**Synthesis, Luminescence and Magnetic Properties of Novel Fe0.5Gd0.5(MoO4)1.5:Eu3+ Micro/Nano (3D) Structures**

Rajagopalan Krishnan and Jagannathan Thirumalai

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/61323

### **Abstract**

[70] Kawanishi S, Hiraku Y, Murata M, Oikawa S. The role of metals in site-specific DNA damage with reference to carcinogenesis. *Free Rad Biol Med*. 2002;32:822-832. DOI:

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O2 within the mitochondrial matrix space. Arch Bio‐

[74] Petrat F, Pindiur S, Kirsch M, de Groot H. "Mitochondrial" photochemical drugs do

O2 in mito‐

[73] Petrat, F, Pindiur S, Kirsch M, de Groot H. NAD(P)H, a primary target of 1

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10.1111/j.1751-1097.1974.tb06506.x

DOI: 10.1021/ja00758a072

not release toxic amounts of 1

jbc.M204230200

32 Nanoparticles Technology

The controlled synthesis of Fe0.5Gd0.5(MoO4)1.5:Eu3+ microcrystals (self-aggregated 3D su‐ perstructures) using hydrothermal method by employing PVP as surfactant along with the state of the art of this field at present. By varying the reaction time, we have reported the morphology selection and the condition to derive novel nanoparticle sheathed bipyr‐ amid-like morphology. The rest of the parameters like, molar ratio between initial precur‐ sor / surfactant and temperature were kept as constant. The Fe0.5Gd0.5(MoO4)1.5:Eu3+ microcrystals are found to be an excellent matrix for photoluminescence property and the Eu3+ ion serves as a good red emitting luminescent center. Further, room temperature (RT) magnetic properties of Fe0.5Gd0.5(MoO4)1.5:Eu3+ were investigated in detail.

**Keywords:** Chemical synthesis, Electron diffraction, Crystal structure, Optical properties, Magnetic measurements

### **1. Introduction**

Size- and shape-controlled synthesis of self-organized, three-dimensional (3D) micro/nano architectures are still challenging and much attracted in the recent years owing to vast scientific and technological interest in multifaceted research areas [11, 27, 14]. The growth and devel‐ opment of uniform, self-aggregated 3D super-structures through plane-to-plane coalition of the basic building blocks have been paid significant attention owing to their unusual physical, chemical properties [29, 3]. The development of self-organized architectures consents to find neoteric and diversified materials, causing momentous discovery that has been developed for the construction of a new class of micro/nanostructured devices [14, 15]. Self-aggregation of hierarchical architectures has proved to be a virtuous and expedient tool to control the

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structure and properties for the fabrication of intricate networks [15]. The peculiar size and morphology of the self-assembled 3D networks determine the attributes of a material which are important for the manufacture of a novel category of micro/nano-structured devices for advanced engineering and sciences [29, 14, 13]. The supreme and hopeful approach to synthesize self-assembled 3D networks is to use a solution-based technique which is a simple route to regulate the physicochemical properties of materials [3]. Specifically, in the hydro‐ thermal method, numerous versatile morphologies have been described by using surfactants or chelating agents [7, 16, 17]. The surface capping agent used in the hydrothermal synthesis method positively assures an effective way to regulate the nucleation, controlled growth, oriented aggregation mechanism, and self-organized 3D morphology [15, 4].

### **2. Recent research scenario**

A wide variety of 3D super-structures can be synthesized by using a suitable amount of surfactant which reduces the particle–particle amalgamation by steric effects, thereby control‐ ling the size and shape [15]. Up till now, a number of research works have been reported on controlled synthesis of 3D hierarchical architectures using hydrothermal technique, for example, bowknot-like Y2(WO4)3:Ln3+ [9], almond-like (Na0.5La0.5)MoO4:RE3+ [14], monodis‐ perse CaMoO4:Eu3+, M+ (M = Li, Na, K) microspheres [18], micron-sized flower-like NaCe(MoO4)2 architectures [27], tetragonal bipyramid-like (Na0.5Gd0.5)MoO4:Ln3+ [15], rugbylike NaEu(MoO4)2 microstructures [26], sheaf-like Gd2(MoO4)3 [24], etc. Large-scale assert for phosphor materials as efficient sources of energy that can supply sustained proficiency is growing gradually. Current research in phosphors is focused on new materials having both luminescence and magnetic properties in a single entity based on the presence of trivalent lanthanide ions (Ln3+) with magnetic Fe3+ ions. These bifunctional materials have attracted worldwide intense research, because their up/down conversion luminescence and para/ ferromagnetic properties are strongly required for therapeutics, contrast agents in magnetic resonance (MRI), fluorescent bioimaging in biophotonics, drug delivery and lumino-magnetic applications [6, 13, 19]. The phosphor particles can be magnetically directed, aligned, tracked, after applying the external magnetic fields. Subsequently, their luminescence can be visualized under optical excitation, which can be used widely for biological labeling, imaging of human and animal cells, tissues, and in vitro, in vivo applications [22, 25].


**Table 1.** Recent investigation of magnetic and luminescence properties of various phosphors.

structure and properties for the fabrication of intricate networks [15]. The peculiar size and morphology of the self-assembled 3D networks determine the attributes of a material which are important for the manufacture of a novel category of micro/nano-structured devices for advanced engineering and sciences [29, 14, 13]. The supreme and hopeful approach to synthesize self-assembled 3D networks is to use a solution-based technique which is a simple route to regulate the physicochemical properties of materials [3]. Specifically, in the hydro‐ thermal method, numerous versatile morphologies have been described by using surfactants or chelating agents [7, 16, 17]. The surface capping agent used in the hydrothermal synthesis method positively assures an effective way to regulate the nucleation, controlled growth,

A wide variety of 3D super-structures can be synthesized by using a suitable amount of surfactant which reduces the particle–particle amalgamation by steric effects, thereby control‐ ling the size and shape [15]. Up till now, a number of research works have been reported on controlled synthesis of 3D hierarchical architectures using hydrothermal technique, for example, bowknot-like Y2(WO4)3:Ln3+ [9], almond-like (Na0.5La0.5)MoO4:RE3+ [14], monodis‐

NaCe(MoO4)2 architectures [27], tetragonal bipyramid-like (Na0.5Gd0.5)MoO4:Ln3+ [15], rugbylike NaEu(MoO4)2 microstructures [26], sheaf-like Gd2(MoO4)3 [24], etc. Large-scale assert for phosphor materials as efficient sources of energy that can supply sustained proficiency is growing gradually. Current research in phosphors is focused on new materials having both luminescence and magnetic properties in a single entity based on the presence of trivalent lanthanide ions (Ln3+) with magnetic Fe3+ ions. These bifunctional materials have attracted worldwide intense research, because their up/down conversion luminescence and para/ ferromagnetic properties are strongly required for therapeutics, contrast agents in magnetic resonance (MRI), fluorescent bioimaging in biophotonics, drug delivery and lumino-magnetic applications [6, 13, 19]. The phosphor particles can be magnetically directed, aligned, tracked, after applying the external magnetic fields. Subsequently, their luminescence can be visualized under optical excitation, which can be used widely for biological labeling, imaging of human

(M = Li, Na, K) microspheres [18], micron-sized flower-like

**Observed magnetism Reference**

oriented aggregation mechanism, and self-organized 3D morphology [15, 4].

and animal cells, tissues, and in vitro, in vivo applications [22, 25].

**investigated**

**Table 1.** Recent investigation of magnetic and luminescence properties of various phosphors.

Gd2O2S:Yb/Er Upconversion Paramagnetic [1] NaGdF4:Yb3+, Ln3+ Upconversion Paramagnetic [8] YAG:Ce3+ Downconversion Ferromagnetism [10] (Na0.5La0.5)MoO4:Eu3+ Up/down conversion Ferromagnetism [15] (Na0.5Gd0.5)MoO4:Eu3+ Up/down conversion Paramagnetic [14]

**Phosphor Type of luminescence**

**2. Recent research scenario**

34 Nanoparticles Technology

perse CaMoO4:Eu3+, M+

The luminescence and magnetic properties of lanthanide ions doped iron-based inorganic compounds that are synthesized by different research groups are detailed as follows. The novel Fe3O4@YPO4:Re (Re = Tb, Eu) magnetic-fluorescent hybrid spheres were prepared using solvothermal and co-precipitation techniques [25]. The obtained nanostructures are an excellent platform to integrate fluorescent materials and magnetite into one single entity to exhibit both magnetic and fluorescent properties [25]. Wang et al. [25] reported that human cervical carcinoma Hela cells were successfully labeled by using the Fe3O4@YPO4:Re hybrid spheres. Zhang et al. [29] synthesized orthorhombic cubic GdFeO3 particles by using a facile hydrothermal method. The emission color of cubic GdFeO3 particles could be adjusted effectively by doping different lanthanide activators [29]. Also, cubic GdFeO3 particles exhibiting paramagnetic properties are reported to be used as an excellent luminescence and magnetic material [29]. To replace the existing halide-based upconverting phosphors for infrared-based biomedical imaging [1], nontoxic Gd2O2S:Yb/Er is synthesized using flux fusion method. The synthesized phosphor not only enables near-IR imaging features but also at the same time could be used as a contrast agent in MRI imaging [1]. Multifunctional luminescentmagnetic YAG:Ce nano-phosphors were synthesized by post-heat solidified combustion technique, post-heated hydrothermal-homogenous precipitation, and post-heated auto-clave techniques, respectively, using inexpensive aluminum and yttrium nitrates as the starting materials and urea as the homogenizing precipitant [10]. Since the yttrium and aluminium are non-magnetic elements, the origin of ferromagnetism of these nonmagnetic oxides could be attributed solely to the defects and/or oxygen vacancies [10]. The novel YAG:Ce nanophosphor exhibits the ferromagnetic behavior which finds the way for a new class of magneto-optic and spintronic materials for a wide variety of applications. The multicolor upconversion emissions and paramagnetic nature were observed in β-NaGdF<sup>4</sup> crystals, which were synthesized by one-step precipitation method at room temperature [8]. To control the crystal size and morphology of the β-NaGdF4 crystals, different types of surfactants, namely Na2EDTA, PVP, SDS, and Na2tar, were introduced during the synthesis procedure. Due to the presence of paramagnetic Gd3+ ions, β-NaGdF<sup>4</sup> crystals exhibit paramagnetic behavior, which is a favora‐ ble bifunctional candidate for bioimaging applications [8].

Scheelite-type metal molybdates are considered as an ideal candidate, and are receiving accelerating interest due to their remarkable luminescence, magnetic, catalytic, sensing, and photocatalytic properties [13]. Scheelite-type crystalline structures, with the general chemical formula such as (Na0.5R0.5)MoO4:Ln3+ [15, 14]; BaMoO4:Ln3+ [23; Ca0.5R(MoO4)2:Ln3+ [20]; Fe0.5R0.5(MoO4)1.5:Ln3+ [13] (where R = La or Gd or Y and Ln = Eu, Tb, Dy, etc.) etc., where molybdenum atoms are coordinated tetrahedrally, have recently attracted great attention because of their promising applications in the opto-electronics [14, 15]. In the crystal structure of Fe0.5R0.5(MoO4)1.5:Ln3+, trivalent Fe3+, R3+/Ln3+ jointly occupy the dodecahedral positions and molybdenum (Mo6+) atoms populated at the center of the tetrahedral symmetry coordinated with four equivalent oxygen atoms. Scheelite-type metal molybdates with the general formula have been synthesized by various techniques, which include hydrothermal method [14, 15], solid-state reaction technique [12], hydrothermal-assisted solid-state reaction [28], and pulsed laser deposition technique [21]. The random distribution of R3+ (La or Gd or Y) in the scheelite tetragonal structures show an intense and broad absorption band in the near UV- region [15]. Among all these molybdates, (Fe0.5Gd0.5)MoO4:Eu3+ is an important opto-magnetic material and possess good luminescence property. Due to the notable emission spectral lines in Ln3+ doped phosphors extending from ultraviolet to infra-red regionand its magnetic properties may open up a new platform to design and formulate other inorganic functional materials with for potential magnetic and luminescence applications [13], there is a great deal of scientific and technological interest in developing novel opto-electronic applications. In general, synthesis of any kind of phosphor materials using the conventional/sol–gel technique would require laborious conditions such as hefty manual grinding, annealing, or calcination of samples to high temperature, intermediate grinding process, and heat treatment for several times, etc. In order to avoid the above mentioned difficulties, hydrothermal technique is the alternative and most promising approach to synthesize the variety of phosphor materials with desired size, shape, and dimensions, crystallinity by changing the experimental parameters.

Herein, we report the controlled synthesis of (Fe0.5Gd0.5)MoO4:Eu3+ microcrystals (selfaggregated 3D superstructures) using hydrothermal method by employing PVP as the surfactant. By varying the reaction time, we have reported the morphology selection and the condition to derive novel nanoparticle sheathed bipyramid-like morphology. The rest of the parameters like molar ratio between initial precursor / surfactant and temperature were kept as constant. The (Fe0.5Gd0.5)MoO4:Eu3+ microcrystals are found to be excellent matrix for photoluminescence property and the Eu3+ ion serves as a good red-emitting luminescent center. Further, room-temperature (RT) magnetic properties of Fe0.5Gd0.5(MoO4)1.5:Eu3+ were investi‐ gated in detail.

### **3. Synthesis, morphology, structure of the Fe0.5Gd0.5(MoO4)1.5:Eu3+ 3D hierarchical microstructures**

### **3.1. Synthesis procedure**

All the chemicals were purchased from Sigma Aldrich with 99.99% purity. In the typical synthesis, a stoichiometric amount of FeCl3 6H2O was dissolved in 15 mL of double-distilled water through vigorous stirring. Then, the appropriate amount of Na2MoO4 was initially dissolved in 30 mL of double- distilled water under vigorous stirring. Then, GdCl3 and EuCl3 were prepared by dissolving the corresponding rare-earth sesquioxides (Gd2O3, Eu2O3) in diluted hydrochloric acid and stirred for 15 min. In order to remove the excess of HCl, the solution was heated to 60–80°C for 30 min. Then, the rare-earth chloride solution was carefully mixed into the Na2MoO4 solution. A white colloidal precipitate was obtained under vigorous stirring for 15 min. Subsequently, iron chloride solution was gradually introduced into the above white precipitate solution, forming a brown mixture solution. Then, 0.1 mM of polyvi‐ nylpyrrolidone (PVP40) was dissolved in 20 mL of double-distilled water, and was added to the resultant colloidal solution. Then, the pH value of the final product was consequently, adjusted to a value of 7 -8 by adding the proper amount of sodium hydroxide solution. After additional stirring for 30 min, the acquired brown mixture solution was transferred into a 100 mL Teflon-lined stainless steel autoclave, and heated at 200°C for 6, 12, and 24 h. Finally, the autoclave was allowed to cool at room temperature, and the final product was centrifugally cleaned (5, 000 rpm for 30 min to remove organic containments and NaCl), separated from the solution, rinsed with distilled water and absolute ethanol several times, and dried at 60°C in the air for 5 h.

### **3.2. Characterization**

Among all these molybdates, (Fe0.5Gd0.5)MoO4:Eu3+ is an important opto-magnetic material and possess good luminescence property. Due to the notable emission spectral lines in Ln3+ doped phosphors extending from ultraviolet to infra-red regionand its magnetic properties may open up a new platform to design and formulate other inorganic functional materials with for potential magnetic and luminescence applications [13], there is a great deal of scientific and technological interest in developing novel opto-electronic applications. In general, synthesis of any kind of phosphor materials using the conventional/sol–gel technique would require laborious conditions such as hefty manual grinding, annealing, or calcination of samples to high temperature, intermediate grinding process, and heat treatment for several times, etc. In order to avoid the above mentioned difficulties, hydrothermal technique is the alternative and most promising approach to synthesize the variety of phosphor materials with desired size,

Herein, we report the controlled synthesis of (Fe0.5Gd0.5)MoO4:Eu3+ microcrystals (selfaggregated 3D superstructures) using hydrothermal method by employing PVP as the surfactant. By varying the reaction time, we have reported the morphology selection and the condition to derive novel nanoparticle sheathed bipyramid-like morphology. The rest of the parameters like molar ratio between initial precursor / surfactant and temperature were kept as constant. The (Fe0.5Gd0.5)MoO4:Eu3+ microcrystals are found to be excellent matrix for photoluminescence property and the Eu3+ ion serves as a good red-emitting luminescent center. Further, room-temperature (RT) magnetic properties of Fe0.5Gd0.5(MoO4)1.5:Eu3+ were investi‐

**3. Synthesis, morphology, structure of the Fe0.5Gd0.5(MoO4)1.5:Eu3+ 3D**

All the chemicals were purchased from Sigma Aldrich with 99.99% purity. In the typical synthesis, a stoichiometric amount of FeCl3 6H2O was dissolved in 15 mL of double-distilled water through vigorous stirring. Then, the appropriate amount of Na2MoO4 was initially dissolved in 30 mL of double- distilled water under vigorous stirring. Then, GdCl3 and EuCl3 were prepared by dissolving the corresponding rare-earth sesquioxides (Gd2O3, Eu2O3) in diluted hydrochloric acid and stirred for 15 min. In order to remove the excess of HCl, the solution was heated to 60–80°C for 30 min. Then, the rare-earth chloride solution was carefully mixed into the Na2MoO4 solution. A white colloidal precipitate was obtained under vigorous stirring for 15 min. Subsequently, iron chloride solution was gradually introduced into the above white precipitate solution, forming a brown mixture solution. Then, 0.1 mM of polyvi‐ nylpyrrolidone (PVP40) was dissolved in 20 mL of double-distilled water, and was added to the resultant colloidal solution. Then, the pH value of the final product was consequently, adjusted to a value of 7 -8 by adding the proper amount of sodium hydroxide solution. After additional stirring for 30 min, the acquired brown mixture solution was transferred into a 100 mL Teflon-lined stainless steel autoclave, and heated at 200°C for 6, 12, and 24 h. Finally, the autoclave was allowed to cool at room temperature, and the final product was centrifugally

shape, and dimensions, crystallinity by changing the experimental parameters.

gated in detail.

