nanoparticles with thiol-functionalised surfaces hold great potential for biomedical applications. Such as fluorescence-free SERS-enhanced nanoparticles are useful for ultrasensitive bioimaging and biodetection applications.

surface of Fe3O4@SiO2 particles oxidized the pyrrole monomer to give PPy-Pd layer surrounding Fe3O4@SiO2-SO3H spheres. The nanocomposite Fe3O4@SiO2-SO3H@PPy-Pd nanoparticles

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Cui et al. in 2011 [24] proposed that highly stable Au@SiO2 nanoparticles with functional groups were easily fabricated by pre-hydrolysis of MPTES in a water together with gold nanoparticles. In the preparation of organosilica shell, heterofunctional poly(ethyleneglycol) (SH-PEG-COOH) solution was added into solution containing Au nanoparticles and then MPTES is added. The solubility of MPTES in the aqueous solution is very low and tends to form an emulsion. NH4OH was then added to adjust the solution pH to 8.5–9. The synthesis performed under static conditions for 48 h without stirring or shaking is the unique advantage of simplicity of the present method. The size of organosilica outer layer can be easily tuned by controlling the MPTES added. Steric hindrance caused by SH-PEG-COOH polymers, which are formed on gold surface, can inhibit gold particles aggregation, and therefore, polymer brushes perform a critical role in core-shell nanoparticles. Without any SH-PEG-COOH steric hindrance, gold particles will agglomerate and nucleate from the mixture even in MPTES solution. Au@SiO2 particles appear to be rounded in shape and singular in nature. Every particle is sized about 55 nm and has a single gold core. The silica layer is about 15–75 nm in thickness (Figure 11). The silica thickness could be simply tuned by the concentration of MPTES. The surface charge of these Au@SiO2 core-shell particles was measured to have 51.7 mV in zeta potential. Conventional Au@SiO2 core-shell particles are measured to have 38 mV in zeta potential. These zeta measurements confirm the improved stability of the

Gao et al. in 2015 [25] synthesized multifunctional gold nanostars through the direct silica coating in aqueous solution using "pre-hydrolyzed" MPTES. Organosilica shells are synthesized by first hydrolyzing 3-mercaptopropyltriethoxysilane (MPTES) as silane precursor in water. Due to the large amount of SH groups provided by MPTES, the tedious step of surface modification is avoided, and the synthesis technique is easily controlled and simple. Gao reported this simple synthesis of highly functional anisotropic Au nanostars via directly coating silica on the surface

Figure 11. (Left) Proposed growth mechanism of Au@organosilica nanoparticles by Cui et al. using MPTES in water medium. (Right) TEM images of gold@silica core-shells with the 55 nm gold particles and varying silica thicknesses:

reduced 4-nitrophenol and undergo Suzuki coupling reactions as a nanocatalyst.

4. MPTES-assisted silica coating

Au@SiO2 nanoparticles upon organosilica shell coating.

(a = 15 and b = 20 nm) [24].

Shang et al. in 2016 [23] synthesized iron oxide Fe3O4 nanoparticles coated with 10–20 nm thick silica shells using the same method reported by Shah et al. [22] after reacting for 2 h at room temperature with constant stirring (Figure 10). Hydrolysed MPTMS is formed by adding MPTM into deionized water and stirred [22]. Thereafter, Fe3O4 spheres were added into prehydrolysed MPTMS solution followed by adding ammonium hydroxide. Fe3O4@SiO2-SH spheres (0.1 g) were dispersed in H2O2, in order to oxidize SH groups to SO3H groups to form Fe3O4@SiO2-SO3H sulfonic acid-functionalised spheres. When Fe3O4@SiO2 spheres and PdCl2 solution were added, Pd2+ ions could be surface immobilized onto the surface of silica shell via coordination interaction between SO3H groups and Pd2+ ions. Pd2+ ions immobilized on the

Figure 9. (Top row) Schematic of aqueous synthesis procedure used by Shah et al. to coat thin silica shell on Ag nanospheres using MPTMS. (Bottom row) Ag@SiO2 (A) and Ag@SiO2 after coating for (B) 1 h, (C) 3 h and (D) 5 h [22].

Figure 10. (Left) Schematic of aqueous synthesis procedure used by Shang et al. to coat thin silica shell on Ag nanospheres using MPTMS. (Right) (A, B) TEM micrographs of resulting Fe3O4@SiO2-SH spheres [23].

surface of Fe3O4@SiO2 particles oxidized the pyrrole monomer to give PPy-Pd layer surrounding Fe3O4@SiO2-SO3H spheres. The nanocomposite Fe3O4@SiO2-SO3H@PPy-Pd nanoparticles reduced 4-nitrophenol and undergo Suzuki coupling reactions as a nanocatalyst.

#### 4. MPTES-assisted silica coating

fluorescence background, which is used for bioimaging as SERS markers. The facile silicaencapsulation procedure developed here presents a highly potential encapsulating method and protection for the metallic core nanoparticles. The resulting highly negatively charged and SERS-enhanced metal@SiO2 nanoparticles with thiol-functionalised surfaces hold great potential for biomedical applications. Such as fluorescence-free SERS-enhanced nanoparticles

Shang et al. in 2016 [23] synthesized iron oxide Fe3O4 nanoparticles coated with 10–20 nm thick silica shells using the same method reported by Shah et al. [22] after reacting for 2 h at room temperature with constant stirring (Figure 10). Hydrolysed MPTMS is formed by adding MPTM into deionized water and stirred [22]. Thereafter, Fe3O4 spheres were added into prehydrolysed MPTMS solution followed by adding ammonium hydroxide. Fe3O4@SiO2-SH spheres (0.1 g) were dispersed in H2O2, in order to oxidize SH groups to SO3H groups to form Fe3O4@SiO2-SO3H sulfonic acid-functionalised spheres. When Fe3O4@SiO2 spheres and PdCl2 solution were added, Pd2+ ions could be surface immobilized onto the surface of silica shell via coordination interaction between SO3H groups and Pd2+ ions. Pd2+ ions immobilized on the

Figure 9. (Top row) Schematic of aqueous synthesis procedure used by Shah et al. to coat thin silica shell on Ag nanospheres using MPTMS. (Bottom row) Ag@SiO2 (A) and Ag@SiO2 after coating for (B) 1 h, (C) 3 h and (D) 5 h [22].

Figure 10. (Left) Schematic of aqueous synthesis procedure used by Shang et al. to coat thin silica shell on Ag

nanospheres using MPTMS. (Right) (A, B) TEM micrographs of resulting Fe3O4@SiO2-SH spheres [23].

are useful for ultrasensitive bioimaging and biodetection applications.

102 Novel Nanomaterials - Synthesis and Applications

Cui et al. in 2011 [24] proposed that highly stable Au@SiO2 nanoparticles with functional groups were easily fabricated by pre-hydrolysis of MPTES in a water together with gold nanoparticles. In the preparation of organosilica shell, heterofunctional poly(ethyleneglycol) (SH-PEG-COOH) solution was added into solution containing Au nanoparticles and then MPTES is added. The solubility of MPTES in the aqueous solution is very low and tends to form an emulsion. NH4OH was then added to adjust the solution pH to 8.5–9. The synthesis performed under static conditions for 48 h without stirring or shaking is the unique advantage of simplicity of the present method. The size of organosilica outer layer can be easily tuned by controlling the MPTES added. Steric hindrance caused by SH-PEG-COOH polymers, which are formed on gold surface, can inhibit gold particles aggregation, and therefore, polymer brushes perform a critical role in core-shell nanoparticles. Without any SH-PEG-COOH steric hindrance, gold particles will agglomerate and nucleate from the mixture even in MPTES solution. Au@SiO2 particles appear to be rounded in shape and singular in nature. Every particle is sized about 55 nm and has a single gold core. The silica layer is about 15–75 nm in thickness (Figure 11). The silica thickness could be simply tuned by the concentration of MPTES. The surface charge of these Au@SiO2 core-shell particles was measured to have 51.7 mV in zeta potential. Conventional Au@SiO2 core-shell particles are measured to have 38 mV in zeta potential. These zeta measurements confirm the improved stability of the Au@SiO2 nanoparticles upon organosilica shell coating.

Gao et al. in 2015 [25] synthesized multifunctional gold nanostars through the direct silica coating in aqueous solution using "pre-hydrolyzed" MPTES. Organosilica shells are synthesized by first hydrolyzing 3-mercaptopropyltriethoxysilane (MPTES) as silane precursor in water. Due to the large amount of SH groups provided by MPTES, the tedious step of surface modification is avoided, and the synthesis technique is easily controlled and simple. Gao reported this simple synthesis of highly functional anisotropic Au nanostars via directly coating silica on the surface

Figure 11. (Left) Proposed growth mechanism of Au@organosilica nanoparticles by Cui et al. using MPTES in water medium. (Right) TEM images of gold@silica core-shells with the 55 nm gold particles and varying silica thicknesses: (a = 15 and b = 20 nm) [24].

of 900 nm and shell layer of 70 nm thick (Figure 13). During the hydrolysis and sol-gel reactions between MTMS and APTMS, MTMS diffused from micro-droplet to the oil surface. Microparticle size and shell dimensions can be easily adjusted by tuning the precursors dosage and concentrations, subjected to oil loading and sphere hardness requirements. In summary, silicacoated nanospheres were synthesized in a facile way [26]. In this one-step sol-gel technique, no

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Jin et al. in 2010 [28] presented a one-step "sol-gel" method for the formation of silica-coated core-shell microcapsules of phase change materials (PCM) in aqueous medium without using surfactants or dispersants. The main issues associated with the use of polymeric shell materials for encapsulating core PCM materials are their low thermal conductivity and undesirable residues such as formaldehyde. Silica shell material has a much higher thermal conductivity and chemical stability than the polymeric materials. There have been many reports on the use of silica as the shell material of microcapsules or nanocapsules for controlled release and targeted drug delivery, and protecting active agents. Various methods have been proposed for making silica-shelled PCM microcapsules, including Pickering stabilization, water/oil/ water emulsion templating using sodium silicate as the precursor, water/oil or oil/water emulsion templating using tetraethylorthosilicate (TEOS) as the precursor under acidic or basic conditions and hydrolysis-condensation reactions using tetraethylorthosilicate (TEOS) as the silica precursor with surfactants. Jin silica-coated nonadecane using methyltrimethoxysilane (MTMS) and 3-aminopropyl trimethoxysilane (APTMS) as silica precursors without any surfactants in an aqueous solution of distilled water. The synthesis was done under magnetic stirring to produce clear spherical, core-shell microparticles. A clear core-shell microstructure of few microns and a shell thickness about 500 nm (Figure 14). Positively charged amine groups of the APTMS suggest that MTMS in the oil phase diffuses to the oil-water interface where it reacts with the APTMS in the water phase to form a silica shell around the oil droplet. This effect was also observed by Ottenbrite [27] in their synthesis of organosilica nanoparticles, using aminopropyltriethoxysilane as the silica precursor. Jin found that increasing the volume

of APTMS in water above 2% results in gelling, and no microcapsules were formed.

Figure 13. (Left) Illustration of the design of gold nanostars@SiO2 coated using MPTES in water medium by Fei et al. SEM images of monodisperse organosilica nanocapsules of 900 nm diameter and 70 nm thick. (Right) SEM images of organosilica nanocapsules containing DEET about 300 nm diameter and 30 nm thickness (a) Scalebar 1 µm, (b) TEM

image, and (c) Scalebar 160 nm [26].

use of intermediate agents or surfactant were required.

Figure 12. (Left) Illustration of the design of gold nanostars@SiO2 coated using MPTES in water medium by Gao et al. (Right) TEM images of synthesized gold nanostars@SiO2 (a) Scalebar 50 nm (b) Scalebar 200 nm [25].

of nano-sized star-shaped gold cores and subsequently immobilization of gadolinium chelates. Highly dispersed and uniformly sized Au-Gd nanostars@SiO2 absorb strongly within the NIR spectrum and give out strong SERS waves which can strengthen magnetic resonance imaging. More silica coating was grown by adding MPTES and NH4OH to the water-nanostars mixture and stirring for 1–2 min. The suspension was left for half a day to produce star-shaped Au@Silica. Obviously, 60 nm Au@SiO2 nanostars and 20 nm silica shell, which are highly homogenous, are seen under TEM imaging (Figure 12). In addition, using this one-step procedure, in situ surface immobilization of SH groups onto Au@SiO2 nanostars on the outer surface can be easily done for subsequent bioconjugation. This nanocomposite with distinct core-shell structure preserves the star shape and as a result gives efficient conversion from photo to thermal effect. The silica outer coating could prevent Raman markers from detachment loss. This ensures that Au nanostars@SiO2 has greater Raman signals strength and sensing ability. A facile way via aqueous-coating silica directly on Au star-shaped cores to fabricate functional nanoparticles for applications in bioimaging and detection of cancer supports future research in silica-coated nanostructures for cancer therapy and diagnosis applications.

#### 5. MTMS-assisted silica coating

Core-shell nano/microcapsule is a fast expanding area of research because of their wide applications, such as effective storage, controlled release and strong adsorption. Batch production of oil-loaded nano/microcapsules favors facile one-step fabrication methods.

Fei et al. in 2006 [26] synthesized monodisperse functionalised silica microcapsules using a onepot procedure at room temperature. Silica encapsulation of oil microdroplet is reported. This is done without using any mediating agent or surfactant. Fei synthesized silica nano/microcapsules using self-silanization reaction at water-oil interface. N,N-diethyl-m-toluamide (DEET, Aldrich) was chosen as PCM to be silica-encapsulated. A water-oil solution of methyltrimethoxysilane (MTMS) and DEET was mixed into water and 3-aminopropyltrimethoxysilane (APTMS) and sonicated to get a white dispersion. When MTMS and APTMS are mixed, sol-gel reaction takes place and the solution reacts to give white-colored suspension. Adding MTMS or APTMS alone into water, there is no whitening or sol-gel reactions to be observed. Ottenbrite et al. [27] reported similar sol-gel effect. Monodisperse spheres are about 900 nm in diameter. The spheres had sizes of 900 nm and shell layer of 70 nm thick (Figure 13). During the hydrolysis and sol-gel reactions between MTMS and APTMS, MTMS diffused from micro-droplet to the oil surface. Microparticle size and shell dimensions can be easily adjusted by tuning the precursors dosage and concentrations, subjected to oil loading and sphere hardness requirements. In summary, silicacoated nanospheres were synthesized in a facile way [26]. In this one-step sol-gel technique, no use of intermediate agents or surfactant were required.

Jin et al. in 2010 [28] presented a one-step "sol-gel" method for the formation of silica-coated core-shell microcapsules of phase change materials (PCM) in aqueous medium without using surfactants or dispersants. The main issues associated with the use of polymeric shell materials for encapsulating core PCM materials are their low thermal conductivity and undesirable residues such as formaldehyde. Silica shell material has a much higher thermal conductivity and chemical stability than the polymeric materials. There have been many reports on the use of silica as the shell material of microcapsules or nanocapsules for controlled release and targeted drug delivery, and protecting active agents. Various methods have been proposed for making silica-shelled PCM microcapsules, including Pickering stabilization, water/oil/ water emulsion templating using sodium silicate as the precursor, water/oil or oil/water emulsion templating using tetraethylorthosilicate (TEOS) as the precursor under acidic or basic conditions and hydrolysis-condensation reactions using tetraethylorthosilicate (TEOS) as the silica precursor with surfactants. Jin silica-coated nonadecane using methyltrimethoxysilane (MTMS) and 3-aminopropyl trimethoxysilane (APTMS) as silica precursors without any surfactants in an aqueous solution of distilled water. The synthesis was done under magnetic stirring to produce clear spherical, core-shell microparticles. A clear core-shell microstructure of few microns and a shell thickness about 500 nm (Figure 14). Positively charged amine groups of the APTMS suggest that MTMS in the oil phase diffuses to the oil-water interface where it reacts with the APTMS in the water phase to form a silica shell around the oil droplet. This effect was also observed by Ottenbrite [27] in their synthesis of organosilica nanoparticles, using aminopropyltriethoxysilane as the silica precursor. Jin found that increasing the volume of APTMS in water above 2% results in gelling, and no microcapsules were formed.

of nano-sized star-shaped gold cores and subsequently immobilization of gadolinium chelates. Highly dispersed and uniformly sized Au-Gd nanostars@SiO2 absorb strongly within the NIR spectrum and give out strong SERS waves which can strengthen magnetic resonance imaging. More silica coating was grown by adding MPTES and NH4OH to the water-nanostars mixture and stirring for 1–2 min. The suspension was left for half a day to produce star-shaped Au@Silica. Obviously, 60 nm Au@SiO2 nanostars and 20 nm silica shell, which are highly homogenous, are seen under TEM imaging (Figure 12). In addition, using this one-step procedure, in situ surface immobilization of SH groups onto Au@SiO2 nanostars on the outer surface can be easily done for subsequent bioconjugation. This nanocomposite with distinct core-shell structure preserves the star shape and as a result gives efficient conversion from photo to thermal effect. The silica outer coating could prevent Raman markers from detachment loss. This ensures that Au nanostars@SiO2 has greater Raman signals strength and sensing ability. A facile way via aqueous-coating silica directly on Au star-shaped cores to fabricate functional nanoparticles for applications in bioimaging and detection of cancer supports future research in silica-coated

Figure 12. (Left) Illustration of the design of gold nanostars@SiO2 coated using MPTES in water medium by Gao et al.

(Right) TEM images of synthesized gold nanostars@SiO2 (a) Scalebar 50 nm (b) Scalebar 200 nm [25].

Core-shell nano/microcapsule is a fast expanding area of research because of their wide applications, such as effective storage, controlled release and strong adsorption. Batch produc-

Fei et al. in 2006 [26] synthesized monodisperse functionalised silica microcapsules using a onepot procedure at room temperature. Silica encapsulation of oil microdroplet is reported. This is done without using any mediating agent or surfactant. Fei synthesized silica nano/microcapsules using self-silanization reaction at water-oil interface. N,N-diethyl-m-toluamide (DEET, Aldrich) was chosen as PCM to be silica-encapsulated. A water-oil solution of methyltrimethoxysilane (MTMS) and DEET was mixed into water and 3-aminopropyltrimethoxysilane (APTMS) and sonicated to get a white dispersion. When MTMS and APTMS are mixed, sol-gel reaction takes place and the solution reacts to give white-colored suspension. Adding MTMS or APTMS alone into water, there is no whitening or sol-gel reactions to be observed. Ottenbrite et al. [27] reported similar sol-gel effect. Monodisperse spheres are about 900 nm in diameter. The spheres had sizes

tion of oil-loaded nano/microcapsules favors facile one-step fabrication methods.

nanostructures for cancer therapy and diagnosis applications.

5. MTMS-assisted silica coating

104 Novel Nanomaterials - Synthesis and Applications

Figure 13. (Left) Illustration of the design of gold nanostars@SiO2 coated using MPTES in water medium by Fei et al. SEM images of monodisperse organosilica nanocapsules of 900 nm diameter and 70 nm thick. (Right) SEM images of organosilica nanocapsules containing DEET about 300 nm diameter and 30 nm thickness (a) Scalebar 1 µm, (b) TEM image, and (c) Scalebar 160 nm [26].