36 Nanoparticles Technology

**hierarchical microstructures**

**3.1. Synthesis procedure**

XRD pattern of as-synthesized phosphor was analyzed using PANalytical's X'Pert PRO materials research, X-ray diffractometer equipped with a Cu-Kα radiation (*λ* = 1.54060 A°) at a scanning rate of 0.02° s-1 performed in the range of 20–60 degrees. Morphology and energydispersive X-ray spectra (EDX) of the product were investigated by scanning electron micro‐ scope (FESEM (Quanta 3D FEG)). The spacing between the two adjacent lattice planes was measured using high-resolution transmission electron microscope (HRTEM JEOL 3010). Furthermore, the room-temperature (RT) photoluminescence studies were carried out using Horiba-Jobin Yvon Fluromax-4P bench-top Spectrofluorometer. Room-temperature magnetic properties of the samples were analyzed using (Lake Shore 7307 model) vibrating sample magnetometer (VSM).

### **3.3. Morphological investigations**

To explore the 3D surface morphology, particle size distribution, and formation mechanism for the evolution of nanoparticle sheathed bipyramid-like structures, a series of time-depend‐ ent experiments were carried out using a PVP employed hydrothermal route.

**Figure 1.** Field emission scanning electron microscopic images of Fe0.5Gd0.5(MoO4)1.5:Eu3+ synthesized by PVP employed hydrothermal method with different time intervals, i.e., (a) 1 h, (b) 2 h, (c) 6 h, (d) 12 h at 200°C with fixed pH (~7).

The FESEM and TEM analyses were used to characterize the morphology of the final prod‐ ucts. By changing the experimental reaction time interval with fixed temperature (200°C), pH (~7), and molar concentration of the PVP (0.1 mM), a substantial difference was observed in the final product. Figure (1–3)shows the low, high magnifications FESEM image of the Fea0.5Gd0.5(MoO4)1.5:Eu3+ samples synthesized by hydrothermal route at different time inter‐ vals (1, 3, 6, 12, 24, and 48 h) with fixed temperature (200°C), pH (~7), and PVP molar concentra‐ tion. During the initial period of 1 h time interval, numerous 2D nanoflakes with few tens of nanometer were observed, which acts as a primary building block for the subsequent growth of 3D microstructures (Figure 1 (a)). It was noticed from Figure 1 (b, c), when the hydrother‐ mal time period was increased to 3 and 6 h, nanoflakes began to self-organize continuously through plane-to-plane coalition and try to form a bipyramidal morphology (average diame‐ ter of 1 μm). However, the morphology and size distribution was found to be less uniform. When the reaction time interval was kept at 12 h, bipyramid-like morphology with an aver‐ age of 1.0 μm in length and an average of 1.2 μm in diameter was obtained and shown in Fig. 1 (d). When the reaction time period was raised to 24 h, highly ordered and self-assembled porous nanoparticle sheathed bipyramid-like morphology with the size approximately equal to ~1.2 μm diameter was observed (Figure 2 (a, b)). Further, increasing the reaction time to 48 h, the final product morphology consisted of irregular porous bipyramid-like structures in which nanoparticles are sheathed and shown in Figure 2 (c, d). This may be due to the prolonged time interval, resulting in thermal collision between the particles (growth stage; the self-assembly process predominates between these crystals because of its high surface energies) [14].

**Figure 2.** Low and high magnification SEM images of self-assembled nanoparticle sheathed Fe0Gd0.5(MoO4)1.5:Eu3+ structures synthesized at (a, b) 24 and (c, d) 48 h.

The FESEM and TEM analyses were used to characterize the morphology of the final prod‐ ucts. By changing the experimental reaction time interval with fixed temperature (200°C), pH (~7), and molar concentration of the PVP (0.1 mM), a substantial difference was observed in the final product. Figure (1–3)shows the low, high magnifications FESEM image of the Fea0.5Gd0.5(MoO4)1.5:Eu3+ samples synthesized by hydrothermal route at different time inter‐ vals (1, 3, 6, 12, 24, and 48 h) with fixed temperature (200°C), pH (~7), and PVP molar concentra‐ tion. During the initial period of 1 h time interval, numerous 2D nanoflakes with few tens of nanometer were observed, which acts as a primary building block for the subsequent growth of 3D microstructures (Figure 1 (a)). It was noticed from Figure 1 (b, c), when the hydrother‐ mal time period was increased to 3 and 6 h, nanoflakes began to self-organize continuously through plane-to-plane coalition and try to form a bipyramidal morphology (average diame‐ ter of 1 μm). However, the morphology and size distribution was found to be less uniform. When the reaction time interval was kept at 12 h, bipyramid-like morphology with an aver‐ age of 1.0 μm in length and an average of 1.2 μm in diameter was obtained and shown in Fig. 1 (d). When the reaction time period was raised to 24 h, highly ordered and self-assembled porous nanoparticle sheathed bipyramid-like morphology with the size approximately equal to ~1.2 μm diameter was observed (Figure 2 (a, b)). Further, increasing the reaction time to 48 h, the final product morphology consisted of irregular porous bipyramid-like structures in which nanoparticles are sheathed and shown in Figure 2 (c, d). This may be due to the prolonged time interval, resulting in thermal collision between the particles (growth stage; the self-assembly process predominates between these crystals because of its high surface energies) [14].

**Figure 2.** Low and high magnification SEM images of self-assembled nanoparticle sheathed Fe0Gd0.5(MoO4)1.5:Eu3+

structures synthesized at (a, b) 24 and (c, d) 48 h.

38 Nanoparticles Technology

This involves growth from 0D primary particles to 2D nanosheets, followed by the formation of 3D hierarchical networks through an assembly–disassembly process [14]. The TEM image of the sample synthesized at 24 h time interval with 0.1 mM of PVP molar concentration is shown in Figure 3a and the corresponding inset shows the FFT pattern. The formation mechanism for the evolution of a nanoparticle-sheathed bipyramid-like Fe0.5Gd0.5(MoO4)1.5:Eu3+ structure can be described as follows. Primarily, smaller crystalline nuclei formed in a supersaturated solution starts the crystallization process for subsequent growth of self-organization. During the self-organization process, smaller nanocrystals were grown-up constantly to form bigger particles via an Ostwald ripening phenomenon at higher reaction time intervals, due to interactions between the amorphous nanoparticles, resulting in 2D nanoflakes [14]. Initially, PVP acts as a surface capping agent to form a stable complex with Gd3+/Fe3+ and then kinetically controls the reaction rate, germination of nuclei, the growth rate, and the oriented aggregation mechanism. When the reaction temperature gradually increases to higher temperature, PVP would decompose gradually and has no obvious effect on the possible growth of nuclei. When the time period was increased, these 2D nanoflakes joined each other along a mutual crystallographic orientation through plane-to-plane conjunction of neighboring particles [14]. In general, such amalgamation process between the atoms/ions in molecules/crystals can induce dipole moments, which can be associated with certain specific crystallographic orientations [13]. The interaction among dipole moments in nearby nanopar‐ ticles is highly probable and acts as a driving factor for the oriented attachment procedure by allocation of a common crystallographic facet. The growth rate in Fe0Gd0.5(MoO4)1.5:Eu3+ crystals is oriented preferentially along the [001] direction. Donnay-Harker et al. [5] reported that, for tetragonal structures, the surface free energy of the [001] face is larger than that of the [101] face [5, 2]. As a result, the faster growth rate in the [100], [010], and [001] faces in comparison with that in the [101] direction accelerates the formation of self-assembled structures.

**Figure 3.** (a) TEM image of the Fe0Gd0.5(MoO4)1.5:Eu3+ sample prepared at 24 h time interval and the corresponding in‐ set shows the FFT pattern, (b) HRTEM image (inset typical HRTEM image taken on the marked part).

The as-grown 2D nanoflakes are continuously self-assembled via a layer-by-layer stacking process to form 3D bipyramid-like structures. Then, the completely developed bipyramid particles are readily disassembled into smaller nanoparticles, which are finally sheathed on individual bipyramid-like particles. After aging for a longer time period, the disassembly of well-organized structures could increase the entropy of the system and thus could decrease the free energy of the whole system [16], resulting in the nanoparticles arbitrarily sheathed on individual bipyramid-like particles. The assembly–disassembly process in Fe0Gd0.5(MoO4)1.5:Eu3+ was described in our previous work [13] by using EDTA as a surface capping agent.

**Figure 4.** Schematic diagram illustrating the formation mechanism of nanoparticle sheathed Fe0Gd0.5(MoO4)1.5:Eu3+ 3D hierarchical structures.

### **3.4. Structural analysis**

The chemical phase purity and crystalline nature of the product were investigated using X-ray diffraction patterns. Figure 5 shows the XRD pattern of Fe0.5Gd0.5(MoO4)1.5:Eu3+ samples synthesized by a facile hydrothermal route at different time intervals (6, 12, 24, and 48 h). All the peaks are perfectly matched with the standard JCPDS card no. 25-0828 of (Na0.5Gd0.5)MoO4. It could be clearly observed that the XRD patterns of as-synthesized samples are highly crystalline, and the strongest intensity peak is observed at 2θ = 28.67 degrees analogous to (112) plane. The XRD pattern reveals that they belong to a tetragonal phase corresponding to the scheelite structure with the space group I41/a. In this crystal structure Fe3+, Gd3+/Eu3+ ions jointly occupy the dodecahedral sites of the tetrahedra. Molybdenum atoms (Mo6+) populated in the center of tetrahedral symmetry is surrounded by four equivalent oxygen (O2–) atoms [13] and forming isolated [MoO4] 2– group. Owing to their similar ionic radius and valance state, the Eu3+ ion successfully replaces the Gd3+ ion and does not change the lattice site of the host Fe0.5Gd0.5(MoO4)1.5:Eu3+ and the crystal phase is very much pure. Further, no other extra peaks of impurity were observed.

Figure 5 (a, b) shows the representative EDX analysis, taking single bipyramid-like particle and single nanoparticle sheathed on bipyramidal structure synthesized at 24 h time interval for Eu3+ doped Fe0.5Gd0.5(MoO4)1.5. The energy-dispersive X-ray spectrum was used to confirm the presence of iron (Fe3+), gadolinium (Gd3+), europium (Eu3+), molybdenum (Mo6+), and Synthesis, Luminescence and Magnetic Properties of Novel Fe0.5Gd0.5(MoO4)1.5:Eu3+ Micro/Nano (3D) Structures http://dx.doi.org/10.5772/61323 41

particles are readily disassembled into smaller nanoparticles, which are finally sheathed on individual bipyramid-like particles. After aging for a longer time period, the disassembly of well-organized structures could increase the entropy of the system and thus could decrease the free energy of the whole system [16], resulting in the nanoparticles arbitrarily sheathed on individual bipyramid-like particles. The assembly–disassembly process in Fe0Gd0.5(MoO4)1.5:Eu3+ was described in our previous work [13] by using EDTA as a surface

**Figure 4.** Schematic diagram illustrating the formation mechanism of nanoparticle sheathed Fe0Gd0.5(MoO4)1.5:Eu3+ 3D

The chemical phase purity and crystalline nature of the product were investigated using X-ray diffraction patterns. Figure 5 shows the XRD pattern of Fe0.5Gd0.5(MoO4)1.5:Eu3+ samples synthesized by a facile hydrothermal route at different time intervals (6, 12, 24, and 48 h). All the peaks are perfectly matched with the standard JCPDS card no. 25-0828 of (Na0.5Gd0.5)MoO4. It could be clearly observed that the XRD patterns of as-synthesized samples are highly crystalline, and the strongest intensity peak is observed at 2θ = 28.67 degrees analogous to (112) plane. The XRD pattern reveals that they belong to a tetragonal phase corresponding to the scheelite structure with the space group I41/a. In this crystal structure Fe3+, Gd3+/Eu3+ ions jointly occupy the dodecahedral sites of the tetrahedra. Molybdenum atoms (Mo6+) populated in the center of tetrahedral symmetry is surrounded by four equivalent

radius and valance state, the Eu3+ ion successfully replaces the Gd3+ ion and does not change the lattice site of the host Fe0.5Gd0.5(MoO4)1.5:Eu3+ and the crystal phase is very much pure.

Figure 5 (a, b) shows the representative EDX analysis, taking single bipyramid-like particle and single nanoparticle sheathed on bipyramidal structure synthesized at 24 h time interval for Eu3+ doped Fe0.5Gd0.5(MoO4)1.5. The energy-dispersive X-ray spectrum was used to confirm the presence of iron (Fe3+), gadolinium (Gd3+), europium (Eu3+), molybdenum (Mo6+), and

2– group. Owing to their similar ionic

capping agent.

40 Nanoparticles Technology

hierarchical structures.

**3.4. Structural analysis**

oxygen (O2–) atoms [13] and forming isolated [MoO4]

Further, no other extra peaks of impurity were observed.

**Figure 5.** Powder X-ray diffraction (XRD) patterns of the Fe0Gd0.5(MoO4)1.5:Eu3+ samples synthesized at different time intervals using the PVP-mediated hydrothermal method. replaces the Gd3+ ion and does not change the lattice site of the host Fe0.5Gd0.5(MoO4)1.5:Eu3+ and the crystal phase is very much pure. Further, no other extra peaks of impurity were observed.

EDX spectrum of single nanoparticle sheathed on a bipyramid-like particle. **Figure 6.** (a) EDX spectrum of single bipyramid-like Fe0.5Gd0.5(MoO4)1.5:Eu3+ particle, (b) EDX spectrum of single nano‐ particle sheathed on a bipyramid-like particle.

Figure. 6 (a) EDX spectrum of single bipyramid-like Fe0.5Gd0.5(MoO4)1.5:Eu3+ particle, (b)

oxygen (O2–) in the final product. The EDX spectrum shown in Figure 5 (a, b) confirms the presence of Fe, Gd, Eu, Mo, and O in the final product.

### **3.5. Optical and magnetic properties of Fe0.5Gd0.5(MoO4)1.5:Eu3+ structures**

**Figure 7.** Room temperature emission spectra of Eu3+ doped Fe0.5Gd0.5(MoO4)1.5 synthesized by the PVP-assisted hydro‐ thermal route.

Figure 7 shows the PL excitation and emission spectra of nanoparticle sheathed bipyramidlike microstructures of Fe0.5Gd0.5(MoO4)1.5:Eu3+. Emission (*λ*ex = 395 nm) spectra of the Fe0.5Gd0.5(MoO4)1.5 doped with Eu3+ ions were recorded within the range from 575 to 700 nm at room temperature. Upon 395 nm UV irradiation, the emission spectrum was governed by the hypersensitive red emission, showing a transition 5 D0→<sup>7</sup> F2 (due to electric dipole transition) higher than 5 D0→<sup>7</sup> F1 (magnetic dipole). In the emission spectrum, the transitions 5 D0→<sup>7</sup> FJ (where J = 1, 2, 3, 4) were observed due to intra-configurational f–f electronic transitions of Eu3+ ions. The presence of electric dipole transition confirmed that Eu3+ ions were located at sites without inversion symmetry (C3v symmetry). The other transitions 5 D0→<sup>7</sup> F3 and 5 D0→<sup>7</sup> F4 were comparatively very weak for all the samples (not shown in PL spectrum). However, the presence of magnetic iron (Fe3+) in the Fe0.5Gd0.5(MoO4)1.5 tetragonal crystal structure consid‐ erably decreases their luminescence, as Fe3+ ions share the dodecahedral sites with Gd3+/Eu3+ [13]. However, it is still unclear for the weak red emission in Eu3+ doped Fe0.5Gd0.5(MoO4)1.5.

Figure 8 shows the room-temperature (RT) magnetic properties of hierarchically self-organ‐ ized Fe0.5Gd0.5(MoO4)1.5:Eu3+ 3D microstructures synthesized by the hydrothermal route at 24 h time interval. Magnetization as a function of applied magnetic field (from 15000 Oe to 15000 Oe) was measured using a vibrating sample magnetometer. Due to the presence of Gd3+ ions, a magnetic hysteresis (M–H) loop was not observed in the case of Eu3+ doped Fe0.5Gd0.5(MoO4)1.5. From Figure 8, it is clear that a straight line crosses the origin, indicating that the phosphor Fe0.5Gd0.5(MoO4)1.5:Eu3+ exhibits a paramagnetic behavior due to the existence of paramagnetic Gd3+ ions and the saturation magnetization (Ms) value is found to be 0.028 emu/gm for RT. The extremely localized nature of the seven unpaired inner 4f electrons of Gd plays an important role which determines its magnetic properties [15]. These unpaired electrons in the outer orbital are closely bound to the nucleus and effectively shielded by the outer closed shell electrons 5s2 5p6 from the crystal field [15]. Even though the Gd3+ ions jointly occupy the dodecahedral sites with Fe3+ ions, number of 4f electrons in the Gd3+ ions, atomic unit cell volume, direct f–f exchange interactions between neighboring Gd3+ and Fe3+ atoms are responsible for setting up the magnetic moments [15, 14]. Since, Gd3+ ions have total orbital angular momentum *L* = 0 [13], the spin orbit coupling between the partially filled 4f electrons in the Gd3+ and Fe3+ ions is weak, which inhibits sufficient orbital overlaps [8].

**Figure 8.** Magnetization vs. magnetic field of the as-synthesized bipyramid-like structures of Fe0.5Gd0.5(MoO4)1.5:Eu3+ synthesized by hydrothermal treatment at 200°C for 24 h.

### **4. Conclusion**

oxygen (O2–) in the final product. The EDX spectrum shown in Figure 5 (a, b) confirms the

**Figure 7.** Room temperature emission spectra of Eu3+ doped Fe0.5Gd0.5(MoO4)1.5 synthesized by the PVP-assisted hydro‐

Figure 7 shows the PL excitation and emission spectra of nanoparticle sheathed bipyramidlike microstructures of Fe0.5Gd0.5(MoO4)1.5:Eu3+. Emission (*λ*ex = 395 nm) spectra of the Fe0.5Gd0.5(MoO4)1.5 doped with Eu3+ ions were recorded within the range from 575 to 700 nm at room temperature. Upon 395 nm UV irradiation, the emission spectrum was governed by the

(where J = 1, 2, 3, 4) were observed due to intra-configurational f–f electronic transitions of Eu3+ ions. The presence of electric dipole transition confirmed that Eu3+ ions were located at

were comparatively very weak for all the samples (not shown in PL spectrum). However, the presence of magnetic iron (Fe3+) in the Fe0.5Gd0.5(MoO4)1.5 tetragonal crystal structure consid‐ erably decreases their luminescence, as Fe3+ ions share the dodecahedral sites with Gd3+/Eu3+ [13]. However, it is still unclear for the weak red emission in Eu3+ doped Fe0.5Gd0.5(MoO4)1.5. Figure 8 shows the room-temperature (RT) magnetic properties of hierarchically self-organ‐ ized Fe0.5Gd0.5(MoO4)1.5:Eu3+ 3D microstructures synthesized by the hydrothermal route at 24 h time interval. Magnetization as a function of applied magnetic field (from 15000 Oe to 15000 Oe) was measured using a vibrating sample magnetometer. Due to the presence of Gd3+ ions, a magnetic hysteresis (M–H) loop was not observed in the case of Eu3+ doped Fe0.5Gd0.5(MoO4)1.5. From Figure 8, it is clear that a straight line crosses the origin, indicating that the phosphor Fe0.5Gd0.5(MoO4)1.5:Eu3+ exhibits a paramagnetic behavior due to the existence

sites without inversion symmetry (C3v symmetry). The other transitions 5

D0→<sup>7</sup>

F1 (magnetic dipole). In the emission spectrum, the transitions 5

F2 (due to electric dipole transition)

D0→<sup>7</sup>

F3 and 5

D0→<sup>7</sup> FJ

D0→<sup>7</sup> F4

presence of Fe, Gd, Eu, Mo, and O in the final product.

hypersensitive red emission, showing a transition 5

thermal route.

42 Nanoparticles Technology

higher than 5

D0→<sup>7</sup>

**3.5. Optical and magnetic properties of Fe0.5Gd0.5(MoO4)1.5:Eu3+ structures**

In summary, the interaction between the Gd3+ and Fe3+ ions in the crystal structure Fe0.5Gd0.5(MoO4)1.5:Eu3+ is weak and hence insufficient overlap of the orbitals associated with 4f shells gives rise to paramagnetism. Further, more detailed investigations are required to determine the magnetic behavior of Fe0.5Gd0.5(MoO4)1.5:Eu3+ at various temperatures. The nanoparticle sheathed Fe0.5Gd0.5(MoO4)1.5:Eu3+ porous bipyramid-like structures were success‐ fully synthesized using hydrothermal technique by employing PVP as the surfactant. The hydrothermal reaction time period was the key parameter and plays an important role for morphological evolution of self-aggregated 3D hierarchical networks. A possible growth mechanism for the formation of microstructures clearly involves the layer-by-layer selfassembly of 2D nanoflakes with successive disassembly, via an Ostwald ripening phenomena. The photoluminescence emission properties of Eu3+ doped Fe0.5Gd0.5(MoO4)1.5 were investigat‐ ed. Under the excitation of 395 nm UV light, the as-synthesized samples show red emission from the hypersensitive 5 D0→<sup>7</sup> F2 transition at 615 nm. RT magnetic properties of the nanopar‐ ticles sheathed porous bipyramid-like Fe0.5Gd0.5(MoO4)1.5:Eu3+ structures clearly exhibit the paramagnetic behavior due to the presence of paramagnetic Gd3+ ions. The current experi‐ mental results are important and may open up a new scope for the design and formulation of bifunctional magnetic and luminescence applications.

### **Author details**

Rajagopalan Krishnan1 and Jagannathan Thirumalai2\*

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

1 Department of Physics, Rajalakshmi Institute of Technology, Kuthampakkam, Chennai, Tamil Nadu, India

2 Department of Physics, B. S. Abdur Rahman University, Vandalur, Chennai, Tamil Nadu, India

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ticles sheathed porous bipyramid-like Fe0.5Gd0.5(MoO4)1.5:Eu3+ structures clearly exhibit the paramagnetic behavior due to the presence of paramagnetic Gd3+ ions. The current experi‐ mental results are important and may open up a new scope for the design and formulation of

1 Department of Physics, Rajalakshmi Institute of Technology, Kuthampakkam, Chennai,

2 Department of Physics, B. S. Abdur Rahman University, Vandalur, Chennai, Tamil Nadu,

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## **Effect of Argon Carrier Gas Flux on TiO2 Nanostructures**

Reza Sabet Dariani and Zohreh Nafari Qaleh

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/61639

#### **Abstract**

TiO2 nanowires (NWs) have been fabricated by thermal evaporation method, using the vapor–liquid–solid (VLS) growth method and Au thin layer as a catalyst which were coated on titanium/quartz substrate. Optical and microstructure properties of TiO2 NWs have been reviewed under different argon carrier gas flows. The gas flow was varied from 60 to 170 sccm. The X-ray diffraction patterns showed that TiO2 NWs formed in ru‐ tile phase and their preferred orientation was in (110). The field emission scanning elec‐ tron microscope figures indicated that with increasing gas flow, nanowires became thinner and longer. The absorbance curve of TiO2 nanowires has been drawn by using spectrophotometer and their band gap was measured. Our results showed that the band gap has been enhanced in comparison with bulk one due to the quantum confinement. Also, the band gap enhanced from 3.35 to 3.45 eV by increasing the argon flow, since the wires became thinner and longer. The Williamson–Hall mechanism was applied as the vi‐ tal tools for characterization of crystallite size and microstrain. Also, microstructure prop‐ erties have been studied by X-ray diffraction analysis such as; dislocation density and texture coefficient. The crystallinity of TiO2 NWs has been made better by enhancing Ar flow due to increasing of VLS growth method.

**Keywords:** TiO2, Nanostructure, Absorption coefficient

### **1. Introduction**

### **1.1. Semiconducting nanowires**

Based on the geometry, 1D nanostructures typically consist of four categories: nanowire (NW), nanorod (NR), nanobelt (NB), and nanotube. The first three categories all possess a solid structure, while the nanotube has a hollow center. There is a significant different in geometry, property, and synthesis strategies between the nanotube and the first three categories, i.e., NW, NR, and NB. In general, NRs are NWs with a small aspect ratio (i.e., short NWs), which

© 2015 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.

typically appear rigid and straight. NB is a special NW structure with a rectangular crosssection. Typical NB structures are single crystals with well-defined side crystal facets. Some‐ times, nanofibers were also used to describe the NW morphology, particularly when the NWs are very long and not single crystalline [1]. Nanoscaled materials are of great interest for their unique structure and properties. Compared to isotropic nanoparticles and two dimensional thin films, one dimensional nanostructures such as nanofibers and nanowires are much more anisotropic, i.e., the aspect ratio of length to diameter can attain very high values.

Semiconductor one-dimensional structures, i.e., nanorods and nanowires with diameters ranging from 1 to 400 nm and lengths of up to hundreds of micrometers, are known to have many interesting physical properties and great potential applications in semiconductor and electronic technologies, which have stimulated intensive, worldwide research activities. They are perhaps the most versatile building blocks for optical and (opto-)electronic circuits at the nanoscale. Preliminary studies of some of these one-dimensional nanostructures indicate that they may be used as microscopic probes, field emission sources, storage materials, and lightemitting devices with extremely low power consumption. Among these semiconductors, some metal oxides, for instance, stannic oxide, titanium oxide, cuprous oxide, and vanadium pentoxide are the most important base materials in industry for gas sensors, transistors, and electrode material as well as catalysts. There has been extensive research into various "topdown" and "bottom-up" methods to create these nanostructures. In order to investigate their properties, many methods have been developed to synthesize one-dimensional nanomaterials, such as carbon nanotubes confined chemical reaction, Chemical Vapor Deposition (CVD), Molecular Beam Epitaxy (MBE), vapor-liquid-solid (VLS), solution–liquid–solid (SLS), and template-based synthetic approaches. However, complicated apparatus, complex process control, and special condition may be required for these approaches [2, 3]. Here we will describe VLS method in details.

### **1.2. Titanium dioxide (TiO2)**

Based on the geometry, 1D nanostructures typically consist of four categories: nanowire (NW), nanorod (NR), nanobelt (NB), and nanotube. The first three categories all possess a solid structure, while the nanotube has a hollow center. There is a significant difference in geometry, property, and synthesis strategies between the nanotube and the first three types, i.e., NW, NR, and NB. Generally, NRs are NWs with a small aspect ratio (short NWs), which typically become rigid and straight. NB is a special NW structure with a rectangular cross- section. Typical NB structures are single crystals with well-defined side crystal facets. Sometimes, nanofibers were also used to explain the NW morphology, particularly when the NWs are very long and not single crystalline.

Since Fujishima and Honda (1972) reported that n-type titanium dioxide (TiO2) could be used as a photo anode for a photo electrochemical cell, titanium oxide has been comprehensively investigated owing to its great applications in many fields [4].

TiO2 naturally exhibits four different types of polymorphs, i.e., rutile, anatase, brookite, and TiO2 (B). In addition, several metastable polymorphs, such as TiO2 (H), TiO2 II, and perovskite, have been produced synthetically. Different phases have different properties and thus require different conditions to form desired morphologies and offer different performance. In general, rutile is the most stable phase, while the anatase, brookite, and TiO2 (B) are metastable and will transform into rutile under high temperature. This phase stability relationship also exists in TiO2 nanomaterial formations. Anatase phased nanomaterials were commonly obtained but always found in solution-based or low-temperature vapor deposition systems. High- temper‐ ature deposition or annealing would usually yield rutile TiO2 nanostructures. Brookite and TiO2 (B) phases were less common and typically obtained from solution-based growth systems. Other metastable phases have to be synthesized using particular types of precursors and under restrict conditions. They were very rarely observed as a stable nanomaterial phase [1, 5, 6].

typically appear rigid and straight. NB is a special NW structure with a rectangular crosssection. Typical NB structures are single crystals with well-defined side crystal facets. Some‐ times, nanofibers were also used to describe the NW morphology, particularly when the NWs are very long and not single crystalline [1]. Nanoscaled materials are of great interest for their unique structure and properties. Compared to isotropic nanoparticles and two dimensional thin films, one dimensional nanostructures such as nanofibers and nanowires are much more

Semiconductor one-dimensional structures, i.e., nanorods and nanowires with diameters ranging from 1 to 400 nm and lengths of up to hundreds of micrometers, are known to have many interesting physical properties and great potential applications in semiconductor and electronic technologies, which have stimulated intensive, worldwide research activities. They are perhaps the most versatile building blocks for optical and (opto-)electronic circuits at the nanoscale. Preliminary studies of some of these one-dimensional nanostructures indicate that they may be used as microscopic probes, field emission sources, storage materials, and lightemitting devices with extremely low power consumption. Among these semiconductors, some metal oxides, for instance, stannic oxide, titanium oxide, cuprous oxide, and vanadium pentoxide are the most important base materials in industry for gas sensors, transistors, and electrode material as well as catalysts. There has been extensive research into various "topdown" and "bottom-up" methods to create these nanostructures. In order to investigate their properties, many methods have been developed to synthesize one-dimensional nanomaterials, such as carbon nanotubes confined chemical reaction, Chemical Vapor Deposition (CVD), Molecular Beam Epitaxy (MBE), vapor-liquid-solid (VLS), solution–liquid–solid (SLS), and template-based synthetic approaches. However, complicated apparatus, complex process control, and special condition may be required for these approaches [2, 3]. Here we will

Based on the geometry, 1D nanostructures typically consist of four categories: nanowire (NW), nanorod (NR), nanobelt (NB), and nanotube. The first three categories all possess a solid structure, while the nanotube has a hollow center. There is a significant difference in geometry, property, and synthesis strategies between the nanotube and the first three types, i.e., NW, NR, and NB. Generally, NRs are NWs with a small aspect ratio (short NWs), which typically become rigid and straight. NB is a special NW structure with a rectangular cross- section. Typical NB structures are single crystals with well-defined side crystal facets. Sometimes, nanofibers were also used to explain the NW morphology, particularly when the NWs are very

Since Fujishima and Honda (1972) reported that n-type titanium dioxide (TiO2) could be used as a photo anode for a photo electrochemical cell, titanium oxide has been comprehensively

TiO2 naturally exhibits four different types of polymorphs, i.e., rutile, anatase, brookite, and TiO2 (B). In addition, several metastable polymorphs, such as TiO2 (H), TiO2 II, and perovskite, have been produced synthetically. Different phases have different properties and thus require

investigated owing to its great applications in many fields [4].

anisotropic, i.e., the aspect ratio of length to diameter can attain very high values.

describe VLS method in details.

**1.2. Titanium dioxide (TiO2)**

50 Nanoparticles Technology

long and not single crystalline.

Indeed among the semiconductor oxides, titanium oxide (TiO2) has received more attention during the last decade because of its high refractive index (2.5), high band gap (3 eV), trans‐ parency throughout the visible and infrared region, high chemical stability, and hardness [7]. Crystal structures of the two most important phases are anatase (*a* = 3.78 Å, *c* = 9.51 Å; space group of I41/amd) and rutile (*a* = 4.59, *c* = 2.96, space group of P42/mnm). Each cell of rutile consists of two TiO2 molecules, while anatase has four TiO2 molecules. Figure 1 gives crystal structures of rutile (a) and anatase (b). The common features for these structures are as follows: Ti atom is positioned at the center of six O atoms, forming a deformed Ti-O octahedron. The octahedral share top points and a rhombic side join together and thus construct a crystal. The differences between them exist: three sides shared for the brookite, four for anatase and two for rutile. Among them rutile is the most stable and anatase is less while brookite is not stable. Therefore, the first two are commonly concerned in nanomaterials [4].

**Figure 1.** Two kinds of TiO2 crystal structures: (a) rutile, (b) anatase.

TiO2 has showed a wide range of applications in hydrogen production, lithium-ion batteries, fuel cells, gas sensors, detoxification, photovoltaics, photocatalysts, and super capacitors due to its excellent solid-state physical properties [8]. Over the past two decades, there have been remarkable advances in the expansion of one-dimensional (1D) nanostructures. The wire-like structure not only introduces significant enlarged and well-defined crystal surfaces over planar structures but also provides 1D confined channel, which could basically tailor the transportation of electrons, phonons, and photons [9]. The one-dimensional structure such as TiO2 nanowire is considered as a good candidate for getting higher performance in those applications compared to the bulk one. For example, a TiO2 nanowire based electrode can provide a large surface area for effectively collecting photons and electrons [8]. Tao et al. [10] used TiO2 NW array as photo anodes of back-illuminated dye-sensitized solar cells. Shi et al. [11] grew TiO2 nanorods uniformly on dense Si NW array backbones, by using surfacereaction-limited pulsed chemical vapor deposition technique. They developed a 3D highdensity heterogeneous NW architecture, which is useful for photo electrochemical electrodes. Wang et al. [12] reported hydrogen-treated TiO2 NW arrays for photo electrochemical water splitting. Mandal and Bhattacharyya [13] presented through systematic studies the perform‐ ance of TiO2 NWs in sensing, quantification, and photocatalytic degradation of cationic dyes in aqueous solution. Wang et al. [14] indicated that TiO2 NWs have promising application for high energy density lithium-ion batteries. Xi et al. [15] grew single-crystalline rutile TiO2 nanorods on fluorine doped tin oxide glass for organic–inorganic hybrid solar cells by means of hydrothermal mechanisms.