9. MPTMS and MPTES are proven effective for metals and semiconductors-based core-mate-

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This chapter reviewed a series of potentially reliable and reproducible synthesis techniques for high-quality silica-encapsulated core-shell nanomaterials. We compared and provided insights into the development of four water-based techniques to fabricate novel silica-encapsulated nanocomposites with unique core-shell architecture. These synthesis routes take place under mild reaction conditions, at ambient temperature and pressure, which are potentially applicable for large-scale production. The facile synthesis of silica shell under aqueous environment allows greater repeatability, facile steps for silanization, as compared with alcohol-based conventional Stober methods. They possess great potential to upscale batch production and present a biocompatible system that can be utilized for biomedical applications. The universality of these techniques were experimentally proven and demonstrated using a wide variety of core materials including Ag, Au, Pt, Fe2O3, CdTe, Gd, PCM, DEET and many others. In future, we hope to update the scientific community with more up-to-date research studies on the development of similar useful techniques for a simple one-pot aqueous synthesis of silica-

We acknowledge the support and funding of NRF POC Grant NRF2015NRF-POC001-025 and

[1] Chaudhuri RG, Paria S. Core/shell nanoparticles: Classes, properties, synthesis mecha-

[2] Guerrero-Martınez A, Perez-Juste J, Liz-Marzán LM. Recent progress on silica coating of nanoparticles and related Nanomaterials. Advanced Materials. 2010;22:1182-1195. DOI:

nisms. Chemical Reviews. 2012;112:2373-2433. DOI: 10.1021/cr100449n

National University of Singapore for support and funding this work.

Address all correspondence to: bdgskw@nus.edu.sg

National University of Singapore, Singapore

10.1002/adma.200901263

10. MTMS is proven effective in silica encapsulation of oil-based nano/microstructures.

11. MTMS is used in tandem with APTMS as a catalyst under oil/water emulsion.

rials.

encapsulated core-shell nanomaterials.

Acknowledgements

Author details

Kwok Wei Shah

References

Figure 14. (Left) Illustration of the design of silica-coated nonadecane using MTMS in water medium by Jin et al. (Right) SEM images of PCM microencapsulated spherical capsules: (a) an intact microcapsule and (b) a broken microcapsule of core-shell microstructure [28].

#### 6. Conclusion

In this chapter, we summarize up-to-date the four major water-based silica-coating strategies for a variety of core nanomaterials. Currently, there are four successful water-based routes using different water-soluble silane precursors, namely sodium silicate, MPTMS, MPTES and MTMS. We reviewed these techniques in aqueous condition and their final core-shell morphologies. Their effective silica encapsulation leads to improved colloidal properties and creates new emerging materials based on silica-coated core-shell nanostructures. In particular, these simple large-scale aqueous synthesis techniques for silica-coated nanoparticles and their morphologies are important for paving new and practical applications. Four major water-based routes for silica-coating core-shell nanostructures are analyzed and their findings relating to sodium silicate, MPTMS, MPTES and MTMS are summarized below:


This chapter reviewed a series of potentially reliable and reproducible synthesis techniques for high-quality silica-encapsulated core-shell nanomaterials. We compared and provided insights into the development of four water-based techniques to fabricate novel silica-encapsulated nanocomposites with unique core-shell architecture. These synthesis routes take place under mild reaction conditions, at ambient temperature and pressure, which are potentially applicable for large-scale production. The facile synthesis of silica shell under aqueous environment allows greater repeatability, facile steps for silanization, as compared with alcohol-based conventional Stober methods. They possess great potential to upscale batch production and present a biocompatible system that can be utilized for biomedical applications. The universality of these techniques were experimentally proven and demonstrated using a wide variety of core materials including Ag, Au, Pt, Fe2O3, CdTe, Gd, PCM, DEET and many others. In future, we hope to update the scientific community with more up-to-date research studies on the development of similar useful techniques for a simple one-pot aqueous synthesis of silicaencapsulated core-shell nanomaterials.

#### Acknowledgements

6. Conclusion

core-shell microstructure [28].

106 Novel Nanomaterials - Synthesis and Applications

In this chapter, we summarize up-to-date the four major water-based silica-coating strategies for a variety of core nanomaterials. Currently, there are four successful water-based routes using different water-soluble silane precursors, namely sodium silicate, MPTMS, MPTES and MTMS. We reviewed these techniques in aqueous condition and their final core-shell morphologies. Their effective silica encapsulation leads to improved colloidal properties and creates new emerging materials based on silica-coated core-shell nanostructures. In particular, these simple large-scale aqueous synthesis techniques for silica-coated nanoparticles and their morphologies are important for paving new and practical applications. Four major water-based routes for silica-coating core-shell nanostructures are analyzed and their findings relating to

Figure 14. (Left) Illustration of the design of silica-coated nonadecane using MTMS in water medium by Jin et al. (Right) SEM images of PCM microencapsulated spherical capsules: (a) an intact microcapsule and (b) a broken microcapsule of

1. Sodium silicate as a precursor and its resultant shells are highly dependent on pH condi-

2. Sodium silicate precursor is suitable for very thin silica-shell coatings (<2 nm), whose

3. Sodium silicate can be precipitated into silica-shell coatings under fully aqueous conditions (thin coat <6 nm), as well as under ethanol-water conditions (thick coating >10 nm).

4. Sodium silicate provide only –OH groups, which may demand extra functionalisation step to provide additional groups, such as –SH, –NH2, –CHO, –COOH for conjugation.

5. MPTMS and MPTES precursors are effective in silica-coating metallic and semiconductor

6. "Pre-hydrolysis" of MPTMS and MPTES is an essential step to successful silica coating.

7. MPTMS and MPTES precursors produce silica-coated core-shell nanomaterials that are highly monodispersed, spherical and uniform with smooth outer surfaces. Thickness can

8. MPTMS and MPTES precursors provide intrinsic thiol –SH and silanol –OH functional

nanomaterials of different morphologies (nanospheres, nanorods, nanostars)

be easily controlled and thick dense shell can be coated within a few hours.

groups, without any need for extra step of surface modifications.

tions and coating process requires a long period of time, lasting a few days.

resultant silica-shell surface is often uneven, irregular and inhomogeneous.

sodium silicate, MPTMS, MPTES and MTMS are summarized below:

We acknowledge the support and funding of NRF POC Grant NRF2015NRF-POC001-025 and National University of Singapore for support and funding this work.

#### Author details

Kwok Wei Shah Address all correspondence to: bdgskw@nus.edu.sg National University of Singapore, Singapore

#### References


[3] Liu S, Han M-Y. Silica-coated metal nanoparticles. Chemistry, an Asian Journal. 2010;5: 36-45. DOI: 10.1002/adfm.200400427

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[21] Lee JY, Park W, Yib DK. Immunostimulatory effects of gold nanorod and silica-coated gold nanorod on RAW 264.7 mouse macrophages. Toxicology Letters. 2012;209:51-57 DOI:

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**Chapter 7**

Provisional chapter

**Delamination and Longitudinal Cracking in**

on Cutting Tool Wear Mechanism and Tool Life

**Tool Life**

Alexey Vereschaka, Sergey Grigoriev,

Alexey Vereschaka, Sergey Grigoriev,

Anatoliy Aksenenko and Andre Batako

Anatoliy Aksenenko and Andre Batako

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Nikolay Sitnikov, Gaik Oganyan,

Nikolay Sitnikov, Gaik Oganyan,

http://dx.doi.org/10.5772/intechopen.72257

delamination, nanoscale structures

Abstract

**Multilayered Composite Nanostructured Coatings and**

**Their Influence on Cutting Tool Wear Mechanism and**

The wear and failure mechanism for multilayered nanostructured coatings has a number of significant differences from the one typical for monolithic single-layered coatings. In particular, while the strength of adhesion bonds at the "substrate-coating" boundary is important for monolithic coatings, then for multilayered nanostructured coatings, the strength of adhesion and cohesion bonds at interlayer boundaries and boundaries of separate nano-sublayers becomes of significant significance. Meanwhile, the delamination arising in the structure of multilayered nanostructured coatings can have both negative (leading to loss of coating uniformity and subsequent failure of coating) and positive influences (due to decrease of internal stresses and inhibition of transverse cracking). Various mechanisms of formation of longitudinal cracks and delaminations in coatings on rake tool faces, which vary based on the compositions and architectures of the coatings, are studied. In addition, the influence of internal defects, including embedded microdrops and pores, on the formation of cracks and delaminations and the failure of coatings is discussed. The importance of ensuring a balance of the basic properties of coatings to achieve high wear resistance and maximum tool life of coated metal-cutting tools is shown. The properties of coatings and the natures of their failures, as investigated during scratch testing and dry turning of steel C45, are provided.

Keywords: wear-resistant coatings, wear, crack, fracture, tool life, PVD coatings,

© 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 IntechOpen. 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.

Delamination and Longitudinal Cracking in Multilayered

DOI: 10.5772/intechopen.72257

Composite Nanostructured Coatings and Their Influence

**Delamination and Longitudinal Cracking in Multilayered Composite Nanostructured Coatings and Their Influence on Cutting Tool Wear Mechanism and Tool Life** Provisional chapter Delamination and Longitudinal Cracking in Multilayered Composite Nanostructured Coatings and Their Influence

DOI: 10.5772/intechopen.72257

on Cutting Tool Wear Mechanism and Tool Life

Alexey Vereschaka, Sergey Grigoriev, Nikolay Sitnikov, Gaik Oganyan, Anatoliy Aksenenko and Andre Batako Alexey Vereschaka, Sergey Grigoriev, Nikolay Sitnikov, Gaik Oganyan,

Anatoliy Aksenenko and Andre Batako

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.72257

#### Abstract

The wear and failure mechanism for multilayered nanostructured coatings has a number of significant differences from the one typical for monolithic single-layered coatings. In particular, while the strength of adhesion bonds at the "substrate-coating" boundary is important for monolithic coatings, then for multilayered nanostructured coatings, the strength of adhesion and cohesion bonds at interlayer boundaries and boundaries of separate nano-sublayers becomes of significant significance. Meanwhile, the delamination arising in the structure of multilayered nanostructured coatings can have both negative (leading to loss of coating uniformity and subsequent failure of coating) and positive influences (due to decrease of internal stresses and inhibition of transverse cracking). Various mechanisms of formation of longitudinal cracks and delaminations in coatings on rake tool faces, which vary based on the compositions and architectures of the coatings, are studied. In addition, the influence of internal defects, including embedded microdrops and pores, on the formation of cracks and delaminations and the failure of coatings is discussed. The importance of ensuring a balance of the basic properties of coatings to achieve high wear resistance and maximum tool life of coated metal-cutting tools is shown. The properties of coatings and the natures of their failures, as investigated during scratch testing and dry turning of steel C45, are provided.

Keywords: wear-resistant coatings, wear, crack, fracture, tool life, PVD coatings, delamination, nanoscale structures

© 2018 The Author(s). Licensee IntechOpen. 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 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.

#### 1. Introduction

#### 1.1. Background

Further increases in efficiency of machining and cutting speeds as well as tightening of reliability requirements associated with greater levels of automation of production result in the need to create new tool materials with enhanced performance characteristics. One way to improve the performance characteristics of tool materials is to enhance their surface properties by applying modified coatings [1]. In turn, the properties of modified coatings continue to be improved, and their architecture and elemental composition become more complicated. In particular, multilayered composite coatings, nanostructured and gradient coatings, and coatings with multicomponent elemental composition have been used extensively in recent years [2]. The use of a multilayered architecture of coatings and the use of nanostructured technology can significantly improve the performance characteristics of a new generation of coatings. However, along with the use of such coatings come new problems that did not occur with monolithic coatings of the first generation. In particular, problems arose concerning interlayer delamination and formation of specific longitudinal cracks in the structure of coating. A large number of studies examining problems of cracking have been conducted. The general assumption is that the formation of microcracks is associated with the displacement of dislocations [3– 6]. A number of mechanisms for the formation of dislocation microcracks are well known [3, 4, 6]. In principle, those mechanisms provide for blocking of the progress of dislocation by some obstacle (e.g., a grain boundary, a boundary of nanolayers, or inclusion). If in some slip plane dislocations stop before a sufficiently powerful obstacle, then a cluster of dislocations is formed, and it causes a high concentration of stresses at the obstacle. This concentration of stresses results in formation of a dislocation microcrack. It should be noted that the problems of crack formation and delamination in the structures of multilayered coatings have not been studied as thoroughly as have other aspects of operation and wear of such coatings.

and cracking of physical vapor deposition (PVD)-coated (TiN, (Ti,Si)N, (Ti,Al)N, and (Al,Cr)N) polycrystalline cubic boron nitride. Koseki et al. [22] examined the cutting performance of TiNcoated cutting tools. Defects (e.g., droplets, voids) in the coating were found to be the starting point of damage. The breakdown region is enlarged as the work material is caught in the damaged portion of the coating. Kumar and Curtin [23] considered the probable mechanisms of development and inhibition of cracks in microstructures: particularly at crack bridging by ductile ligaments, crack deflection by second-phase particles, microcrack formation, and stressinduced phase transformations. The same paper also includes an overview of methods for modeling the development of cracks using FEM and incorporating cohesive elements at the continuum level, as well as discrete dislocation methodology at the mesoscopic level, and coupled atomistic/continuum methods that transition atomic level information to the microscopic level. A large number of studies have been devoted to the investigation of causes and conditions for the formation of delaminations in multilayered composite macrostructures. To predict the occurrence of delaminations, the methods of layer-wise interface elements [24, 25], classical finite element analysis (FEA) [25, 26], and the virtual crack closure technique (VCCT) [27] are widely used. The issues concerning delamination of multilayered nanostructured coat-

Delamination and Longitudinal Cracking in Multilayered Composite Nanostructured Coatings and Their Influence…

Starting from the theory of crack formation [3–5], the concentration of local tensile stresses σld in the head of a series of edge dislocations caused by the action of a number of edge stresses τ

> ffiffiffiffiffi d 2x r

where 2d is the length of the slip band or the distance between the slip bands (the value may also correspond to crystalline grain diameter), x is the distance from the strip to the head of the cluster of dislocations, and τ<sup>i</sup> is the stress of resistance to movement of dislocations (friction stress). If the local stress reaches the theoretical strength of the crystalline body σtheor, deter-

> ffiffiffiffiffiffi Eγ a0

> > ffiffiffiffiffiffi Eγ a0

s

s

σtheor ¼

Consequently, the criterion for the formation of a microcrack is as follows:

ffiffiffiffiffi d 2x r

where a<sup>0</sup> is the equilibrium distance between atoms, E is the modulus of elasticity, and γ is the plastic shear deformation, then conditions arise for the formation of a dislocation microcrack.

ð Þ τ � τ<sup>i</sup> ≥

The number of positive or negative dislocations in a flat cluster near an obstacle can be expressed by an approximate formula (assuming E ≈ 2G, where G is the shear modulus):

ð Þ τ � τ<sup>i</sup> (1)

http://dx.doi.org/10.5772/intechopen.72257

113

(2)

(3)

σld ¼

ings are also discussed in details in [28].

1.3. Mechanisms of crack formation

mined by the equation

can be determined using the following equation [29]:

#### 1.2. Literature review

Tabakov et al. [7, 8] considered mechanisms of cracking with respect to single-layer macroscale coatings on the basis of systems composed of TiN, TiCN, (Ti,Zr)N, and (Ti,Zr)CN. They discovered that coatings of a complex composition of (Ti,Zr)N and (Ti,Zr)CN are characterized by better resistance to intensive cracking. Tabakov et al. also considered multilayered coatings with macroscale structure: in particular, on the basis of systems composed of TiCN-(Ti,Zr)N-TiN, TiN-(Ti,Zr)N-TiN, TiCN-(Ti,Al)N-TiN, and TiCN-(Ti,Мо)N-TiN [9]. These studies proved that the introduction of zirconium nitride in the coating composition significantly reduces the tendency to cracking. The problems of cracking and brittle fracture of coatings consisting of Ti-TiN- (Ti,Cr,Al)N, Zr-(Zr,Cr)N-CrN, and Ti-TiN-(Ti,Cr,Al)N and Ti-(Al,Cr)N-(Ti,Al)N, Ti-(Al,Cr)N-(Ti, Cr,Al)N, and Zr-(Al,Cr)N-(Zr,Cr,Al)N also were addressed in papers [10–16]. A detailed review of existing papers in the field of crack formation in multilayered coatings, with classification of types of cracks and analysis of the mechanisms of their formation, is given in [17]. The topic of mathematical modeling of cracking in multilayered coatings with the use of an axis-symmetrical finite element method (FEM) model was considered by Skordaris et al. [18]. Wu et al. [19, 20] modeled cracking in single-layer coatings within the framework of linear elastic fracture mechanics (LEFM). M'Saoubi et al. [21] investigated the nature of wear, including brittle fracture and cracking of physical vapor deposition (PVD)-coated (TiN, (Ti,Si)N, (Ti,Al)N, and (Al,Cr)N) polycrystalline cubic boron nitride. Koseki et al. [22] examined the cutting performance of TiNcoated cutting tools. Defects (e.g., droplets, voids) in the coating were found to be the starting point of damage. The breakdown region is enlarged as the work material is caught in the damaged portion of the coating. Kumar and Curtin [23] considered the probable mechanisms of development and inhibition of cracks in microstructures: particularly at crack bridging by ductile ligaments, crack deflection by second-phase particles, microcrack formation, and stressinduced phase transformations. The same paper also includes an overview of methods for modeling the development of cracks using FEM and incorporating cohesive elements at the continuum level, as well as discrete dislocation methodology at the mesoscopic level, and coupled atomistic/continuum methods that transition atomic level information to the microscopic level. A large number of studies have been devoted to the investigation of causes and conditions for the formation of delaminations in multilayered composite macrostructures. To predict the occurrence of delaminations, the methods of layer-wise interface elements [24, 25], classical finite element analysis (FEA) [25, 26], and the virtual crack closure technique (VCCT) [27] are widely used. The issues concerning delamination of multilayered nanostructured coatings are also discussed in details in [28].

#### 1.3. Mechanisms of crack formation

1. Introduction

112 Novel Nanomaterials - Synthesis and Applications

1.1. Background

1.2. Literature review

Further increases in efficiency of machining and cutting speeds as well as tightening of reliability requirements associated with greater levels of automation of production result in the need to create new tool materials with enhanced performance characteristics. One way to improve the performance characteristics of tool materials is to enhance their surface properties by applying modified coatings [1]. In turn, the properties of modified coatings continue to be improved, and their architecture and elemental composition become more complicated. In particular, multilayered composite coatings, nanostructured and gradient coatings, and coatings with multicomponent elemental composition have been used extensively in recent years [2]. The use of a multilayered architecture of coatings and the use of nanostructured technology can significantly improve the performance characteristics of a new generation of coatings. However, along with the use of such coatings come new problems that did not occur with monolithic coatings of the first generation. In particular, problems arose concerning interlayer delamination and formation of specific longitudinal cracks in the structure of coating. A large number of studies examining problems of cracking have been conducted. The general assumption is that the formation of microcracks is associated with the displacement of dislocations [3– 6]. A number of mechanisms for the formation of dislocation microcracks are well known [3, 4, 6]. In principle, those mechanisms provide for blocking of the progress of dislocation by some obstacle (e.g., a grain boundary, a boundary of nanolayers, or inclusion). If in some slip plane dislocations stop before a sufficiently powerful obstacle, then a cluster of dislocations is formed, and it causes a high concentration of stresses at the obstacle. This concentration of stresses results in formation of a dislocation microcrack. It should be noted that the problems of crack formation and delamination in the structures of multilayered coatings have not been

studied as thoroughly as have other aspects of operation and wear of such coatings.