Numerous routes have been explored to the synthesis of 1D TiO2 nanostructures from both bottom up and top down. Representative bottom-up approaches include a large variety of solution- and vapor-based growth strategies. Although bottom-up approaches remain as the major synthesis efforts, several top-down procedures have also been explored for 1D TiO2 growth, such as direct oxidation and electrochemical etching techniques. In general, there are four main approaches, which include solution-based synthetic approaches (such as hydro‐ thermal/ solvothermal and sol-gel methods), vapor-based approaches (chemical/physical vapor deposition, atomic layer deposition-related methods and pulsed laser deposition), templated growth, and top-down fabrication techniques. Solution-based growth techniques offer several major advantages for mass production of nanomaterials including low-cost, simple processing, and good scalability. Vapor deposition is typically conducted in vacuum and under high temperature. Compared to the wet chemistry methods, it offers several unique advantages. First, the high quality of NW's crystallinity is secured due to the utilization of high-temperature and high-vacuum deposition conditions. Second, since the vapor deposition processes are typically performed in conventional thin film deposition systems, the knowledge of thin film growth techniques can be adopted to understanding the NW growth behaviors. Lastly, by engineering catalysts, growth sites, or precursors, the control of location, composi‐ tion, dimension, and organization of NWs could be achieved via vapor deposition processes. Nevertheless, vapor deposition is usually associated with high cost and small- scale produc‐ tion. Template growth is a very versatile synthesis technique that forms nanostructure with a morphology that follows the porous template. Top-down fabrication techniques describe the creation of nanostructures from the bulk form. This approach typically is able to process a large quantity of nanostructures with uniform morphology and orientation. For TiO2 1D nanostruc‐ tures, a top-down process usually starts from a bulk Ti foil followed by selective etching and/ or oxidization. In the following section, the vapor- liquid- solid mechanism is described [1].

### **1.3. VLS mechanism of nanowire growth**

planar structures but also provides 1D confined channel, which could basically tailor the transportation of electrons, phonons, and photons [9]. The one-dimensional structure such as TiO2 nanowire is considered as a good candidate for getting higher performance in those applications compared to the bulk one. For example, a TiO2 nanowire based electrode can provide a large surface area for effectively collecting photons and electrons [8]. Tao et al. [10] used TiO2 NW array as photo anodes of back-illuminated dye-sensitized solar cells. Shi et al. [11] grew TiO2 nanorods uniformly on dense Si NW array backbones, by using surfacereaction-limited pulsed chemical vapor deposition technique. They developed a 3D highdensity heterogeneous NW architecture, which is useful for photo electrochemical electrodes. Wang et al. [12] reported hydrogen-treated TiO2 NW arrays for photo electrochemical water splitting. Mandal and Bhattacharyya [13] presented through systematic studies the perform‐ ance of TiO2 NWs in sensing, quantification, and photocatalytic degradation of cationic dyes in aqueous solution. Wang et al. [14] indicated that TiO2 NWs have promising application for high energy density lithium-ion batteries. Xi et al. [15] grew single-crystalline rutile TiO2 nanorods on fluorine doped tin oxide glass for organic–inorganic hybrid solar cells by means

Numerous routes have been explored to the synthesis of 1D TiO2 nanostructures from both bottom up and top down. Representative bottom-up approaches include a large variety of solution- and vapor-based growth strategies. Although bottom-up approaches remain as the major synthesis efforts, several top-down procedures have also been explored for 1D TiO2 growth, such as direct oxidation and electrochemical etching techniques. In general, there are four main approaches, which include solution-based synthetic approaches (such as hydro‐ thermal/ solvothermal and sol-gel methods), vapor-based approaches (chemical/physical vapor deposition, atomic layer deposition-related methods and pulsed laser deposition), templated growth, and top-down fabrication techniques. Solution-based growth techniques offer several major advantages for mass production of nanomaterials including low-cost, simple processing, and good scalability. Vapor deposition is typically conducted in vacuum and under high temperature. Compared to the wet chemistry methods, it offers several unique advantages. First, the high quality of NW's crystallinity is secured due to the utilization of high-temperature and high-vacuum deposition conditions. Second, since the vapor deposition processes are typically performed in conventional thin film deposition systems, the knowledge of thin film growth techniques can be adopted to understanding the NW growth behaviors. Lastly, by engineering catalysts, growth sites, or precursors, the control of location, composi‐ tion, dimension, and organization of NWs could be achieved via vapor deposition processes. Nevertheless, vapor deposition is usually associated with high cost and small- scale produc‐ tion. Template growth is a very versatile synthesis technique that forms nanostructure with a morphology that follows the porous template. Top-down fabrication techniques describe the creation of nanostructures from the bulk form. This approach typically is able to process a large quantity of nanostructures with uniform morphology and orientation. For TiO2 1D nanostruc‐ tures, a top-down process usually starts from a bulk Ti foil followed by selective etching and/ or oxidization. In the following section, the vapor- liquid- solid mechanism is described [1].

of hydrothermal mechanisms.

52 Nanoparticles Technology

Nanowires are commonly grown using vapor, solution or (template directed) electrodeposi‐ tion methods. High- temperature growth from the vapor phase is often preferred due to the high crystal quality that can be obtained and the ability to grow large quantities of wires at once. A key factor in most vapor- and solution-based methods is the presence of small metal droplets during synthesis. Analysis of the growth of silicon whiskers (hairs) from the vapor phase using gold catalyst particles leads to the postulation of the vapor- liquid- solid (VLS) mechanism of growth. The VLS mechanism consists of three stages which are illustrated in figure 2A. First, a metal particle absorbs semiconductor material and forms an alloy. In this step, the volume of the particle increases and the particle often transition from a solid to a liquid state. Second, the alloy particle absorbs more semiconductor material until it is satu‐ rated. The saturated alloy droplet becomes in equilibrium with the solid phase of the semi‐ conductor and nucleation occurs (i.e., solute/solid phase transition). During the final phase, a steady state is formed in which a semiconductor crystal grows at the solid/liquid interface. The precipitated semiconductor material grows as a wire because it is energetically more favorable than extension of the solid- liquid interface. The semiconductor material is precipitated at the existing solid/liquid interface as opposed to the formation of a new interface.

**Figure 2.** (A) The stages of (I) alloying, (II) nucleation, and (III) growth of nanowire synthesis according to the VLS growth mechanism. (B) Pseudobinary phase diagram of a semiconductor-gold system. The arrows indicate the subsequent phas‐ es when a gold droplet absorbs semiconductor material at a constant temperature. (1) At first, the gold particle is in the sol‐ id or liquid state. By absorbing semiconductor material, a liquid alloy is formed. (II) Subsequent absorption of more semiconductor material allows the liquid alloy to be in equilibrium with the solid semiconductor [2].

In the VLS mechanism, the wire diameter is determined by the diameter of the alloy particle which is in turn determined by the low temperature size of the metal particle and the temper‐ ature. The wire length is determined by the growth rate and time. When the system is cooled, the alloy droplet solidifies at the wire tip. To examine the feasibility of VLS wire growth from a certain semiconductor/metal combination, it is essential to study the (pseudo)binary phase diagram (figure 2B); the metal should form an alloy with the semiconductor at a temperature that also allows the semiconductor to exist in the solid phase. The validity of the VLS mecha‐ nism of wire growth has been proven for germanium nanowires by in- situ high- temperature TEM measurements. Heating neighboring gold and germanium clusters to growth tempera‐ tures and selectively evaporating the germanium clusters with an electron beam are allowed for the direct imaging of the successive steps of alloying and melting, nucleation and wire growth. A study using colloidal gold catalyst particles showed that using the VLS method, (single) crystalline nanowires with diameters as small as 5 nm could be obtained. In addition, it has been shown that by increasing the VLS growth rate by modulating the temperature, crystallization only takes place at the surface of the catalyst particle resulting in the growth of nanotubes. Since the late 1990's, semiconductor nanowire growth using the VLS method has enjoyed an increasing popularity with researchers and industry [2].


Some of studies on TiO2 nanowires are listed in Table 1.

**Table 1.** Some of works on the fabrication of TiO2 nanowires

The vapor–liquid–solid and the vapor–solid growth mechanisms have been used to fabricate high-quality oxide NWs with short heating times. Ha et al. [28] synthesized single-crystalline rutile TiO2 NWs by VLS method on Ti foil substrates patterned with catalytic Sn nano-islands, Lee et al. [18] grew TiO2 NWs in this way but used Au as catalysts, and Park and Lee [27] reported synthesis of TiO2 NWs by Ni as catalysts.

TiO2 nanowire is typically achieved using either of the two different synthetic methods developed over the past few years: one is the wet-chemistry method such as electrochemical, sol-gel electrophoresis, hydrothermal growth on Ti substrates at 180 ℃ utilizing various organic solvents to oxidize Ti [37], direct oxidizing Ti substrate with aqueous hydrogen peroxide solutions [10], and solvothermal method [38]. The other is the dry method such as, growth single-crystalline TiO2 nanowires by thermal evaporation in atmospheric-pressure (101325 Pa) [30], metal-organic chemical vapor deposition [39], pulsed chemical vapor deposition [8], chemical vapor transport process by evaporating Ti metal powders [18], heating titanium grids (substrate) by ethanol-embedding followed by an annealing in the air [40], and by a facile water- assisted chemical vapor deposition method based on commercial metal titanium [4]. TiO2 nanowires produced by the wet-chemistry method due to poor crystallinity have difficulty integrating into device fabrication. Also, they need additional processes in order to improve their crystallinity [27].

In this work, TiO2 nanowires were synthesized by a carbothermal evaporation process on quartz substrate at 850 ℃ for 1 h at atmospheric pressure (101325 Pa), and Au was used as a catalyst. Also, the influence of argon flow as carrier gas on microstructural and optical characteristics was measured. We have been characterized some microstructural parameters like crystallite size, microstrain within the crystallite, dislocation density, and texture coeffi‐ cient by the diffraction pattern.

### **2. Experimental details**

(single) crystalline nanowires with diameters as small as 5 nm could be obtained. In addition, it has been shown that by increasing the VLS growth rate by modulating the temperature, crystallization only takes place at the surface of the catalyst particle resulting in the growth of nanotubes. Since the late 1990's, semiconductor nanowire growth using the VLS method has

**Year Researcher Fabrication method Carrier gas References** 2002 Zhang et al. Hydrothermal synthesis None [16]

2006 Lee et al. VLS 100 sccm (Ar) + 0.5 sccm(O2) [18] 2006 Wu et al. Thermal evaporation 20 sccm (Ar) + 1 sccm (O2) [19] 2006 Gyorgy et al. Pulsed laser deposition None [20]

2007 Berger et al. Chemical bath deposition None [22] 2008 Huo et al. Thermal 150 sccm (Ar) [23] 2008 Baik et al. VLS 200 sccm (Ar) [24] 2009 Huo et al. Direct Growth on ti foil 50 sccm (Ar) [25] 2011 RamezaniSani et al. Thermal evaporation 50- 150 sccm (Ar) [26]

2011 Park et al. VLS and VS 200 sccm (Ar) [27] 2011 Ha et al VLS Ar with 2% H2 upon reaching 1 atm [28] 2012 Wu et al. Solvothermal method None [29] 2012 Shang et al. Thermal Evaporation 100 sccm (Ar) [30]

2013 Dariani et al. VLS 60- 170 sccm (Ar) [32] 2013 Tian et al. Electron beam lithography None [33] 2013 Liu et al Molten-salt flux method None [34] 2014 Nechache et al. Pulsed laser deposition 300 mTorr (O2/Ar) [35]

deposition 500–800 sccm (Ar) [17]

vapor deposition 400 sccm (Ar) [4]

10 sccm (Ar) + 1 sccm (O2) [21]

None [31]

None [36]

enjoyed an increasing popularity with researchers and industry [2].

2007 Amin et al. VLS 10 sccm (Ar)/

water-assisted chemical

hydrothermal method.

low temperature

Some of studies on TiO2 nanowires are listed in Table 1.

2003 Pradhan et al. Metalorganic chemical vapor

2013 Liu et al. Low-temperature

2015 Wang et al. Confined-space synthesis at

**Table 1.** Some of works on the fabrication of TiO2 nanowires

2011 Liu et al.

54 Nanoparticles Technology

In this study, pieces of quartz with dimensions of 1 × 15 × 15 mm3 are used. Initially, the substrates were ultrasonically cleaned with acetone and ethanol for 10 minutes, respectively. Then a layer of titanium was deposited on substrates by electron beam evaporation technique with the evaporation rate of 0.2 Å/s under the pressure of 1.5 × 10 -5 mbar. In this step, voltage and current were kept 4 kV and 100 mA, respectively. The electron beam evaporation device is specified by: HIND HIGH VACUUM CO. Model 12A4D. After removing the samples, a gold thin layer with 3 nm thickness was coated on the titanium layer by sputtering method.

In next step, a tablet containing of titanium and graphite with the ratio of 1:1 as a source material was placed in the center of horizontal furnace. The center of the furnace was heated to 1050℃ with heating rate of 20℃/min. According to phase diagram of the Au-Ti system, the eutectoid temperature is about 832℃ [30]. The eutectoid temperature is a point in which two metals in solid phase form an alloy in solid phase (no in liquid phase). Thus, the sample was placed in distance of 12 - 13 cm (with about 850 ℃ temperature) from the center of the furnace for 1 h at atmospheric pressure and under argon flow. The vapor was generated from source

Figure 3. A schematic of CVD system.

In next step, a tablet containing of titanium and graphite with the ratio of 1:1 as a source material was

**Figure 3.** A schematic of CVD system.

material transported to the substrate and at last condensed on the surface of substrate (see figure 3). Then the substrate was naturally cooled down to room temperature. In this study, three samples under different argon gas flows (60 sccm, 110 sccm, and 170 sccm) were fabricated. Argon gas was also used to control the oxygen content of experiment. It is believed that oxygen comes from leaked air during heating process. Content of argon was kept constant during heating process by using bubbling trap at the end of the vent line. The microstructure and the morphology of the samples were analyzed by x-ray diffraction (XRD: Philips X-Pert PW 3040/60, *λ* (Cukα) = 1.54 ˚A) and field emission scanning electron microscope (FESEM: Hitachi S4160 with operating voltage of 15 kV), respectively. The absorption spectrum of TiO2 nanostructures was evaluated by spectrophotometer (Ocean Optics, Model HR4000 CG-UV-NIR). placed in the center of horizontal furnace. The center of the furnace was heated to 1050Ԩ with heating rate of 20Ԩ/min. According to phase diagram of the Au-Ti system, the eutectoid temperature is about 832Ԩ [30]. The eutectoid temperature is a point in which two metals in solid phase form an alloy in solid phase (no in liquid phase). Thus, the sample was placed in distance of 12 - 13 cm (with about 850 Ԩ temperature) from the center of the furnace for 1 h at atmospheric pressure and under argon flow. The vapor was generated from source material transported to the substrate and at last condensed on the surface of substrate (see figure 3). Then the substrate was naturally cooled down to room temperature. In this study, three samples under different argon gas flows (60 sccm, 110 sccm, and 170 sccm) were fabricated. Argon gas was also used to control the oxygen content of experiment. It is believed that oxygen comes from leaked air during heating process. Content of argon was kept constant during heating process by using bubbling trap at the end of the vent line. The microstructure and the morphology of the samples were analyzed by x-ray diffraction (XRD: Philips X-Pert PW 3040/60, *λ* (Cukα) = 1.54 A) and field emission scanning electron microscope (FESEM: Hitachi S4160 with operating voltage of 15 kV), respectively. The absorption spectrum of TiO2 nanostructures was evaluated by spectrophotometer (Ocean Optics, Model HR4000 CG-

#### **3. Results and discussion III. Results and discussion 1. SEM observations**

#### **3.1. SEM observations** The morphologic characteristics of the samples under different carrier gas flows are determined by SEM images. The cross-sectional views of the samples are shown in figure 4.

UV-NIR).

The morphologic characteristics of the samples under different carrier gas flows are deter‐ mined by SEM images. The cross-sectional views of the samples are shown in figure 4.

The average range of length and diameter of nanowires are given in Table 2. It can be seen that nanowires become thinner and longer with increasing the carrier gas flow. Since fabrication parameters except argon flow were fixed, it can be concluded that the amount of O2 inside the reaction chamber plays a vital role in the growth of TiO2 1D nanostruc‐ tures. When the O2 concentration was high, i.e., Ar flow was low, the VLS growth (will describe later) was suppressed because Ti vapor could readily react with O2 before forming alloy droplets with Au as catalyst [12]. Therefore, in this condition, a relatively low concentration of O2 is necessary [13]. Top View 

7

Top View Cross section view

In next step, a tablet containing of titanium and graphite with the ratio of 1:1 as a source material was placed in the center of horizontal furnace. The center of the furnace was heated to 1050Ԩ with heating rate of 20Ԩ/min. According to phase diagram of the Au-Ti system, the eutectoid temperature is about 832Ԩ [30]. The eutectoid temperature is a point in which two metals in solid phase form an alloy in solid phase (no in liquid phase). Thus, the sample was placed in distance of 12 - 13 cm (with about 850 Ԩ temperature) from the center of the furnace for 1 h at atmospheric pressure and under argon flow. The vapor was generated from source material transported to the substrate and at last condensed on the surface of substrate (see figure 3). Then the substrate was naturally cooled down to room temperature. In this study, three samples under different argon gas flows (60 sccm, 110 sccm, and 170 sccm) were fabricated. Argon gas was also used to control the oxygen content of experiment. It is believed that oxygen comes from leaked air during heating process. Content of argon was kept constant during heating process by using bubbling trap at the end of the vent line. The microstructure and the morphology of the samples were analyzed by x-ray diffraction (XRD: Philips X-Pert PW 3040/60, *λ* (Cukα) = 1.54 A) and field emission scanning electron microscope (FESEM: Hitachi S4160 with operating voltage of 15 kV), respectively. The absorption spectrum of TiO2 nanostructures was evaluated by spectrophotometer (Ocean Optics, Model HR4000 CG-

Figure 4. Top and cross-sectional SEM views of TiO2 nanostructures. Argon gas flow increases from left to right. **Figure 4.** Top and cross-sectional SEM views of TiO2 nanostructures. Argon gas flow increases from left to right.