Tabakov et al. [7, 8] considered mechanisms of cracking with respect to single-layer macroscale coatings on the basis of systems composed of TiN, TiCN, (Ti,Zr)N, and (Ti,Zr)CN. They discovered that coatings of a complex composition of (Ti,Zr)N and (Ti,Zr)CN are characterized by better resistance to intensive cracking. Tabakov et al. also considered multilayered coatings with macroscale structure: in particular, on the basis of systems composed of TiCN-(Ti,Zr)N-TiN, TiN-(Ti,Zr)N-TiN, TiCN-(Ti,Al)N-TiN, and TiCN-(Ti,Мо)N-TiN [9]. These studies proved that the introduction of zirconium nitride in the coating composition significantly reduces the tendency to cracking. The problems of cracking and brittle fracture of coatings consisting of Ti-TiN- (Ti,Cr,Al)N, Zr-(Zr,Cr)N-CrN, and Ti-TiN-(Ti,Cr,Al)N and Ti-(Al,Cr)N-(Ti,Al)N, Ti-(Al,Cr)N-(Ti, Cr,Al)N, and Zr-(Al,Cr)N-(Zr,Cr,Al)N also were addressed in papers [10–16]. A detailed review of existing papers in the field of crack formation in multilayered coatings, with classification of types of cracks and analysis of the mechanisms of their formation, is given in [17]. The topic of mathematical modeling of cracking in multilayered coatings with the use of an axis-symmetrical finite element method (FEM) model was considered by Skordaris et al. [18]. Wu et al. [19, 20] modeled cracking in single-layer coatings within the framework of linear elastic fracture mechanics (LEFM). M'Saoubi et al. [21] investigated the nature of wear, including brittle fracture Starting from the theory of crack formation [3–5], the concentration of local tensile stresses σld in the head of a series of edge dislocations caused by the action of a number of edge stresses τ can be determined using the following equation [29]:

$$
\sigma\_{\rm lld} = \sqrt{\frac{d}{2\chi}} (\tau - \tau\_i) \tag{1}
$$

where 2d is the length of the slip band or the distance between the slip bands (the value may also correspond to crystalline grain diameter), x is the distance from the strip to the head of the cluster of dislocations, and τ<sup>i</sup> is the stress of resistance to movement of dislocations (friction stress). If the local stress reaches the theoretical strength of the crystalline body σtheor, determined by the equation

$$
\sigma\_{\text{theor}} = \sqrt{\frac{\text{E}\overline{\text{\textdegree}}}{a\_0}} \tag{2}
$$

where a<sup>0</sup> is the equilibrium distance between atoms, E is the modulus of elasticity, and γ is the plastic shear deformation, then conditions arise for the formation of a dislocation microcrack. Consequently, the criterion for the formation of a microcrack is as follows:

$$
\sqrt{\frac{d}{2\chi}}(\tau - \tau\_i) \ge \sqrt{\frac{E\gamma}{a\_0}}\tag{3}
$$

The number of positive or negative dislocations in a flat cluster near an obstacle can be expressed by an approximate formula (assuming E ≈ 2G, where G is the shear modulus):

$$m \approx \frac{(\pi - \pi\_i)}{bE} \tag{4}$$

intergranular fracture can occur in monolithic single-layer coatings (in particular, transverse cracks in the columnar crystalline structure of TiN are formed under the above mechanism). As a rule, several mechanisms of failure take place simultaneously, and as a consequence, a mixed type of failure occurs. Typically, three types of loading or displacement of the points of crack surfaces under the influence of an external load are considered [3, 4, 6]. The first type (type 1) includes the formation of normal detachment cracks, characterized by movement of the points of the crack surface under the action of a load in the direction perpendicular to the plane of the crack. In this case, the crack tends to open. Cracks of the transverse type (type 2) are cracks in which points of surfaces are displaced across the front of the crack (leading edge of the crack). Finally, longitudinal shear cracks (type 3) are characterized by displacement of the points of the crack surface along its front. It should be noted that these types of loading can be combined, thus forming complex types of loading. The conditions of operation of multilayered nanostructured modified coatings for a metal-cutting tool are most typically characterized by loading of type 1 (typical for the rake face of the tool, due to the constant formation and failure of adhesion bridges with the tool being machined) and type 2 (typical for the flank face of the tool, due to longitudinal compressive stresses and the resulting plastic deformation of a substrate). To describe the delamination process, it is also possible to use strain energy release rate (SERR), which represents energy dissipated during fracture per unit of newly created fracture surface area [31]. Delamination growth rates were corre-

http://dx.doi.org/10.5772/intechopen.72257

115

Delamination and Longitudinal Cracking in Multilayered Composite Nanostructured Coatings and Their Influence…

lated with the SERR by means of the Paris relation [32, 33]:

Figure 1. Mechanisms of failure: (a) viscous fracture, (b) transcrystallite cleavage, and (c) intercrystallite fracture.

Figure 2. Pattern of the micromechanism of growth of a viscous crack: (a) inclusions at the crack tip, (b) growth of

micropores in front of the crack tip, and (c) micropores merging with the crack tip.

Assuming the additional condition of x ≈ a0, from the joint solution of Eqs. (3) and (4), we obtain the condition necessary for the formation of a microcrack under the dislocation mechanism:

$$(\pi - \pi\_i)nb = 2\gamma \tag{5}$$

Analysis of the given conditions for the formation of a dislocation microcrack results in the following conclusion. Local tensile stresses in the head of a number of dislocations are formed primarily because of tangential stresses τ and are not related in any way to tensile stresses (i.e., only shear stresses are crucial for the initiation of a microcrack). The defects already existing or emerging in the early stages of deformation of solid bodies result in the initiation and development of the failure processes. Various mechanisms of failure are realized depending on the structural and stress strain states of a solid body and also depending on the external medium [3, 6, 30]. The following are the most common mechanisms of microfailure of metals (Figure 1):


<sup>n</sup> <sup>≈</sup> ð Þ <sup>τ</sup> � <sup>τ</sup><sup>i</sup>

114 Novel Nanomaterials - Synthesis and Applications

Assuming the additional condition of x ≈ a0, from the joint solution of Eqs. (3) and (4), we obtain the condition necessary for the formation of a microcrack under the dislocation mechanism:

Analysis of the given conditions for the formation of a dislocation microcrack results in the following conclusion. Local tensile stresses in the head of a number of dislocations are formed primarily because of tangential stresses τ and are not related in any way to tensile stresses (i.e., only shear stresses are crucial for the initiation of a microcrack). The defects already existing or emerging in the early stages of deformation of solid bodies result in the initiation and development of the failure processes. Various mechanisms of failure are realized depending on the structural and stress strain states of a solid body and also depending on the external medium [3, 6, 30]. The following are the most common mechanisms of microfailure of metals (Figure 1): 1. Viscous fracture (Figure 1a). This mechanism of failure is caused by the formation of micropores near inclusions or particles of the second phase, their growth, localization of microplastic deformation in the crosspieces between the pores, and, in the final stage, the fusion of micropores and the break of the bridges. The model and criterion for the formation of micropores are as follows: a micropore is formed when the cohesive stresses (bond stresses between the inclusion particles and the matrix) reach a critical stress. The existing models for the formation, growth, and fusion of micropores can be used to analyze micromechanisms of viscous fracture of a solid body in front of a crack tip. When the stresses in front of the inclusions (before a crack tip) reach critical values, the micropores are being formed. Further growth of micropores and localization of plastic deformation results in plastic blunting of crack tip, merging of micropores with crack tip, and subsequent growth of a viscous crack (Figure 2). In multilayered composite nanostructured coatings, such a mechanism of cracking can occur primarily in the formation of delaminations at interlayer boundaries, as well as at the boundaries of nano-sublayers. 2. Transcrystallite cleavage (Figure 1b). This mechanism is characterized by failure of a solid body (spread of a crack) along certain crystallographic planes. In polycrystalline bodies, the process of transcrystallite cleavage is realized not in one crystallographic plane, but through the distribution and subsequent integration of a multitude of microcracks of the cleavage that arises in a certain family of crystallographic grain planes. As a rule, transcrystallite cleavage is of a brittle nature, although plastic deformation processes are also possible. The described mechanism can be typical for the formation of cracks in monolithic single-layer coatings (TiN, ZrN, etc.), as well as in monolithic layers of multilayered coatings. In coatings with nano-sublayers, the development of cracks under the above mechanism is significantly constrained by the boundaries of nano-sublayers. 3. Intercrystallite fracture (Figure 1c). This mechanism consists of the initiation and propagation of microcracks along grain boundaries. This failure mechanism is related to the fact that the fracture energy necessary for propagation of a crack along the grain boundaries is lower than the corresponding energy of the transcrystallite cleavage. In the coatings,

bE (4)

ð Þ τ � τ<sup>i</sup> nb ¼ 2γ (5)

intergranular fracture can occur in monolithic single-layer coatings (in particular, transverse cracks in the columnar crystalline structure of TiN are formed under the above mechanism). As a rule, several mechanisms of failure take place simultaneously, and as a consequence, a mixed type of failure occurs. Typically, three types of loading or displacement of the points of crack surfaces under the influence of an external load are considered [3, 4, 6]. The first type (type 1) includes the formation of normal detachment cracks, characterized by movement of the points of the crack surface under the action of a load in the direction perpendicular to the plane of the crack. In this case, the crack tends to open. Cracks of the transverse type (type 2) are cracks in which points of surfaces are displaced across the front of the crack (leading edge of the crack). Finally, longitudinal shear cracks (type 3) are characterized by displacement of the points of the crack surface along its front. It should be noted that these types of loading can be combined, thus forming complex types of loading. The conditions of operation of multilayered nanostructured modified coatings for a metal-cutting tool are most typically characterized by loading of type 1 (typical for the rake face of the tool, due to the constant formation and failure of adhesion bridges with the tool being machined) and type 2 (typical for the flank face of the tool, due to longitudinal compressive stresses and the resulting plastic deformation of a substrate). To describe the delamination process, it is also possible to use strain energy release rate (SERR), which represents energy dissipated during fracture per unit of newly created fracture surface area [31]. Delamination growth rates were correlated with the SERR by means of the Paris relation [32, 33]:

Figure 1. Mechanisms of failure: (a) viscous fracture, (b) transcrystallite cleavage, and (c) intercrystallite fracture.

Figure 2. Pattern of the micromechanism of growth of a viscous crack: (a) inclusions at the crack tip, (b) growth of micropores in front of the crack tip, and (c) micropores merging with the crack tip.

$$\frac{dB\_d}{dN} = \mathsf{C}f(\mathsf{G})^{n\_p} \tag{6}$$

• Formation of the nanoscale structure of the deposited coating layers (grain size, sublayer thickness) with high density due to the energy supplied to the deposited condensate and transformation of the kinetic energy of the bombarding ions into thermal energy in local surface volumes of carbide material at an extremely high rate of approximately 1014 K s<sup>1</sup>

Delamination and Longitudinal Cracking in Multilayered Composite Nanostructured Coatings and Their Influence…

When choosing the composition of NMCC layers, in forming the coating of the three-layered architecture [2, 10], the Hume-Rothery rule was used. This rule states that the difference in atomic dimensions in contacting compounds should not exceed 20% [34]. The parameters used

An uncoated carbide tool and a carbide tool with "reference" coating TiN, deposited via standard vacuum-arc technology of arc-PVD, were used as objects for comparative studies of tool life.

For microstructural studies of samples of carbide with coatings, a raster electron microscope FEI Quanta 600 FEG was used. The studies of chemical composition were conducted using the same raster electron microscope. To perform X-ray microanalysis, characteristic X-ray emissions resulting from electron bombardment of a sample were examined. The hardness (HV) of coatings was determined by measuring the indentation at low loads according to the method of Oliver and Pharr [35], which was conducted on a micro-indentometer microhardness tester (CSM Instruments) at a fixed load of 300 mN. The penetration depth of the indenter was monitored so that it did not exceed 10–20% of the coating thickness to limit the influence of the substrate. The adhesion characteristics were studied on a Nanovea scratch tester, which represents a diamond cone with apex angle of 120 and radius of top curvature of 100 μm. The tests were conducted with the load linearly increasing from 0.05 to 40 N. Crack length was 5 mm. Each sample was subjected to three trials. The obtained curves were used to determine two parameters: the first critical load, LC1, at which the first cracks appeared in the coating,

A study of the cutting properties of the tool made of carbide with developed NMCC was conducted using a lathe CU 500 MRD for longitudinal turning of steel C45 (HB 200). In the experiment, the cutters featured mechanical fastening of inserts made of carbide (WC + 15%

Process pN (Pa) U (V) IAl (A) IZrNb (A) ITi (A) ICr (A) Pumping and heating of vacuum chamber 0.06 +20 120 80 65 75

Deposition of coating 0.36 800 DC 160 75 55 70 Cooling of products 0.06 — —— ——

Note: ITi = current of titanium cathode, IAl = current of aluminum cathode, IZrNb = current of zirconium-niobium cathode,

f = 10 kHz, 2:1

80 — ——

at each stage of the deposition process of NMCC are shown in Table 1.

and the second critical load, LC2, which caused the total failure of the coating.

Heating and cleaning of products with gaseous plasma 2.0 100 DC/900 AC

Table 1. Parameters of stages of the technological process of deposition of NMCC.

ICr = current of chromium cathode, pN = gas pressure in chamber, and U = voltage on substrate.

2.2. Microstructural studies

2.3. Study of cutting properties

.

117

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where B is delamination length (mm), N is the number of cycles, C is the Paris coefficient for delamination growth, G is the strain energy release rate (N/mm), np is the Paris exponent for delamination growth, and C and n are empirically determined parameters that depend on the materials and the temperature (and possibly other factors). As yet, no consensus has been reached on the correct form of f(G). In particular, f(G) can be represented as SERR at maximum fatigue load, Gmax, and within the SERR range, ΔG. In turn, ΔG can be determined from the following formula [25]:

$$
\Delta \mathbf{G} = \left(\sqrt{\mathbf{G}\_{\text{max}}} - \sqrt{\mathbf{G}\_{\text{min}}}\right)^2 \tag{7}
$$

The dependence (1) also can be represented by [26] as follows:

$$\frac{dB\_d}{dN} = \mathbb{C}(\mathbf{G}\_{max})^{n\_p} \tag{8}$$

where Gmax is the strain energy release rate at maximum fatigue load (N/mm).

Thus, the task of the present work is to study the process of delamination between layers of multilayer coatings and between nano-sublayers of nanostructured coatings. These processes have a significant impact on the overall performance of modifying coatings and products with such coatings (in particular, metal-cutting tools). An important feature of this work is that not only laboratory samples but also cutting tool samples that underwent cutting tests in real production conditions were considered. The peculiarities of delamination were investigated depending both on the elemental composition of the coatings and on their architecture (total coating thickness and thickness of the nanolayers).

#### 2. Materials and methods

#### 2.1. Deposition method

For deposition of nanoscale multilayered composite coatings (NMCC), a vacuum-arc VIT-2 unit [2], which was designed for the synthesis of coatings on substrates of various tool materials, was used. The unit was equipped with an arc evaporator with filtration of vapor-ion flow. In this study, the process, termed filtered cathodic vacuum-arc deposition (FCVAD) [10–16], was used for deposition of coatings on the tools to significantly reduce the formation of the droplet phase during formation of the coating. The use of the FCVAD process does not cause structural changes in carbide and provides the following:


• Formation of the nanoscale structure of the deposited coating layers (grain size, sublayer thickness) with high density due to the energy supplied to the deposited condensate and transformation of the kinetic energy of the bombarding ions into thermal energy in local surface volumes of carbide material at an extremely high rate of approximately 1014 K s<sup>1</sup> .

When choosing the composition of NMCC layers, in forming the coating of the three-layered architecture [2, 10], the Hume-Rothery rule was used. This rule states that the difference in atomic dimensions in contacting compounds should not exceed 20% [34]. The parameters used at each stage of the deposition process of NMCC are shown in Table 1.

An uncoated carbide tool and a carbide tool with "reference" coating TiN, deposited via standard vacuum-arc technology of arc-PVD, were used as objects for comparative studies of tool life.

#### 2.2. Microstructural studies

dBd

<sup>Δ</sup><sup>G</sup> <sup>¼</sup> ffiffiffiffiffiffiffiffiffiffi

dBd

where Gmax is the strain energy release rate at maximum fatigue load (N/mm).

The dependence (1) also can be represented by [26] as follows:

coating thickness and thickness of the nanolayers).

structural changes in carbide and provides the following:

stresses in the surface layers of the carbide material.

• High adhesive strength of the coating in relation to the carbide substrate.

2. Materials and methods

2.1. Deposition method

Gmax <sup>p</sup> � ffiffiffiffiffiffiffiffiffi

Thus, the task of the present work is to study the process of delamination between layers of multilayer coatings and between nano-sublayers of nanostructured coatings. These processes have a significant impact on the overall performance of modifying coatings and products with such coatings (in particular, metal-cutting tools). An important feature of this work is that not only laboratory samples but also cutting tool samples that underwent cutting tests in real production conditions were considered. The peculiarities of delamination were investigated depending both on the elemental composition of the coatings and on their architecture (total

For deposition of nanoscale multilayered composite coatings (NMCC), a vacuum-arc VIT-2 unit [2], which was designed for the synthesis of coatings on substrates of various tool materials, was used. The unit was equipped with an arc evaporator with filtration of vapor-ion flow. In this study, the process, termed filtered cathodic vacuum-arc deposition (FCVAD) [10–16], was used for deposition of coatings on the tools to significantly reduce the formation of the droplet phase during formation of the coating. The use of the FCVAD process does not cause

• Control of the level of the "healing" of energy impact on surface defects in carbide in the form of microcracks and micropores and formation of favorable residual compressive

� � p <sup>2</sup>

Gmin

following formula [25]:

116 Novel Nanomaterials - Synthesis and Applications

where B is delamination length (mm), N is the number of cycles, C is the Paris coefficient for delamination growth, G is the strain energy release rate (N/mm), np is the Paris exponent for delamination growth, and C and n are empirically determined parameters that depend on the materials and the temperature (and possibly other factors). As yet, no consensus has been reached on the correct form of f(G). In particular, f(G) can be represented as SERR at maximum fatigue load, Gmax, and within the SERR range, ΔG. In turn, ΔG can be determined from the

dN <sup>¼</sup> Cf Gð Þnp (6)

dN <sup>¼</sup> C Gð Þ max np (8)

(7)

For microstructural studies of samples of carbide with coatings, a raster electron microscope FEI Quanta 600 FEG was used. The studies of chemical composition were conducted using the same raster electron microscope. To perform X-ray microanalysis, characteristic X-ray emissions resulting from electron bombardment of a sample were examined. The hardness (HV) of coatings was determined by measuring the indentation at low loads according to the method of Oliver and Pharr [35], which was conducted on a micro-indentometer microhardness tester (CSM Instruments) at a fixed load of 300 mN. The penetration depth of the indenter was monitored so that it did not exceed 10–20% of the coating thickness to limit the influence of the substrate. The adhesion characteristics were studied on a Nanovea scratch tester, which represents a diamond cone with apex angle of 120 and radius of top curvature of 100 μm. The tests were conducted with the load linearly increasing from 0.05 to 40 N. Crack length was 5 mm. Each sample was subjected to three trials. The obtained curves were used to determine two parameters: the first critical load, LC1, at which the first cracks appeared in the coating, and the second critical load, LC2, which caused the total failure of the coating.