**Table 2.** Average values of diameter and length of nanowires.

necessary [13].

UV-NIR).

**III. Results and discussion 1. SEM observations** 

material transported to the substrate and at last condensed on the surface of substrate (see figure 3). Then the substrate was naturally cooled down to room temperature. In this study, three samples under different argon gas flows (60 sccm, 110 sccm, and 170 sccm) were fabricated. Argon gas was also used to control the oxygen content of experiment. It is believed that oxygen comes from leaked air during heating process. Content of argon was kept constant during heating process by using bubbling trap at the end of the vent line. The microstructure and the morphology of the samples were analyzed by x-ray diffraction (XRD: Philips X-Pert PW 3040/60, *λ* (Cukα) = 1.54 ˚A) and field emission scanning electron microscope (FESEM: Hitachi S4160 with operating voltage of 15 kV), respectively. The absorption spectrum of TiO2 nanostructures was evaluated by spectrophotometer (Ocean Optics, Model HR4000 CG-

Figure 3. A schematic of CVD system.

In next step, a tablet containing of titanium and graphite with the ratio of 1:1 as a source material was placed in the center of horizontal furnace. The center of the furnace was heated to 1050Ԩ with heating rate of 20Ԩ/min. According to phase diagram of the Au-Ti system, the eutectoid temperature is about 832Ԩ [30]. The eutectoid temperature is a point in which two metals in solid phase form an alloy in solid phase (no in liquid phase). Thus, the sample was placed in distance of 12 - 13 cm (with about 850 Ԩ temperature) from the center of the furnace for 1 h at atmospheric pressure and under argon flow. The vapor was generated from source material transported to the substrate and at last condensed on the surface of substrate (see figure 3). Then the substrate was naturally cooled down to room temperature. In this study, three samples under different argon gas flows (60 sccm, 110 sccm, and 170 sccm) were fabricated. Argon gas was also used to control the oxygen content of experiment. It is believed that oxygen comes from leaked air during heating process. Content of argon was kept constant during heating process by using bubbling trap at the end of the vent line. The microstructure and the morphology of the samples were analyzed by x-ray diffraction (XRD: Philips X-Pert PW 3040/60, *λ* (Cukα) = 1.54 A) and field emission scanning electron microscope (FESEM: Hitachi S4160 with operating voltage of 15 kV), respectively. The absorption spectrum of TiO2 nanostructures was evaluated by spectrophotometer (Ocean Optics, Model HR4000 CG-

The morphologic characteristics of the samples under different carrier gas flows are deter‐ mined by SEM images. The cross-sectional views of the samples are shown in figure 4.

The morphologic characteristics of the samples under different carrier gas flows are determined by SEM

The average range of length and diameter of nanowires are given in Table 2. It can be seen that nanowires become thinner and longer with increasing the carrier gas flow. Since fabrication parameters except argon flow were fixed, it can be concluded that the amount of O2 inside the reaction chamber plays a vital role in the growth of TiO2 1D nanostruc‐ tures. When the O2 concentration was high, i.e., Ar flow was low, the VLS growth (will describe later) was suppressed because Ti vapor could readily react with O2 before forming alloy droplets with Au as catalyst [12]. Therefore, in this condition, a relatively low

7

UV-NIR).

Top View

UV-NIR).

**3. Results and discussion**

**Figure 3.** A schematic of CVD system.

56 Nanoparticles Technology

images. The cross-sectional views of the samples are shown in figure 4.

concentration of O2 is necessary [13].

**3.1. SEM observations**

**III. Results and discussion 1. SEM observations** 

Figure 5 shows compositional energy-dispersive X-ray spectroscopy (EDAX) analysis of the NW tip and the confirmation of the existence of Au alloy.

7

**Figure 5.** A typical EDAX of our samples which shows the existence of Au.

### **3.2. XRD studies**

To examine the effect of argon flow on structural properties of TiO2 1D nanostructure, the XRD analyses were performed. Figure 6 shows XRD patterns of TiO2 grown on Ti/quartz substrates.

**Figure 6.** XRD patterns of TiO2 nanowires at different carrier gas flows.

Rutile phase is formed on the samples and (110) direction is preferred direction which is consistent with SEM images. In the TiO2 tetragonal structure, the plane spacing (*d*) is related to lattice constants (*a* and *c*) and the Miller indices (*h*, *k*, and *l*) by the following relation [42]:

$$1\left(d\_{hkl}\right)^2 = h^2 + k^2 \;/\; a^2 + l^2 \;/\; c^2 \tag{1}$$

The results are presented in Table 3 and are consistent with standard card of JCPDS (73-1765).


**Table 3.** Lattice constants and E g of nanowires at different carrier gas flows.

### **3.3. Optical properties**

**3.2. XRD studies**

58 Nanoparticles Technology

To examine the effect of argon flow on structural properties of TiO2 1D nanostructure, the XRD analyses were performed. Figure 6 shows XRD patterns of TiO2 grown on Ti/quartz substrates.

Rutile phase is formed on the samples and (110) direction is preferred direction which is consistent with SEM images. In the TiO2 tetragonal structure, the plane spacing (*d*) is related to lattice constants (*a* and *c*) and the Miller indices (*h*, *k*, and *l*) by the following relation [42]:

The results are presented in Table 3 and are consistent with standard card of JCPDS (73-1765).

**Sample** *a (Å) c (Å) Eg (eV)* S1 4.586 2.951 3.35±0.02 S2 4.583 2.955 3.39±0.02 S3 4.589 2.949 3.45±0.02

( )<sup>2</sup> 2 2222 1 / / / *hkl d h ka lc* =+ + (1)

**Figure 6.** XRD patterns of TiO2 nanowires at different carrier gas flows.

JCPDS 4.589 2.954

**Table 3.** Lattice constants and E g of nanowires at different carrier gas flows.

The absorbance spectrum of TiO2 nanowires is measured by spectrophotometer. According to Beer-Lambert's law (Eq. 2), transmittance (Eq. 3), and absorbance (Eq. 4) definitions [43], the absorption coefficients of samples according to absorbance are obtained by Eq. 5:

$$I\left(\mathcal{X}\right) = I\_0\left(\mathcal{X}\right)e^{-\alpha(\lambda)d} \tag{2}$$

$$\mathbf{T(x)} = \mathbf{I(x)} \land \mathbf{I'(x)}\tag{3}$$

$$A\begin{pmatrix} \lambda \end{pmatrix} = \begin{pmatrix} -\log\left(T\_{\lambda}\right) \end{pmatrix} \tag{4}$$

$$\alpha = \begin{pmatrix} -1/d \end{pmatrix} \operatorname{Ln} \left( 10^{\circ \lambda} \right) = \begin{pmatrix} \operatorname{A}/\operatorname{d} \end{pmatrix} \operatorname{Ln} \left( 10 \right) \tag{5}$$

where *d* is the thickness of samples derived from the cross-sectional images of SEM and *A* is the absorbance of samples.

It is known that the relationship between absorption coefficient α near the absorption edge and the optical band gap *Eg* is determined by Tauc relation [15]:

$$
abla \nu = \mathbb{C} \left( h \nu - E\_{\frac{\nu}{\lambda}} \right)^{\mathbb{N}} \tag{6}$$

where *C* is a constant, *n* is equal to 0.5 or 2 for direct and indirect allowed transitions, respec‐ tively, and *hν* is photon energy. The *Eg* can be graphically estimated by a linear fit of the highenergy tail of (*αhν*) 1/*<sup>n</sup>*. Since the samples have rutile structure, rutile titania has direct band gap (*n* = 0.5).

Figure 7 shows (*αhν*) <sup>2</sup> plots versus photon energy *hν* for direct transition and intercepts value of the linear-fit extrapolations with the horizontal axes. The values of *Eg* are given in Table 2.

The optical results have shown that all the samples are characterized by energy gap higher than bulk rutile TiO2 (*Eg* = 3 eV for bulk) [28]. As the carrier gas concentration increases, the band gap increases from 3.35 to 3.45 eV. The higher value of the rutile titania energy gap are attributable to the well-known relationship between band gap and crystallites size. The appearance of quantum confinement effect in small sized nano crystalline titania is resulting in a blue shift of the band gap [37]. Among the samples, sample 3 has the highest *Eg* value which is the thinnest one.

curves versus photon energy (*hν*) for the samples.

**Figure 7.** (*αhν*) 2 curves versus photon energy (*hν*) for the samples.

Figure 7. (*αhν*)

2

#### One-dimensional nanostructures are grown in this study by VLS mechanism. In the mechanism, the role of the impurity (catalyst) is to form a liquid alloy droplet of eutectoid temperature. Metal catalyst used in the **3.4. Growth mechanism**

**4. Growth mechanism** 

titanium form. On the other hand, during the thermal evaporation process, at first, titanium powder in the presence of oxygen atmosphere was converted toTiO2 powder (see Eq. 7) [44], then the TiO2 powder was reduced by graphite to produce Ti vapor as the following chemical reaction [45]. Ti(s)+O2(g)=TiO2(s) (298-2500Ԩ) (7) TiO2(s)+C(s)=Ti(g)+CO2(g) (8) We have used Ti metal and carbon instead of TiO2 in carbothermal reduction process. In fact, Ti in the tablet converts to TiO2 by oxygen in air (Eq. 7), and then TiO2 is reduced in carbothermal reduction process by carbon (Eq. 8). Indeed, carbon mainly was used for TiO2 reduction. Also there are some papers that have used titanium and carbon as a source material [26, 46]. Previously, TiO2 and C as a source material were examined in our lab but NWs did not grow. Amin et al. [21] also reported that continuous increase of the amount of oxygen further reduced the yield of TiO2 1D nanostructures. When the O2 concentration was One-dimensional nanostructures are grown in this study by VLS mechanism. In the mecha‐ nism, the role of the impurity (catalyst) is to form a liquid alloy droplet of eutectoid tempera‐ ture. Metal catalyst used in the mechanism was gold with 3 nm thickness. At high temperatures (about 850℃) alloy liquid droplets of Au-titanium form. On the other hand, during the thermal evaporation process, at first, titanium powder in the presence of oxygen atmosphere was converted toTiO2 powder (see Eq. 7) [44], then the TiO2 powder was reduced by graphite to produce Ti vapor as the following chemical reaction [45].

high, the VLS growth was suppressed because Ti vapor could readily react with O2 before forming alloy

absorbed into the alloy. The liquid alloy acts as a preferred sink for arriving nanostructure components. The Ti and O2 are dissolved into the mediating alloy liquid droplets. When alloy droplets reach to saturation of

mechanism until the Au is consumed or until the growth conditions are changed [47]. Here, by over an hour, furnace temperature began to reduce and growth condition of one-dimensional titanium dioxide

mechanism was gold with 3 nm thickness. At high temperatures (about 850Ԩ) alloy liquid droplets of Au-

$$\text{Ti(s)} \star \text{O}\_2\text{(g)} \text{=} \text{TiO}\_2\text{(s)} \text{ (298-2500 \quad )}\tag{7}$$

$$\text{TiO}\_2(s) \star \text{C}(s) = \text{Ti}(\text{g}) \star \text{CO}\_2(\text{g}) \tag{8}$$

10 nanostructures is stopped. Usually droplets of catalyst are observed on the top of the nanostructures in a VLS- based growth system. We have used Ti metal and carbon instead of TiO2 in carbothermal reduction process. In fact, Ti in the tablet converts to TiO2 by oxygen in air (Eq. 7), and then TiO2 is reduced in carbo‐ thermal reduction process by carbon (Eq. 8). Indeed, carbon mainly was used for TiO2 reduction. Also there are some papers that have used titanium and carbon as a source material [26, 46]. Previously, TiO2 and C as a source material were examined in our lab but NWs did not grow. Amin et al. [21] also reported that continuous increase of the amount of oxygen further reduced the yield of TiO2 1D nanostructures. When the O2 concentration was high, the VLS growth was suppressed because Ti vapor could readily react with O2 before forming alloy droplets with catalytic materials.

Since the adsorption of liquid is more than solid, the reactant components (vapor of titanium and oxygen) absorbed into the alloy. The liquid alloy acts as a preferred sink for arriving nanostructure components. The Ti and O2 are dissolved into the mediating alloy liquid droplets. When alloy droplets reach to saturation of reactive component, the 1D nanostructures grow by precipitation of TiO2 from the droplets in the interface liquid- solid to minimize free energy of alloy system. The 1D nanostructure grows in length by the mechanism until the Au is consumed or until the growth conditions are changed [47]. Here, by over an hour, furnace temperature began to reduce and growth condition of one-dimensional titanium dioxide nanostructures is stopped. presence of oxygen atmosphere was converted toTiO2 powder (see Eq. 7) [44], then the TiO2 powder was reduced by graphite to produce Ti vapor as the following chemical reaction [45]. Ti(s)+O2(g)=TiO2(s) (298-2500Ԩ) (7) TiO2(s)+C(s)=Ti(g)+CO2(g) (8) We have used Ti metal and carbon instead of TiO2 in carbothermal reduction process. In fact, Ti in the tablet converts to TiO2 by oxygen in air (Eq. 7), and then TiO2 is reduced in carbothermal reduction process by carbon (Eq. 8). Indeed, carbon mainly was used for TiO2 reduction. Also there are some papers that have used titanium and carbon as a source material [26, 46]. Previously, TiO2 and C as a source material were examined in our lab but NWs did not grow. Amin et al. [21] also reported that continuous increase of the amount of oxygen further reduced the yield of TiO2 1D nanostructures. When the O2 concentration was high, the VLS growth was suppressed because Ti vapor could readily react with O2 before forming alloy

Usually droplets of catalyst are observed on the top of the nanostructures in a VLS- based growth system. However, the position of catalyst at the tip or root of the nanostructure depends on the interaction of catalyst nanoparticles and the substrate [48]. droplets with catalytic materials. Since the adsorption of liquid is more than solid, the reactant components (vapor of titanium and oxygen) absorbed into the alloy. The liquid alloy acts as a preferred sink for arriving nanostructure components. The Ti and O2 are dissolved into the mediating alloy liquid droplets. When alloy droplets reach to saturation of reactive component, the 1D nanostructures grow by precipitation of TiO2 from the droplets in the interface liquid- solid to minimize free energy of alloy system. The 1D nanostructure grows in length by the

SEM images (Figure 4) show by increasing argon flow or decreasing oxygen (one of reactants), nanowires become thinner and longer. O2 plays a main role in the growth of TiO2 nanowires. When O2 concentration is high, i.e., Ar flow is low, Ti vapor could readily react with O2, before they adsorb on alloy droplets [30]. Therefore, in this condition, a relatively low concentration of O2 is necessary [15, 31]. According to other findings [17, 27, 48], we suggest that thinner and longer NWs with increasing Ar flow are due to the following reasons: mechanism until the Au is consumed or until the growth conditions are changed [47]. Here, by over an hour, furnace temperature began to reduce and growth condition of one-dimensional titanium dioxide nanostructures is stopped. Usually droplets of catalyst are observed on the top of the nanostructures in a VLS- based growth system. However, the position of catalyst at the tip or root of the nanostructure depends on the interaction of catalyst nanoparticles and the substrate [48]. SEM images (Figure 4) show by increasing argon flow or decreasing oxygen (one of reactants), nanowires become thinner and longer. O2 plays a main role in the growth of TiO2 nanowires. When O2 concentration is high, i.e., Ar flow is low, Ti vapor could readily react with O2, before they adsorb on alloy droplets [30]. Therefore, in this condition, a relatively low concentration of O2 is necessary [15, 31]. According to other

findings [17, 27, 48], we suggest that thinner and longer NWs with increasing Ar flow are due to the

following reasons:

Figure 8. Schematic illustration of growth procedure for TiO2 nanowires on Ti/quartz substrate in (a) low and (b) high argon carrier gas flow [32]. Argon flux plays a major role in combining Ti and O2 and producing TiO2. A schematic of the proposed **Figure 8.** Schematic illustration of growth procedure for TiO2 nanowires on Ti/quartz substrate in (a) low and (b) high argon carrier gas flow [32].

10 growth mechanism is shown in Figure 8. When Ar flux is low (Figure 8a), the possibility of producing TiO2 is much higher than when Ar flux is high (Figure 8b). The reason could be due to moving Ti and O2 gases from furnace tube when Ar flux is high and so the possibility of producing TiO2 becomes lower. On the Argon flux plays a major role in combining Ti and O2 and producing TiO2. A schematic of the proposed growth mechanism is shown in Figure 8. When Ar flux is low (Figure 8a), the possibility of producing TiO2 is much higher than when Ar flux is high (Figure 8b). The reason could be due to moving Ti and O2 gases from furnace tube when Ar flux is high and so the possibility of producing TiO2 becomes lower. On the other hand, in presence of O2, titanium layer on the substrate surface in 850℃ is easily converted to TiO2 (according to XRD diffraction patterns and Eq. 7). Thus, there is a competitive growth between TiO2 layer (epitaxial growth: TiO2 molecules on TiO2 layer) and alloy droplets (VLS growth). TiO2 islands that are formed around the NWs by VLS method become larger with decreasing Ar flow. Since the other parameters of growth, such as source and time are constant, NW diameters and lengths become larger and shorter, respectively.