#### 2.3. Study of cutting properties

A study of the cutting properties of the tool made of carbide with developed NMCC was conducted using a lathe CU 500 MRD for longitudinal turning of steel C45 (HB 200). In the experiment, the cutters featured mechanical fastening of inserts made of carbide (WC + 15%


Note: ITi = current of titanium cathode, IAl = current of aluminum cathode, IZrNb = current of zirconium-niobium cathode, ICr = current of chromium cathode, pN = gas pressure in chamber, and U = voltage on substrate.

Table 1. Parameters of stages of the technological process of deposition of NMCC.

TiC + 6% Co) with square shapes (SNUN ISO 1832:2012) and with the following figures for the geometric parameters of the cutting part: γ = �8�, α = 6�, K = 45�, λ = 0, and R = 0.8 mm. The study was performed for the following cutting modes: f = 0.2 mm/rev, аp = 1.0 mm, and vc = 250 m min�<sup>1</sup> . Flank wear-land values (VBc) were measured with a toolmaker's microscope MBS-10 as the arithmetic mean of four to five tests. A value of VBc = 0.4 mm was taken as failure criterion. The study included statistical processing of tests of wear of cutting tools, sample mean value of wear, and sample mean square deviation of tool wear, which are random variables with different values in repeated experiments. Of note, during the experiments, outlying results were excluded. To exclude outlying results of the experiments, Irwin's criterion was used. To do that, the value of Irwin's criterion K<sup>λ</sup> was defined, if the outlying result was the maximum value VBmax:

$$\mathbf{K}\lambda = (VB\_{\mathcal{C}} - VB\_{\max})/\mathcal{K}\_{\mathcal{O}} \tag{9}$$

and if the doubts were provoked by the wear value with minimum value VBmin:

$$K = (VB\_{\mathcal{C}} - VB\_{\min})/K\_{\sigma} \tag{10}$$

NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N (Figure 4) and (NMCC Ti-TiN-(Ti,Al)N) (Figure 5) are characterized by a number of significant differences. The failures of those coatings occur under the mechanism of "wedging spallation." Meanwhile, NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N shows extensive interlayer delaminations, whereas in NMCC Ti-TiN-(Ti,Al)N, similar delaminations are less pronounced, and delaminations between nano-sublayers also occur. Generally, this picture correlates with the nature of the failure of those coatings observed during cutting tests.

substrate" delamination, and (4) the boundary of the coating zone, pressed by the tip of the scratch tester.

Figure 4. The nature of failure of NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N along a longitudinal crack, caused by a diamond indenter at critical (breaking) load. Vsc, scribing direction [28]. (1) The boundary of the wedging spallation zone, (2) the delamination boundary between the intermediate and wear-resistant layers, (3) the boundary of "adhesion coating layer-

Figure 3. The nature of failure of coating TiN along a longitudinal crack, caused by a diamond indenter at critical (breaking) load [28]. Vsc, scribing direction. (1) "Substrate-coating" boundary, (2) boundary of the brittle fracture zone of

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the coating, (3) boundary of the scribing groove, and (4) splintered section of the coating.

The calculated value K<sup>λ</sup> was compared to the critical value KλΑ, defined theoretically for a given level of significance level Α and selection criterion n. If K<sup>λ</sup> < KλΑ, then deviation of questionable value VBc was considered as valid.

#### 3. Results and discussion

#### 3.1. Adhesion characteristics

The classical test that enables determination of the strength of the adhesive bond of a coating with a substrate by the scratch-test method also can be used for qualitative evaluation of the strength of the adhesive bond between individual coating layers and cohesive bond between nano-sublayers. The tests were conducted on a Nanovea scratch tester. The indenter was a diamond cone with an apex angle of 120� and radius of top curvature of 100 μm. The tests were performed with a load linearly increasing from 0.05 N to the final load (40 N). Crack length was 5 mm. Each sample was subjected to three trials. The obtained curves were used to determine two parameters: the first critical load LC1, at which first cracks appeared in NMCC, and the second critical load LC2, which caused the total failure of NMCC. Typical types of failure are presented in Figure 3 (standard coating TiN), Figure 4 (NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N), and Figure 6 (NMCC Ti-TiN-(Ti,Al)N). All investigated coatings showed a sufficiently high level of adhesion bonds with substrate. Numerous research efforts and the experience of the authors of this paper show that a scratch test does not have a unique correlation with the tool life of a coated tool [1]; the test allows only "rejecting" coatings with insufficient strength of adhesion bonds. However, this test enables a study of the nature of the coating failure, particularly from the point of view of delaminations that occur in its structure. Let us consider the nature of the failure of the single-layer monolithic TiN coating (Figure 3). A fairly smooth scribing groove is clear, with a clearly visible area of brittle fracture of the outer area of the coating. On the edges of the groove, cracks and splintered areas of the coating are visible. Patterns of failure of

Delamination and Longitudinal Cracking in Multilayered Composite Nanostructured Coatings and Their Influence… http://dx.doi.org/10.5772/intechopen.72257 119

TiC + 6% Co) with square shapes (SNUN ISO 1832:2012) and with the following figures for the geometric parameters of the cutting part: γ = �8�, α = 6�, K = 45�, λ = 0, and R = 0.8 mm. The study was performed for the following cutting modes: f = 0.2 mm/rev, аp = 1.0 mm, and

scope MBS-10 as the arithmetic mean of four to five tests. A value of VBc = 0.4 mm was taken as failure criterion. The study included statistical processing of tests of wear of cutting tools, sample mean value of wear, and sample mean square deviation of tool wear, which are random variables with different values in repeated experiments. Of note, during the experiments, outlying results were excluded. To exclude outlying results of the experiments, Irwin's criterion was used. To do that, the value of Irwin's criterion K<sup>λ</sup> was defined, if the

The calculated value K<sup>λ</sup> was compared to the critical value KλΑ, defined theoretically for a given level of significance level Α and selection criterion n. If K<sup>λ</sup> < KλΑ, then deviation of

The classical test that enables determination of the strength of the adhesive bond of a coating with a substrate by the scratch-test method also can be used for qualitative evaluation of the strength of the adhesive bond between individual coating layers and cohesive bond between nano-sublayers. The tests were conducted on a Nanovea scratch tester. The indenter was a diamond cone with an apex angle of 120� and radius of top curvature of 100 μm. The tests were performed with a load linearly increasing from 0.05 N to the final load (40 N). Crack length was 5 mm. Each sample was subjected to three trials. The obtained curves were used to determine two parameters: the first critical load LC1, at which first cracks appeared in NMCC, and the second critical load LC2, which caused the total failure of NMCC. Typical types of failure are presented in Figure 3 (standard coating TiN), Figure 4 (NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N), and Figure 6 (NMCC Ti-TiN-(Ti,Al)N). All investigated coatings showed a sufficiently high level of adhesion bonds with substrate. Numerous research efforts and the experience of the authors of this paper show that a scratch test does not have a unique correlation with the tool life of a coated tool [1]; the test allows only "rejecting" coatings with insufficient strength of adhesion bonds. However, this test enables a study of the nature of the coating failure, particularly from the point of view of delaminations that occur in its structure. Let us consider the nature of the failure of the single-layer monolithic TiN coating (Figure 3). A fairly smooth scribing groove is clear, with a clearly visible area of brittle fracture of the outer area of the coating. On the edges of the groove, cracks and splintered areas of the coating are visible. Patterns of failure of

and if the doubts were provoked by the wear value with minimum value VBmin:

. Flank wear-land values (VBc) were measured with a toolmaker's micro-

K<sup>λ</sup> ¼ ð Þ VBc � VBmax =K<sup>σ</sup> (9)

K ¼ ð Þ VBc � VBmin =K<sup>σ</sup> (10)

vc = 250 m min�<sup>1</sup>

118 Novel Nanomaterials - Synthesis and Applications

outlying result was the maximum value VBmax:

questionable value VBc was considered as valid.

3. Results and discussion

3.1. Adhesion characteristics

Figure 3. The nature of failure of coating TiN along a longitudinal crack, caused by a diamond indenter at critical (breaking) load [28]. Vsc, scribing direction. (1) "Substrate-coating" boundary, (2) boundary of the brittle fracture zone of the coating, (3) boundary of the scribing groove, and (4) splintered section of the coating.

Figure 4. The nature of failure of NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N along a longitudinal crack, caused by a diamond indenter at critical (breaking) load. Vsc, scribing direction [28]. (1) The boundary of the wedging spallation zone, (2) the delamination boundary between the intermediate and wear-resistant layers, (3) the boundary of "adhesion coating layersubstrate" delamination, and (4) the boundary of the coating zone, pressed by the tip of the scratch tester.

NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N (Figure 4) and (NMCC Ti-TiN-(Ti,Al)N) (Figure 5) are characterized by a number of significant differences. The failures of those coatings occur under the mechanism of "wedging spallation." Meanwhile, NMCC Zr-ZrN-(Nb,Zr,Ti,Al)N shows extensive interlayer delaminations, whereas in NMCC Ti-TiN-(Ti,Al)N, similar delaminations are less pronounced, and delaminations between nano-sublayers also occur. Generally, this picture correlates with the nature of the failure of those coatings observed during cutting tests.

Figure 5. The nature of failure of NMCC Ti-TiN-(Ti,Al)N along a longitudinal crack, caused by a diamond indenter at critical (breaking) load [28]. Vsc, scribing direction.

The study of the scribing process for nanostructured coatings of large thickness (exceeding 10 μm) is of particular interest. In this case, it is possible to observe both coating failure caused by violation of adhesion bonds between layers and cohesive bonds between nano-sublayers and failure of a coating as a whole, when failure is not accompanied by delamination. Signs of failure of coating Ti-TiN-(Ti,Al)N (with coating thickness 13 μm) at scribing are shown in Figures 6 and 7.

3.2. Determination of basic properties of NMCC under cutting tests

the following conditions:

favored the formation of the NMCC.

a diamond indenter at critical (breaking) load [28].

mental data are shown in Figure 8.

[10–16].

This study was focused on the NMCC containing nitrides of Ti, Al, Cr, Zr, and Nb in its composition. For the detailed studies of various properties, NMCC were selected based on

Figure 7. The nature of failure of NMCC Ti-TiN-(Ti,Al)N (coating thickness 13 μm) along a longitudinal crack, caused by

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• If earlier studies show significant increase in cutting properties and reliability of the tool

• If the thermodynamic criterion ΔrG (Gibbs free energy change per mole of reaction)

To accomplish the research tasks, NMCC of various compositions were selected to meet the above conditions and were deposited using the FCVAD technology. The thicknesses of the coatings used in the studies were 2.4–5.0 μm. A wide range of thicknesses were selected on the basis of previous studies (in particular [10–16]), indicating the improvement in cutting performance with increase in coating thickness. The basic properties of the NMCC under study are presented in Table 2. Curves obtained by mathematical processing of the experi-

The NMCC Ti-TiN-(Ti,Al)N shows better resistance for approximately 19 min of operation due to its high surface hardness; however, subsequently, the tool with such a coating begins to experience intensive wear. This fact can be related to the start of intense cracking and wear of this coating. As a result, the tool with NMCC Zr-ZrN-(Zr,Cr,Al)N showed better resistance, and it was characterized by a balanced combination of sufficiently high hardness and resistance to brittle fracture. Let us consider in detail the mechanism of cracking and failure of coatings, paying special attention to such aspects of those processes as longitudinal cracks and

In particular, Figure 6 shows both violation of the interlayer interface between layers TiN and (Ti,Al)N and persistence of strong adhesion bonds between nano-sublayers of layer (Ti,Al)N. At zoom in, it is possible to notice in Figure 7 that in some cases, at critical loads, there is also failure of cohesive bonds between nano-sublayers, and that fact results in formation of a kind of "terraces," i.e., flat microsites with surface structure of a nano-sublayer. It is also possible to see signs of a tear-out of microdroplets embedded in the coating structure. Figure 7 shows the "terrace-like" structure of failure zone of a nanostructured coating. A general structure of the coating under the study and the nature of cracking in it during the cutting tests are shown below, in Figure 20.

Figure 6. The nature of failure of NMCC Ti-TiN-(Ti,Al)N (coating thickness 13 μm) along a longitudinal crack, caused by a diamond indenter at critical (breaking) load [28].

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Figure 7. The nature of failure of NMCC Ti-TiN-(Ti,Al)N (coating thickness 13 μm) along a longitudinal crack, caused by a diamond indenter at critical (breaking) load [28].

#### 3.2. Determination of basic properties of NMCC under cutting tests

The study of the scribing process for nanostructured coatings of large thickness (exceeding 10 μm) is of particular interest. In this case, it is possible to observe both coating failure caused by violation of adhesion bonds between layers and cohesive bonds between nano-sublayers and failure of a coating as a whole, when failure is not accompanied by delamination. Signs of failure of coating Ti-TiN-(Ti,Al)N (with coating thickness 13 μm) at scribing are shown in

Figure 5. The nature of failure of NMCC Ti-TiN-(Ti,Al)N along a longitudinal crack, caused by a diamond indenter at

In particular, Figure 6 shows both violation of the interlayer interface between layers TiN and (Ti,Al)N and persistence of strong adhesion bonds between nano-sublayers of layer (Ti,Al)N. At zoom in, it is possible to notice in Figure 7 that in some cases, at critical loads, there is also failure of cohesive bonds between nano-sublayers, and that fact results in formation of a kind of "terraces," i.e., flat microsites with surface structure of a nano-sublayer. It is also possible to see signs of a tear-out of microdroplets embedded in the coating structure. Figure 7 shows the "terrace-like" structure of failure zone of a nanostructured coating. A general structure of the coating under the study and the nature of cracking in it during the cutting tests are shown below,

Figure 6. The nature of failure of NMCC Ti-TiN-(Ti,Al)N (coating thickness 13 μm) along a longitudinal crack, caused by

Figures 6 and 7.

critical (breaking) load [28]. Vsc, scribing direction.

120 Novel Nanomaterials - Synthesis and Applications

a diamond indenter at critical (breaking) load [28].

in Figure 20.

This study was focused on the NMCC containing nitrides of Ti, Al, Cr, Zr, and Nb in its composition. For the detailed studies of various properties, NMCC were selected based on the following conditions:


To accomplish the research tasks, NMCC of various compositions were selected to meet the above conditions and were deposited using the FCVAD technology. The thicknesses of the coatings used in the studies were 2.4–5.0 μm. A wide range of thicknesses were selected on the basis of previous studies (in particular [10–16]), indicating the improvement in cutting performance with increase in coating thickness. The basic properties of the NMCC under study are presented in Table 2. Curves obtained by mathematical processing of the experimental data are shown in Figure 8.

The NMCC Ti-TiN-(Ti,Al)N shows better resistance for approximately 19 min of operation due to its high surface hardness; however, subsequently, the tool with such a coating begins to experience intensive wear. This fact can be related to the start of intense cracking and wear of this coating. As a result, the tool with NMCC Zr-ZrN-(Zr,Cr,Al)N showed better resistance, and it was characterized by a balanced combination of sufficiently high hardness and resistance to brittle fracture. Let us consider in detail the mechanism of cracking and failure of coatings, paying special attention to such aspects of those processes as longitudinal cracks and


Table 2. The basic properties of NMCC and periods of tool life of the carbide tools under study with the NMCC under study.

Figure 8. Dependence of wear VB on cutting time for dry turning of steel C45 at ap = 1.0 mm, f = 0.2 mm/rev, and vc = 250 m/min. (1) Uncoated, (2) TiN, (3) Zr-ZrN-(Zr,Cr,Al,Nb)N, (4) NMCC Zr-ZrN-(Zr,Cr,Al)N, (5) NMCC Ti-TiN-(Ti, Al)N, and (6) Zr-ZrN-(Zr,Cr,Al)N.

delaminations (interlayer delaminations and delaminations between nano-sublayers) form. Basic mechanism for the formation of longitudinal cracks and delaminations can be distinguished in a nanostructured multilayered coating because of the tearing force related to the adhesion interaction between the outer boundary of the coating and the material being machined (Figure 9), which has a prevailing fatigue characteristic and results in the formation of fatigue cracks due to the alternating processes of formation and failure of adhesion bridges in the system of "coating-material being machined." The considered mechanism is more typical for coatings on the rake face of the tool.

An important distinctive feature of the development of longitudinal cracks in nanostructured coatings is the formation of bridges in the process of cracking due to the alternation of less plastic sublayers with more plastic ones in the coating structure. Such bridges inhibit the development of a crack by exerting a positive influence on coating crack resistance and, consequently, on the tool life of a cutting tool (Figure 11). This mechanism of inhibition of

Figure 10. An example of the failure of the upper layer of NMCC Zr-ZrN-(Zr,Cr,Al)N because of the tearing force related to the adhesion interaction between the outer boundary of the coating and the material being machined (steel C45) [28].

Figure 9. The mechanism of formation of longitudinal cracks and delaminations in a nanostructured multilayered coating during cutting due to the tearing force associated with adhesion interaction between the outer boundary of the

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coating and the material being machined [28].

The action of the mechanism shown in Figure 9 can result not only in the formation of longitudinal cracks and delaminations but also in the destruction of the surface layers of the coating and, consequently, in the deterioration of the tool life of the metal-cutting tool (Figure 10).

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Figure 9. The mechanism of formation of longitudinal cracks and delaminations in a nanostructured multilayered coating during cutting due to the tearing force associated with adhesion interaction between the outer boundary of the coating and the material being machined [28].

delaminations (interlayer delaminations and delaminations between nano-sublayers) form. Basic mechanism for the formation of longitudinal cracks and delaminations can be distinguished in a nanostructured multilayered coating because of the tearing force related to the adhesion interaction between the outer boundary of the coating and the material being machined (Figure 9), which has a prevailing fatigue characteristic and results in the formation of fatigue cracks due to the alternating processes of formation and failure of adhesion bridges in the system of "coating-material being machined." The considered mechanism is more

Figure 8. Dependence of wear VB on cutting time for dry turning of steel C45 at ap = 1.0 mm, f = 0.2 mm/rev, and vc = 250 m/min. (1) Uncoated, (2) TiN, (3) Zr-ZrN-(Zr,Cr,Al,Nb)N, (4) NMCC Zr-ZrN-(Zr,Cr,Al)N, (5) NMCC Ti-TiN-(Ti,

The action of the mechanism shown in Figure 9 can result not only in the formation of longitudinal cracks and delaminations but also in the destruction of the surface layers of the coating and, consequently, in the deterioration of the tool life of the metal-cutting tool (Figure 10).

typical for coatings on the rake face of the tool.

Al)N, and (6) Zr-ZrN-(Zr,Cr,Al)N.