10

curves versus photon energy (*hν*) for the samples.

One-dimensional nanostructures are grown in this study by VLS mechanism. In the mechanism, the role of the impurity (catalyst) is to form a liquid alloy droplet of eutectoid temperature. Metal catalyst used in the mechanism was gold with 3 nm thickness. At high temperatures (about 850Ԩ) alloy liquid droplets of Autitanium form. On the other hand, during the thermal evaporation process, at first, titanium powder in the presence of oxygen atmosphere was converted toTiO2 powder (see Eq. 7) [44], then the TiO2 powder was

One-dimensional nanostructures are grown in this study by VLS mechanism. In the mecha‐ nism, the role of the impurity (catalyst) is to form a liquid alloy droplet of eutectoid tempera‐ ture. Metal catalyst used in the mechanism was gold with 3 nm thickness. At high temperatures (about 850℃) alloy liquid droplets of Au-titanium form. On the other hand, during the thermal evaporation process, at first, titanium powder in the presence of oxygen atmosphere was converted toTiO2 powder (see Eq. 7) [44], then the TiO2 powder was reduced by graphite to

2.5 3 3.5 4 4.5 5

hν (eV)

We have used Ti metal and carbon instead of TiO2 in carbothermal reduction process. In fact, Ti in the tablet converts to TiO2 by oxygen in air (Eq. 7), and then TiO2 is reduced in carbothermal reduction process by carbon (Eq. 8). Indeed, carbon mainly was used for TiO2 reduction. Also there are some papers that have used titanium and carbon as a source material [26, 46]. Previously, TiO2 and C as a source material were examined in our lab but NWs did not grow. Amin et al. [21] also reported that continuous increase of the amount of oxygen further reduced the yield of TiO2 1D nanostructures. When the O2 concentration was high, the VLS growth was suppressed because Ti vapor could readily react with O2 before forming alloy

Since the adsorption of liquid is more than solid, the reactant components (vapor of titanium and oxygen) absorbed into the alloy. The liquid alloy acts as a preferred sink for arriving nanostructure components. The Ti and O2 are dissolved into the mediating alloy liquid droplets. When alloy droplets reach to saturation of reactive component, the 1D nanostructures grow by precipitation of TiO2 from the droplets in the interface liquid- solid to minimize free energy of alloy system. The 1D nanostructure grows in length by the mechanism until the Au is consumed or until the growth conditions are changed [47]. Here, by over an hour, furnace temperature began to reduce and growth condition of one-dimensional titanium dioxide

() ( ) () 2 2 Ti s +O g =TiO s 298-2( 500 ) (7)

TiO s +C s =Ti g +CO g 2 2 () () ( ) ( ) (8)

Usually droplets of catalyst are observed on the top of the nanostructures in a VLS- based growth system.

We have used Ti metal and carbon instead of TiO2 in carbothermal reduction process. In fact, Ti in the tablet converts to TiO2 by oxygen in air (Eq. 7), and then TiO2 is reduced in carbo‐ thermal reduction process by carbon (Eq. 8). Indeed, carbon mainly was used for TiO2 reduction. Also there are some papers that have used titanium and carbon as a source material [26, 46]. Previously, TiO2 and C as a source material were examined in our lab but NWs did not grow. Amin et al. [21] also reported that continuous increase of the amount of oxygen further reduced the yield of TiO2 1D nanostructures. When the O2 concentration was high, the VLS growth was suppressed because Ti vapor could readily react with O2 before forming alloy

Figure 7. (*αhν*)

curves versus photon energy (*hν*) for the samples.

Sample 1 Sample 2 Sample 3

Ti(s)+O2(g)=TiO2(s) (298-2500Ԩ) (7) TiO2(s)+C(s)=Ti(g)+CO2(g) (8)

produce Ti vapor as the following chemical reaction [45].

droplets with catalytic materials.

nanostructures is stopped.

droplets with catalytic materials.

**4. Growth mechanism** 

2

**3.4. Growth mechanism**

**Figure 7.** (*αhν*)

0

5

10

(αhν)

2

(eV/cm)2

15

20

25

60 Nanoparticles Technology

2

reduced by graphite to produce Ti vapor as the following chemical reaction [45].

Moreover, TiO2 is heavier than Ti and O2, so probably exists in lower height than them. Therefore, TiO2 can be absorbed into the nanostructure bodies more likely than their tips (alloy droplets), to be absorbed, so when TiO2 is more than Ti and O2, nanostructures will be shorter and thicker. Therefore, with increasing Ar gas flow, the crystallinity improves due to VLS growth method.

RamezaniSani et al. [26] have grown TiO2 nanowires by thermal evaporation on a silicon substrate with the major reflection along (101) direction. Their results indicate that a convenient gas flow for controlling diameter size of nanowires is about 100 sccm, While our substrate is quartz and the major reflection growth along is (110) direction. Also we see that with increasing gas flow from 60 to 170 sccm, nanowires become thinner and longer.

### **3.5. Microstructure characterization (theoretical background)**

Deviation from the ideal crystallinity such as finite crystallite size, strain (at the atomic level), and extended defects (stacking faults and dislocations) leads to broadening of the diffraction lines. Crystallite size is a measure of the size of coherently diffracting domains and is not generally the same as particle size due to polycrystalline aggregates. Strain is defined as change in length per unit length and is measured as the change in d spacing of a strained sample compared to the unstrained state [49].

According to Scherrer [50], the apparent crystallite size can be obtained as follows:

$$D\_{\vee} = \begin{pmatrix} \text{K}\lambda \text{ } / \text{ FWHM} \end{pmatrix} \cos \theta \tag{9}$$

where *K* is a constant close to unity, *θ* is the Bragg angle of the (*hkl*) reflection, and *λ* is the wavelength of X-rays used. *DV* is a volume-weighted quantity. Wilson in 1963 used integral breadth *β* instead of FWHM in Eq. (9).

The dependence of strain ε to line broadening is defined as follows [51]:

$$
\varepsilon = \begin{pmatrix} \beta \ \vdots \ 4 \end{pmatrix} \cot \theta = \beta \ / \begin{pmatrix} 4 \tan \theta \end{pmatrix} \tag{10}
$$

When the XRD patterns are adjusted to a combination of Cauchy and Gaussian line shapes, the Halder and Wagner approximation is better suited for obtaining the physical broaden‐ ing [52]:

$$
\beta\_f \approx \left( \left( \beta\_h \right)^2 - \left( \beta\_g \right)^2 \right) / \beta\_h \tag{11}
$$

where *βg*, *βh*, and *β<sup>f</sup>* refer to the integral breadths of the instrumental, observed, and measured (physical) profiles, respectively.

### *3.5.1. The Williamson–Hall integral breadth method*

Moreover, TiO2 is heavier than Ti and O2, so probably exists in lower height than them. Therefore, TiO2 can be absorbed into the nanostructure bodies more likely than their tips (alloy droplets), to be absorbed, so when TiO2 is more than Ti and O2, nanostructures will be shorter and thicker. Therefore, with increasing Ar gas flow, the crystallinity improves due to VLS

RamezaniSani et al. [26] have grown TiO2 nanowires by thermal evaporation on a silicon substrate with the major reflection along (101) direction. Their results indicate that a convenient gas flow for controlling diameter size of nanowires is about 100 sccm, While our substrate is quartz and the major reflection growth along is (110) direction. Also we see that with increasing

Deviation from the ideal crystallinity such as finite crystallite size, strain (at the atomic level), and extended defects (stacking faults and dislocations) leads to broadening of the diffraction lines. Crystallite size is a measure of the size of coherently diffracting domains and is not generally the same as particle size due to polycrystalline aggregates. Strain is defined as change in length per unit length and is measured as the change in d spacing of a strained sample

where *K* is a constant close to unity, *θ* is the Bragg angle of the (*hkl*) reflection, and *λ* is the wavelength of X-rays used. *DV* is a volume-weighted quantity. Wilson in 1963 used integral

> qb

( ) ( ) <sup>2</sup> <sup>2</sup> / *f hgh*

bbb æ ö ç ÷

where *βg*, *βh*, and *β<sup>f</sup>* refer to the integral breadths of the instrumental, observed, and measured

When the XRD patterns are adjusted to a combination of Cauchy and Gaussian line shapes, the Halder and Wagner approximation is better suited for obtaining the physical broaden‐

 q

> q

= = / 4 / 4 ( ) *cot tan* ( ) (10)

è ø » - (11)

) *cos* (9)

According to Scherrer [50], the apparent crystallite size can be obtained as follows:

*D K <sup>V</sup>* = / FWHM ( l

The dependence of strain ε to line broadening is defined as follows [51]:

 b

b

e

gas flow from 60 to 170 sccm, nanowires become thinner and longer.

**3.5. Microstructure characterization (theoretical background)**

compared to the unstrained state [49].

breadth *β* instead of FWHM in Eq. (9).

(physical) profiles, respectively.

ing [52]:

growth method.

62 Nanoparticles Technology

Williamson and Hall in 1953 proposed a method for resolving size and strain broadening. Williamson–Hall plots can be applied to a Gaussian profile as follows [53]:

$$\left(\left(\boldsymbol{\mathcal{B}}\_{f}\right)^{\*}\right)^{2} = \left(\boldsymbol{\mathcal{B}}\_{f}\cos\theta \mid \boldsymbol{\mathcal{A}}\right)^{2} = \left(\mathbf{1} \mid \left(\boldsymbol{D}\_{v}\right)^{2}\right) + 4\boldsymbol{\mathcal{E}}^{2}\boldsymbol{d}^{\*2} \tag{12}$$

A plot of (*β<sup>f</sup>* \* ) 2 against 4*d*\*2 gives a straight line. From the intercept and slope of line size, strain can be calculated.

### **3.6. Microstructure study in TiO2 nanowires**

X-ray diffraction (XRD), in which strong scattered intensity is observed at specific angles of scattering, has been widely used for decades in the determination of structural and micro‐ structural parameters. By the diffraction pattern, we have characterized some microstructural parameters like crystallite size, microstrain within the grain, dislocation density, and texture coefficient.

### **3.7. The Williamson–Hall integral breadth method**

From an analysis of XRD line broadening, the contributions due to crystallite size and lattice strain can be determined. Here, crystallite size and microstrain contributions are separated using the Williamson–Hall (W-H) method. Assuming contribution to the Bragg peak having both Lorentz and Gauss profile [54]:

$$\text{Lorentz fit} \; : \beta\_h = \beta\_f + \beta\_g \beta \cos \theta \;/\; \lambda = \; 1/D + \; 4 \& \sin \theta \;/\; \lambda \tag{13}$$

$$\text{Gauss fit:} \left\| \boldsymbol{\beta}\_{h}^{\,^{2}} = \boldsymbol{\beta}\_{f}^{\,^{2}} + \boldsymbol{\beta}\_{g}^{\,^{2}} \left( \boldsymbol{\beta} \cos \theta \,/\,\lambda \right)^{2} = \left( 1 / D \right)^{2} + 16 \boldsymbol{\varepsilon}^{2} \left( \sin \theta \,/\,\lambda \right)^{2} \tag{14}$$

By plotting *βcosθ*/*λ* vs *sinθ*/*λ* (Lorentz) or (*βcosθ*/*λ*) 2 vs (*sinθ*/*λ*) <sup>2</sup> (Gauss), one can estimate the microstrain from the line slope and inverse of the crystallite size from its *y*-intercept. Figures 9 and 10 indicate W-H curves for the samples in Lorentz and Gauss fits, respectively. The obtained values from W-H curves are given in Table 4. The plotted W-H curves indicate that strain and size contributions exist simultaneously in the samples.


**Table 4.** Crystallite size, mean microstrain, and dilatation values of the samples.

the samples.

curves for the samples in Lorentz and Gauss fits, respectively. The obtained values from W-H curves are given in Table 4. The plotted W-H curves indicate that strain and size contributions exist simultaneously in

Figure 9. Williamson-Hall curves for the samples: Lorentz fit. **Figure 9.** Williamson-Hall curves for the samples: Lorentz fit.

0.00007

Figure 9. Williamson-Hall curves for the samples: Lorentz fit.

Figure 10. Williamson-Hall curves for the samples: Gauss fit. **Figure 10.** Williamson-Hall curves for the samples: Gauss fit.

*θ*≡Δ*v*/*v* (15)

#### **3.8. Dilatation**

Sample *ε* (%) (Lorentz) *ε* (%) (Guass) *D* (nm) (Lorentz) *D* (nm) (Guass) Volume strain or dilatation is a change in volume that defines as [55]:

Volume strain or dilatation is a change in volume that defines as [55]:

$$
\theta \equiv \Delta v / v \tag{15}
$$

θ(%)

**8. Dilatation**  If three strain *ε*xx, *ε*yy, and *ε*zz have been treated as small quantities, the dilatation is [55]

Table 4. Crystallite size, mean microstrain, and dilatation values of the samples.

$$
\Theta = \hat{\varepsilon}\mathcal{E}/\hat{\varepsilon}\mathbf{x} + \hat{\varepsilon}\eta \;/\,\hat{\varepsilon}\mathcal{Y} + \hat{\mathcal{O}}\mathcal{L} \;/\,\hat{\varepsilon}\mathbf{z} = \mathfrak{e}\_{\infty} + \mathfrak{e}\_{\text{yy}} + \mathfrak{e}\_{\text{zz}} \tag{16}
$$

12

S1 0.3750.015 0.4480.020 14.280.38 14.140.42 0.232

12

The horizontal strain *ε*ll and vertical strain *ε*⊥ are calculated by the expression as follows: *ε*ll=(*a*epi-*a*0)/*a*0 and *ε*⊥=(*c*epi-*c*0)/*c*<sup>0</sup> [56], where a0 and c0 are the lattice constants of bulk TiO2 (*a*0= 4.587 Å, *c*0= 2.954 Å [57]). Here *ε*xx=*ε*yy=*ε*ll and *ε*zz= *ε*⊥; thus, the dilatation becomes.

$$
\theta = \mathcal{Z}\mathfrak{e}\_{\text{ll}} + \mathfrak{e}\_{\perp} \tag{17}
$$

The calculated dilatations are given in Table 4. The slopes of the W-H curve for the samples are negative, i.e., their strain is negative and there is a contraction of epilayer on substrates which means compressive stress. The different slopes represent different mean microstrains.

It is well known that strain has direct relation with elastic energy [55]. Evaporated particles from source due to their temperature have certain kinetic energy. In low Ar flow (60 sccm), Ar particles is lower than evaporated particles then collision between them is negligible. How‐ ever, with increasing Ar flow to 110 sccm, collision between Ar particles and evaporated particles increases therefore it could have reduced the evaporated particles energies and microstrain. However, with increasing Ar flow up to 170 sccm, the number of Ar particles enhanced compared to evaporated particles, so they can help to increase the evaporated particles movement. Therefore, their velocity and energy increase and also mircostrain increases.

Both dilatation and microsrain are negative and at the same order. We compare these two quantities, because the microstrain (obtained from W-H curve) is average of strain in volume not in specific direction. Thus, calculated strain from main definition must consider in volume.

We can compare volume strain (which found from lattice parameters) and microstrain (which found from XRD pattern) because the microstrain from the W-H curve does not have any particular direction. By comparing them, it reveals that they have approximately good agreement both in magnitude and sign.

### **3.9. Dislocation density**

12

S1 0.3750.015 0.4480.020 14.280.38 14.140.42 0.232 S2 0.2510.008 0.3590.012 15.160.52 15.580.53 0.228 S3 0.3610.012 0.4270.014 15.270.51 15.850.53 0.169

Sample *ε* (%) (Lorentz) *ε* (%) (Guass) *D* (nm) (Lorentz) *D* (nm) (Guass)

Table 4. Crystallite size, mean microstrain, and dilatation values of the samples.

q

Volume strain or dilatation is a change in volume that defines as [55]:

 h

Volume strain or dilatation is a change in volume that defines as [55]:

*θ*=∂ξ/∂x+∂η/∂y+∂ζ/∂z=*ε*xx+*ε*yy+*ε*zz (16)

If three strain *ε*xx, *ε*yy, and *ε*zz have been treated as small quantities, the dilatation is [55]

 z

If three strain *ε*xx, *ε*yy, and *ε*zz have been treated as small quantities, the dilatation is [55]

12

 ee

xx yy zz

 e=¶ ¶ +¶ ¶ +¶ ¶ = + + /// *xyz* (16)

Figure 10. Williamson-Hall curves for the samples: Gauss fit.

(sinθ/λ)2 (nm)-2

Figure 10. Williamson-Hall curves for the samples: Gauss fit.

0 0.02 0.04 0.06 0.08 0.1 0.12 0.14

0 0.02 0.04 0.06 0.08 0.1 0.12 0.14

(sinθ/λ)2 (nm)-2

curves for the samples in Lorentz and Gauss fits, respectively. The obtained values from W-H curves are given in Table 4. The plotted W-H curves indicate that strain and size contributions exist simultaneously in

> S1 ■ S2 ▲ S3 ♦

> S1 ■ S2 ▲ S3 ♦

S1 ■ S2 ▲ S3 ♦

> θ(%)

º D*v v*/ (15)

Figure 9. Williamson-Hall curves for the samples: Lorentz fit.

Figure 9. Williamson-Hall curves for the samples: Lorentz fit.

0 0.05 0.1 0.15 0.2 0.25

curves for the samples in Lorentz and Gauss fits, respectively. The obtained values from W-H curves are given in Table 4. The plotted W-H curves indicate that strain and size contributions exist simultaneously in

sinθ/λ (nm-1)

the samples.

64 Nanoparticles Technology

0

**Figure 9.** Williamson-Hall curves for the samples: Lorentz fit.

0

0

**Figure 10.** Williamson-Hall curves for the samples: Gauss fit.

0.00001 0.00002 0.00003 0.00004 0.00005 0.00006 0.00007

0.00001 0.00002 0.00003 0.00004 0.00005 0.00006 0.00007

**8. Dilatation** 

**3.8. Dilatation**

*θ*≡Δ*v*/*v* (15)

qx

(βcosθ/λ)2 (nm)-2

(βcosθ/λ)2 (nm)-2 0.002

the samples.

0.004

βcosθ/λ (nm-1)

0.006

0.008

0.01

**Lorentz fit**

**Gauss fit**

**Gauss fit**

To estimate the dislocation density, *ρD* (dislocation density due to domain size), *ρε* (dislocation density due to strain) and *ρ* (real dislocation density), the simple approach of Williamson and Smallman (1956) is followed, relating *ρ* to crystallite size *D* and strain <*ε*<sup>2</sup> >1/2 [58]:

$$
\rho\_D = \Im / \,\mathrm{D}^2 \tag{18}
$$

$$
\rho\_{\varepsilon} = 4 < \varepsilon^2 > /b^2 \tag{19}
$$

$$\rho = \left(\rho\_D, \rho\_\varepsilon\right)^{1/2} = 2\sqrt{3} < \varepsilon^2 >^{1/2} \ / \left(bD\right) \tag{20}$$

where *b* is the Burger's vector, which determines distance and direction of displacement. We assumed that the magnitude of Burger's vector is calculated by [(*a*<sup>2</sup> /4)+(*a*<sup>2</sup> /4)+(*c*<sup>2</sup> /4)]1/2, where *a* and *c* are lattice parameters. Dislocation densities in actual crystals depend on the preparation of the specimen, but can range from 102 to 1012 cm-2 [59], and the dislocation values of our samples are in the same range (Table 5).

According to Eq. 16, behavior of dislocation density is similar to microstrain, i.e., the value of dislocation density decreases from sample S1 to S2, 4and then increases from S2 to S3, since dislocation density and microstrain are related to displacement. It is noted that the results of Lorentz and Gauss fits for our samples are the same.


**Table 5.** Burger's vectors and dislocation densities of the samples.

### **3.10. Texture coefficient**

In order to clearly understand the preferred orientation of the samples, the texture coefficient TC(*hkl*) of each XRD pattern is calculated according to the following formula [60]:

$$\text{TCC} \{ hkl \} = \left[ I \{ hkl \} / I\_0 \{ hkl \} \right] / \left( 1 / N \right) \Sigma\_N \left[ I \{ hkl \} / I\_0 \{ hkl \} \right] \tag{21}$$

where *I*(*hkl*) is measured intensity of the (*hkl*) diffraction peak and *N* is the amount of crystal directions which is 11, 10, and 12 for samples 1, 2, and 3, respectively. If TC(*hkl*) equals to 1, there is no preferred orientation on (*hkl*) direction of the film. When this value is larger than 1, a preferred orientation exists along the (*hkl*) plane. The standard deviation TC(*hkl*) of the TC values from powder condition is [60]:

$$\sigma = \left[ \left( 1/N \right) \Sigma\_N \left\{ \text{TC} \left( hkl \right) \right\}^2 - 1 \right]^{1/2} \tag{22}$$

If *σ* equals to zero, the crystallite orientation of a film sample is the same as that of the related powder-shaped sample and there is no preferred direction. Therefore, *σ* value indicates the preferred orientation level of thin film samples [61]. The values of texture coefficient and its standard deviation for the samples are shown in Table 6. As expected, the preferred direction in the samples is (110). It is concluded that the different carrier gas flows do not influence the preferential orientation. The rutile (110) surface is the most stable crystal face and has lowest surface energy [62].


According to Table 6, the behavior of TC (110) is similar to dislocation density, i.e., dislocation helps the growth of (110) plane, which is the most stable plane in the rutile phase.

**Table 6.** Texture coefficient and its standard deviation for the samples.

In order to better determine the contribution of TC (110) to other TC (*hkl*), the proportion of its value to total value of all TC(*hkl*) was calculated, (i.e., TC(110)/∑TC(*hkl*)). The values obtained for samples 1, 2 and, 3 were 38.9, 38.5, and 35.3, respectively. It is observed that by increasing argon flow, the value of TC(110)/∑TC(*hkl*) is decreased. Since in the case of argon with 60 sccm flow, epitaxial growth has occurred more than the others, so it can make possible growth in the (110) plane, which is the most stable one in rutile TiO2. However, in further argon flow, the epitaxial growth becomes lower which causes the reduction of TC (110).

### **4. Conclusion**

where *b* is the Burger's vector, which determines distance and direction of displacement. We

and *c* are lattice parameters. Dislocation densities in actual crystals depend on the preparation

According to Eq. 16, behavior of dislocation density is similar to microstrain, i.e., the value of dislocation density decreases from sample S1 to S2, 4and then increases from S2 to S3, since dislocation density and microstrain are related to displacement. It is noted that the results of

**ρ (cm-2) × 1011 - Lorentz** *ρ (cm-2) × 1011- gauss b (Å)* **Sample**

In order to clearly understand the preferred orientation of the samples, the texture coefficient

where *I*(*hkl*) is measured intensity of the (*hkl*) diffraction peak and *N* is the amount of crystal directions which is 11, 10, and 12 for samples 1, 2, and 3, respectively. If TC(*hkl*) equals to 1, there is no preferred orientation on (*hkl*) direction of the film. When this value is larger than 1, a preferred orientation exists along the (*hkl*) plane. The standard deviation TC(*hkl*) of the TC

If *σ* equals to zero, the crystallite orientation of a film sample is the same as that of the related powder-shaped sample and there is no preferred direction. Therefore, *σ* value indicates the preferred orientation level of thin film samples [61]. The values of texture coefficient and its standard deviation for the samples are shown in Table 6. As expected, the preferred direction in the samples is (110). It is concluded that the different carrier gas flows do not influence the preferential orientation. The rutile (110) surface is the most stable crystal face and has lowest

TC(*hkl*) of each XRD pattern is calculated according to the following formula [60]:

( ) { ( )}

1 / TC 1 *N hkl* é ù

N

2.55 3.08 3.563 S1 1.61 2.24 3.562 S2 2.30 2.62 3.564 S3

TC(*hkl I hkl I hkl N I hkl I hkl* ) é ù ( ) / / 1/ <sup>0</sup> ( ) ( )åN 0 ( ) / ( ) ë û <sup>=</sup> (21)

1/2 <sup>2</sup>

ê ú ë û =å - (22)

/4)+(*a*<sup>2</sup>

to 1012 cm-2 [59], and the dislocation values of our

/4)+(*c*<sup>2</sup>

/4)]1/2, where *a*

assumed that the magnitude of Burger's vector is calculated by [(*a*<sup>2</sup>

of the specimen, but can range from 102

66 Nanoparticles Technology

samples are in the same range (Table 5).

Lorentz and Gauss fits for our samples are the same.

**Table 5.** Burger's vectors and dislocation densities of the samples.

**3.10. Texture coefficient**

values from powder condition is [60]:

surface energy [62].

s

In this study, single-crystalline rutile TiO2 nanowires were fabricated under different argon gas flows. The results indicated that TiO2 nanowires form in rutile phase and their preferred growth is in the (110) direction, since (110) plane in rutile phase has lowest surface energy. The band gap energy of TiO2 nanowires have increased due to the quantum confinement compared with bulk one. The wires become thinner and longer by increasing the argon flow. In 170 sccm, which has the least amount of oxygen, the average value of diameter and the length of TiO2 were respectively 72 and 1020 nm. The band gap of TiO2 nanowires has increased from 3.35 to 3.45 eV and their crystallinity has improved by increasing argon flux due to enhancement of VLS growth mechanism.

The advantage of our method is to fabricate TiO2 nanostructures without contamination and impurities. These materials can be used for photovoltaic and gas sensor applications due to their large surface to volume ratio.

### **Author details**

Reza Sabet Dariani\* and Zohreh Nafari Qaleh

\*Address all correspondence to: dariani@alzahra.ac.ir

Department of Physics, Alzahra University, Tehran, Iran

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## **Noble Metal Nanoparticles Prepared by Metal Sputtering into Glycerol and their Grafting to Polymer Surface**

Jakub Siegel, Alena Řezníčková, Petr Slepička and Václav Švorčík

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/61403

### **Abstract**

This chapter summarizes the basic information about elementary characteristics and tech‐ nology of preparation of noble metal nanoparticles. The introduction gives some basic in‐ formation on the history of development in this area, especially in terms of dimensionality of metal nanostructures and their possible applications.The first subsec‐ tion is devoted to the preparation and characterization of Au, Ag, Pt, and Pd nanoparti‐ cles (NPs), which were synthesized by direct metal sputtering in liquid propane-1,2,3, triole (glycerol). This method provides an interesting alternative to time-consuming, wetbased chemical synthesis techniques. Moreover, the suggested technique allows targeted variation of metal nanoparticle size, which is demonstrated in detail in case of AuNPs by variation of capturing media temperature. Nanoparticle size and shape were studied by transmission electron microscopy and dynamic light scattering. Optical properties of nanoparticle solution were determined by measuring its UV–Vis spectra. Concentration of metal nanoparticles in prepared solutions was determined by atomic absorption spec‐ troscopy. Antibacterial properties were tested against two common pollutants (*Escherichia coli*, a Gram-negative bacteria, and *Staphylococcus epidermidis*, a Gram-positive bacteria). In the presence of Ag nanoparticles, the growth of *E. coli* and *S. epidermidis* was complete‐ ly inhibited after 24 h. Any growth inhibition of *E. coli* was observed neither in the pres‐ ence of "smaller" (4–6 nm, AuNP4–6) nor "bigger" (9–12 nm, AuNP8–12) AuNPs during the whole examination period. AuNP4–6, but not AuNP8–12, was able to inhibit the growth of. *S epidermidis*. We also observed significant difference in biological activities of Pt and PdNPs. More specifically, PdNPs exhibited considerable inhibitory potential against both *E. coli* and *S. epidermidis*, which was in contrast to ineffective PtNPs. Our results indicate that Ag, Pd, and partially AuNPs have high potential to combat both Gram-positive and Gram-negative bacterial strains.The second subsection describes the effort to anchor met‐ al nanoparticles onto polyethyleneterephthalate (PET) carrier. Two different procedures of grafting of polymeric carrier, activated by plasma treatment, with Au and AgNPs are described. In the first procedure, the PET foil was grafted with biphenyl-4,4'-dithiol (BPD) and subsequently with Au and AgNPs. In the second one, the PET foil was grafted with Au and AgNPs previously coated by the same BPD. X-ray photoelectron spectrosco‐

© 2015 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.

py, Fourier transform infrared spectroscopy, and electrokinetic analysis were used for characterization of the polymer surface at different modification steps. Au and AgNPs were characterized by UV–Vis spectroscopy. In case of both types of nanoparticles, the first procedure was found to be more effective. It was proved that the BPD was chemical‐ ly bonded to the surface of the plasma-activated PET and it mediates subsequent grafting of the AuNPs.

**Keywords:** Nanoparticles, sputtering, antibacterial effect, chemical anchoring, polymer

### **1. Introduction**

Metal nanoparticles (NPs) nowadays represent the key material in vast range of industri‐ al applications ranging from rough industry to fine medical or biochemical utilization. Gold and silver NPs and their applications are currently of great interest. In recent years, rapid developments in nanotechnology have evolved to the point where inclusion of those entities into smart systems and related technologies has become quite easy. Nowadays, NPs and NPs-embedded materials cover broad spectrum of functional devices with promising properties, which are used in many fields including optoelectronics or catalysis [1]. They are employed in modern nanobiotechnology for the production of biosensors, visualiza‐ tion of cellular structures [2], and targeted transporting of drugs [3]. Colloidal nanoparti‐ cles are studied because they have unique physical and chemical properties that are different from "bulk" materials [4]. In all applications of nanotechnology, the size and shape of the nanoparticles play an important role [5]. Numerous studies [6] describe the unique properties of gold nanoparticles which can be used in applications such as fuel cells, environmental, and other chemical processes [7]. Various preparation techniques have been proposed to cover specific requirements of those man-made entities. Besides classical wetbased preparation techniques of which the pioneering one is that proposed by Turkevich [8], others are intensely studied which would open up new potential applications of NPs in the development of new technologies.

NPs are entities ranging in size between 1 to 100 nanometres. These materials behave as a whole unit in terms of transportation and properties that highly depend on their shape, size, and morphological substructure [9, 10]. Those materials are characterized by high surface areato-volume ratio, high chemical reactivity and physical affinity, as well as other interesting physico-chemical properties (optical, electrical, and magnetic) [11–14]. In addition, different NPs show various properties; likewise, different synthesis techniques of the same type of particle tend to alter or introduce new properties to the material. Consequently, the applica‐ tions of these materials are unlimited and significantly remarkable in human life and in industries. Each nanomaterial with its distinct properties independently finds its way into practical applications such as: drug delivery [15], waste treatment [16], lubricant and surfactant production [17], electrical and sensor devices [18, 19], as well as membrane fuel cells [20]. Noble metal NPs have been widely studied in the past few decades because of their applications in various areas including catalysis, gas and biological sensors, drug delivery, and antimicrobial agents. In specific, due to their biocompatibility and photo-optical distinctiveness, gold NPs (AuNPs) have proven to be very useful tools in several biomedical applications dealing with different aspects of detection, analysis, and treatment.

Recent advances in the study of noble metal NPs have led to their utilization in a number of very important applications including antimicrobial coatings, biosensing, diagnostic imaging, and cancer diagnosis and therapy [21–24]. This chapter surveys the various synthetic methods of broad spectrum of metallic NPs as well as most recent experimental studies focusing on the use of gold, silver, palladium, platinum, and ruthenium NPs in catalysis, food industry, and environmental applications. Moreover, the potent in vitro and in vivo antimicrobial and cytogenotoxic effects of various gold and silver nanomaterials are underlined. Finally, recent advances in functionalization of various solid substrates with gold and silver NPs as effective antimicrobial coatings and promising cell-stimulating agents are summarized. Despite their use in remediating numerous medical and health-related conditions, the efficacy and safety of many gold, silver, palladium, and platinum NPs are still under some scrutiny.

### **2. Nanoparticle synthesis**

py, Fourier transform infrared spectroscopy, and electrokinetic analysis were used for characterization of the polymer surface at different modification steps. Au and AgNPs were characterized by UV–Vis spectroscopy. In case of both types of nanoparticles, the first procedure was found to be more effective. It was proved that the BPD was chemical‐ ly bonded to the surface of the plasma-activated PET and it mediates subsequent grafting

**Keywords:** Nanoparticles, sputtering, antibacterial effect, chemical anchoring, polymer

Metal nanoparticles (NPs) nowadays represent the key material in vast range of industri‐ al applications ranging from rough industry to fine medical or biochemical utilization. Gold and silver NPs and their applications are currently of great interest. In recent years, rapid developments in nanotechnology have evolved to the point where inclusion of those entities into smart systems and related technologies has become quite easy. Nowadays, NPs and NPs-embedded materials cover broad spectrum of functional devices with promising properties, which are used in many fields including optoelectronics or catalysis [1]. They are employed in modern nanobiotechnology for the production of biosensors, visualiza‐ tion of cellular structures [2], and targeted transporting of drugs [3]. Colloidal nanoparti‐ cles are studied because they have unique physical and chemical properties that are different from "bulk" materials [4]. In all applications of nanotechnology, the size and shape of the nanoparticles play an important role [5]. Numerous studies [6] describe the unique properties of gold nanoparticles which can be used in applications such as fuel cells, environmental, and other chemical processes [7]. Various preparation techniques have been proposed to cover specific requirements of those man-made entities. Besides classical wetbased preparation techniques of which the pioneering one is that proposed by Turkevich [8], others are intensely studied which would open up new potential applications of NPs

NPs are entities ranging in size between 1 to 100 nanometres. These materials behave as a whole unit in terms of transportation and properties that highly depend on their shape, size, and morphological substructure [9, 10]. Those materials are characterized by high surface areato-volume ratio, high chemical reactivity and physical affinity, as well as other interesting physico-chemical properties (optical, electrical, and magnetic) [11–14]. In addition, different NPs show various properties; likewise, different synthesis techniques of the same type of particle tend to alter or introduce new properties to the material. Consequently, the applica‐ tions of these materials are unlimited and significantly remarkable in human life and in industries. Each nanomaterial with its distinct properties independently finds its way into practical applications such as: drug delivery [15], waste treatment [16], lubricant and surfactant production [17], electrical and sensor devices [18, 19], as well as membrane fuel cells [20]. Noble metal NPs have been widely studied in the past few decades because of their applications in

of the AuNPs.

in the development of new technologies.

**1. Introduction**

74 Nanoparticles Technology

One of the most common nanoparticle synthesis of noble metals are those developed by Brust– Schiffrin [25]. The technique is based on reduction of AuIII+ complex compound with NaBH4 stabilized by thiols. This technique enables preparation of highly stable particles with narrow distribution and possibility to control their sizes. Also, Ag particles have already been synthesized on the basis of reduction reactions [7,8]. Concurrently, due to specific require‐ ments on newly synthesized NPs, there have arisen other numerous techniques based on both wet and dry processes. Nowadays, most of the wet-based preparation techniques exploit confined reaction area inside which metal ions are reduced to zerovalent metals forming nanoparticles. The smaller and uniform the reaction cavity is, the smaller the dimension and narrower the size distribution of the resulting NPs. Different types of polymer-based agents have been investigated for preparation of metal and non-metal particles [26–31]. Metal NPs, semiconductors, inorganic oxides, and quantum dots synthesis were also investigated, especially in the confined space of surfactants, inorganic templates, and polymeric stabilizing agents [32–35]. The structure of the metal particles can be significantly affected by the prepa‐ ration technique [36–39]. Their size, shape, and geometry strongly depend on the applied technique. Therefore, it is very important to use a template that is modifiable in terms of functional groups and is adjustable with respect to its shape or size. Polymeric hydrogels are able to fulfil these parameters, and they can be also successfully prepared in various sizes with different functional groups [26, 40].

Butun and Sahiner [41] reported synthesis of a novel bulk hydrogel based on acrylamidogly‐ colic acid and its use as a template in the preparation of different metal NPs (Ag, Cu, Ni, Co). Nanoparticle synthesis is based on reduction of the captured metal ions inside the hydrogel. Hydrogels play a key role in tissue engineering because of the resemblance of the threedimensional network structure to the extracellular matrix environment [42–44]. The bulk hydrogels consist mostly of macropores (>50 nm) and even superpores with the pore sizes of up to several microns. The porosity can be modified by the degree of functional groups hydrophilicity or by the cross-linker used during preparation. The tune capability of the pore structure also provides specific advantages for the absorption and desorption of various species [45].

Completely new possibilities for the preparation of metal NPs open up utilization of PVD techniques. The most common capturing media for direct metal sputtering represent ion liquids (ILs) and their significantly less-toxic alternative – vegetable oils. First metal deposition onto liquid substrate was conducted in 1996 by the mean of radiofrequency (RF) magnetron sputtering. Silver was deposited into pure silicon oil [46]. The process went through two stages: (i) creation of percolation structure and (ii) nucleation of first metal clusters on the oil surface. With ongoing deposition, individual clusters spread out over the liquid surface and mutually interconnected forming continuous silver coverage floating on the oil surface. Such Ag layer exhibited considerable roughness with typical morphology. Sputtering process was strongly dependent on deposition conditions; when the sputtering power was held below 30 W, no silver layer was formed. Most probably, the drop of the sputtering power resulted in pene‐ tration of silver atoms into the liquid volume and formation of silver NPs (AgNPs). Never‐ theless, this finding was not confirmed by the authors since their intention was to prepare ultra-thin silver layers on the surface of liquid [47].

Vegetable oils represent perspective liquid substrates for sputtering of metals, since they are commonly available, cheap, biocompatible, and have the ability to stabilize metal NPs [48] (e.g., by the way of chemical reduction [49]). Thus, for the successful preparation of nanopar‐ ticles, no other toxic reducing agents are required, which is very desirable for further in vivo applications. Direct sputtering of Au into castor oil results in formation of biocompatible AuNPs. Application of higher voltage led to creation of larger particles, while the prolongation of deposition time had no impact on nanoparticle size, which was proved by transmission electron microscopy (TEM) and small angle Y-ray scattering (SAXS) [48]. Wender et al. [50] published interesting work correlating the properties of specific oil used with the deposition conditions with regards to formation of thin layers on the surface or in the volume of the oil. Authors showed that formation of AgNPs depends crucially on the applied voltage and specific surface coordination ability of the oil used (castor, canola, and kapron oil). Both lower voltage and weak coordination ability led to formation of continuous coatings on the oil surface. On contrary, higher voltage and strong coordination ability enable preparation of NPs. Higher voltage means higher diffusivity of adsorbed particles (metal atoms) on the surface of liquid which enables (i) particle penetration into the volume of liquid and (ii) anchoring of atoms/clusters to active functional groups. AgNPs were easily formed in castor oil (which contains solely hydroxyl groups) in wide range of applied voltage. On contrary, when using canola oil (which is predominately formed by unsaturated aliphatic chains) and kapron oil, thin silver layer was deposited on the oil surface at lower voltages, while AgNPs were formed only at higher discharge voltage.

Hydrogels play a key role in tissue engineering because of the resemblance of the threedimensional network structure to the extracellular matrix environment [42–44]. The bulk hydrogels consist mostly of macropores (>50 nm) and even superpores with the pore sizes of up to several microns. The porosity can be modified by the degree of functional groups hydrophilicity or by the cross-linker used during preparation. The tune capability of the pore structure also provides specific advantages for the absorption and desorption of various

Completely new possibilities for the preparation of metal NPs open up utilization of PVD techniques. The most common capturing media for direct metal sputtering represent ion liquids (ILs) and their significantly less-toxic alternative – vegetable oils. First metal deposition onto liquid substrate was conducted in 1996 by the mean of radiofrequency (RF) magnetron sputtering. Silver was deposited into pure silicon oil [46]. The process went through two stages: (i) creation of percolation structure and (ii) nucleation of first metal clusters on the oil surface. With ongoing deposition, individual clusters spread out over the liquid surface and mutually interconnected forming continuous silver coverage floating on the oil surface. Such Ag layer exhibited considerable roughness with typical morphology. Sputtering process was strongly dependent on deposition conditions; when the sputtering power was held below 30 W, no silver layer was formed. Most probably, the drop of the sputtering power resulted in pene‐ tration of silver atoms into the liquid volume and formation of silver NPs (AgNPs). Never‐ theless, this finding was not confirmed by the authors since their intention was to prepare

Vegetable oils represent perspective liquid substrates for sputtering of metals, since they are commonly available, cheap, biocompatible, and have the ability to stabilize metal NPs [48] (e.g., by the way of chemical reduction [49]). Thus, for the successful preparation of nanopar‐ ticles, no other toxic reducing agents are required, which is very desirable for further in vivo applications. Direct sputtering of Au into castor oil results in formation of biocompatible AuNPs. Application of higher voltage led to creation of larger particles, while the prolongation of deposition time had no impact on nanoparticle size, which was proved by transmission electron microscopy (TEM) and small angle Y-ray scattering (SAXS) [48]. Wender et al. [50] published interesting work correlating the properties of specific oil used with the deposition conditions with regards to formation of thin layers on the surface or in the volume of the oil. Authors showed that formation of AgNPs depends crucially on the applied voltage and specific surface coordination ability of the oil used (castor, canola, and kapron oil). Both lower voltage and weak coordination ability led to formation of continuous coatings on the oil surface. On contrary, higher voltage and strong coordination ability enable preparation of NPs. Higher voltage means higher diffusivity of adsorbed particles (metal atoms) on the surface of liquid which enables (i) particle penetration into the volume of liquid and (ii) anchoring of atoms/clusters to active functional groups. AgNPs were easily formed in castor oil (which contains solely hydroxyl groups) in wide range of applied voltage. On contrary, when using canola oil (which is predominately formed by unsaturated aliphatic chains) and kapron oil,

species [45].

76 Nanoparticles Technology

ultra-thin silver layers on the surface of liquid [47].

Much more aggressive media (from the toxicity point of view) for potential biomedical application of prepared NPs are already mentioned ILs. Those liquids can be defined as liquid electrolytes exclusively composed of ions. Currently, the most common ILs are pure substances or eutectic mixtures of organic–inorganic salts, which melts at temperatures below 100°C [51]. Among the most widespread ILs belong derivates of organic molecules, e.g., pyrrolidine, imidazole, and pyridine [52]. ILs' attractiveness in the process of NPs preparation consists in unusual physical–chemical properties. In addition to their extreme polarity, they also exhibit very low value of interface tension. Considering this, relatively high speed of nucleation can be achieved during the deposition process, which results in formation of small NPs without undesirable perturbations (Ostwald ripening1 ) [53]. Moreover, ILs very easily change their molecular arrangement so that to adapt to nucleation centres of emerging NPs, which leads to their stabilization. This high degree of adaptability consists in the presence of both hydrophilic and hydrophobic segments together with strong polarization force of ILs, which enables their orientation both perpendicular and parallel to the particles [13].

A large number of authors regard the value of surface tension and viscosity of liquid as critical in the process of preparation of metal NPs in liquid media. However, as evidenced in the work [52], in particular in the case of ILs, other parameters such as composition and the coordinating ability of the used liquid also play a significant role. If in addition the NPs growth occurs in the liquid volume, increase in the NPs size with increasing volume (size) of ILs anionic part can be expected [54–57]. Furthermore, it appears that the difference in size of the synthesized particles cannot be correlated with macroscopic characteristics of ILs, such as surface tension and viscosity [54]. In this connection, a very interesting medium for preparation of NPs by metal sputtering into liquid is pure propane-1,2,3-triole (glycerol) [31]. This substance combines physical properties comparable to those of ILs (i.e., low vapor pressure, thermally manageable viscosity, high coordination ability – per one glycerol molecule falls three hydroxyl groups) with minimal toxicity to living tissues. Biocompatibility of glycerol allows subsequent application of prepared NPs in bioengineering [58]. Additionally, one can easily control the size of emerging particles by the temperature of the capturing media (glycerol). Recent advances in preparation ways of selected noble metal NPs are summarized in Table 1.

Besides all other noble metals, silver in the form of AgNPs possesses strong antimicrobial specificity. The AgNPs use is of great importance, since several pathogenic bacteria have developed resistance to different types of antibiotics. That is why AgNPs have emerged up with varied medical applications reaching from silver-based dressings, silver-coated medicinal devices (nanogels, nanolotions), etc. The antibacterial effects of Ag salts have been noticed since antiquity [59], and Ag is currently used to control bacterial growth in a wide spectrum of applications, including dental work, catheters, and burn wounds [60, 61]. In fact, it is well

<sup>1</sup> Mass transfer from smaller particles to larger particles extending in real polydisperse systems as a result of higher vapour pressure or greater solubility of smaller particles. This reduction of the degree of dispersion proceeds until the conversion of disperse systems into the system with sufficiently coarse dispersion in which the differences in solubility or vapour pressure of particles of different sizes are a very slight, and the process speed is negligibly small.


known that Ag ions and Ag-based compounds are highly toxic to microorganisms, showing strong biocidal effects on as many as more than ten species of bacteria including *E. coli* [62, 63].

**Table 1.** Preparation techniques of selected noble metals with their typical size distributions.

Particle size reduction is an efficient and dependable tool to improve their bactericidal action and cytocompatibility. Nanotechnology techniques are of great assistance in the process of overcoming the limitations of size of nanoparticles [80]. Nanomaterials can be modified with the aim of better efficiency to facilitate their applications in various fields, e.g., bioscience and medicine.

### **3. Gold, silver, platinum, and palladium nanoparticles physically deposited into glycerol**

In this section, a brief overview on results published by our group in the field of noble metal nanoparticles synthesis and investigation of their antibacterial action is introduced. The antimicrobial effects of Au, Ag, Pt, and PdNPs were tested against representative microor‐ ganisms of public concern. When applying metal NPs into living tissues, fundamental attention must be paid to their synthesis process itself. For successive, non-invasive application of metal NPs into living microorganisms, our group has developed a novel, unconventional approach for the physical synthesis of gold, silver, platinum, and palladium NPs based on direct metal sputtering into the liquid media (glycerol) [21, 31]. Considering this, direct metal deposition into the glycerol seems to be a promising technique combining the advantages of non-toxic and environmentally friendly process completely omitting the usage of solvents or reduction agents compared to classical wet-based methods. Furthermore, the possibility of tailoring the NP size via controlling the temperature of capturing media as well as the investigation of the functions of NP size and metal type (gold and silver) on the antibacterial activity are demonstrated. The antibacterial properties of these NPs were tested against two common pollutants of *Escherichia coli* (*E. coli*, Gram-negative bacteria) and *Staphylococcus epidermidis* (*S. epidermidis*, Gram-positive bacteria) naturally occurring on the skin and mucous membranes of human and frequently involved in infections associated with a biofilm forma‐ tion. Prepared NPs were characterized by transmission electron microscopy with or without high-resolution (HRTEM, TEM) and UV–Vis spectroscopy.

known that Ag ions and Ag-based compounds are highly toxic to microorganisms, showing strong biocidal effects on as many as more than ten species of bacteria including *E. coli* [62, 63].

**Type of NPs Synthetic method NP size References**

• • Reduction in acidic environment 3–40 [66]

• pH control of Cu complexes 48–150 [72]

Particle size reduction is an efficient and dependable tool to improve their bactericidal action and cytocompatibility. Nanotechnology techniques are of great assistance in the process of overcoming the limitations of size of nanoparticles [80]. Nanomaterials can be modified with the aim of better efficiency to facilitate their applications in various fields, e.g., bioscience and

In this section, a brief overview on results published by our group in the field of noble metal nanoparticles synthesis and investigation of their antibacterial action is introduced. The antimicrobial effects of Au, Ag, Pt, and PdNPs were tested against representative microor‐ ganisms of public concern. When applying metal NPs into living tissues, fundamental attention must be paid to their synthesis process itself. For successive, non-invasive application of metal NPs into living microorganisms, our group has developed a novel, unconventional approach for the physical synthesis of gold, silver, platinum, and palladium NPs based on direct metal sputtering into the liquid media (glycerol) [21, 31]. Considering this, direct metal deposition into the glycerol seems to be a promising technique combining the advantages of non-toxic and environmentally friendly process completely omitting the usage of solvents or reduction agents compared to classical wet-based methods. Furthermore, the possibility of tailoring the NP size via controlling the temperature of capturing media as well as the investigation of the functions of NP size and metal type (gold and silver) on the antibacterial activity are demonstrated. The antibacterial properties of these NPs were tested against two

• Reduction process 2–40 [25, 67–70]

• • γ-Irradiation 3–30 [71]

• • • Biosynthesis 9–25 [73–77]

**Table 1.** Preparation techniques of selected noble metals with their typical size distributions.

medicine.

78 Nanoparticles Technology

**deposited into glycerol**

• Wet chemistry 20–60 [78, 79]

**3. Gold, silver, platinum, and palladium nanoparticles physically**

• Sol–gel microreactors 5–50 [64, 65] • • • • PVD into liquid substrate 2–5 [21, 31]

Au Ag Pd Pt Cu (nm)

TEM images of Au and Ag NPs are shown in Figure 1 [31]. Both Au and Ag NPs possess spherical shape with average diameter of about 3.5 and 2.4 nm, respectively. It was shown that the NPs size distribution and uniformity remain untouched even in diluted aqueous solutions up to glycerol:water ratio of 1:20. It is a well-known fact that the presence of glycerol –OH group stabilizes aqueous solution of AgNPs [81]. But the stabilization role of glycerol –OH groups in AuNPs is not yet fully understood. Inset in Figure 1A introduces TEM image of AuNPs, which were stored 3 months at ambient conditions. The agglomeration of particles can be seen, which is in accordance with UV–Vis analysis. Contrary to that, this phenomenon was not found in case of AgNPs.

**Figure 1.** TEM images of Au (A) and Ag (B) nanoparticles. Inset in (A) shows TEM image of Au nanoparticles after 3 months storage at laboratory conditions [31].

UV–Vis absorption spectra measured on Au and AgNPs solutions promptly after deposition (solid lines) and after 3 month storage at ambient conditions (dash lines) are introduced in Figure 2 and discussed in detail in [31]. UV–Vis is often used to estimate size, shape, and particle size distribution of colloidal solutions of metal nanoparticles as they show specific absorption band corresponding to localized surface plasmon resonance (LSPR) [82, 83]. Nevertheless, this method must be conducted under precise and reproducible conditions, since the position and intensity of LSPR absorption peak are affected by many factors, e.g., used solvents, type of stabilizers, and counter ions. The shape of the spectra exhibiting significant LSPR peaks corresponds to uniform nanoparticle colloidal solutions and also corresponds to the narrow size distribution [12, 84]. UV–Vis absorption spectrum of Au sample measured directly after deposition (solid line) shows diminishing LSPR peak with maximum at 512 nm. The shape and position of this peak suggest the particle sizes to be in interval from 2 to 10 nm [82], which corresponds with the sample's type (aqueous solution). Nevertheless, the influence of the AuNPs stabilization by –OH groups on the peak shift is not yet fully known. Mild red shift in spectrum after 3 month storage indicates both agglomeration effect and less narrow distribu‐ tioncomparedtooriginal sample.Narrowandhigh-intensitypeakoccurringat 400nmtogether with the pure yellow coloration of the colloid solution indicates the presence of zerovalent Ag. Oxidized AgNPs do not exhibit this peak in their UV–Vis spectrum [85]. Contrary to AuNPs, AgNPs solution is much more stable even after 3 months storage. The corresponding spec‐ trum exhibits almost the same peak position with pretty narrow distribution.

**Figure 2.** UV–Vis absorption spectra of Au and Ag aqueous nanoparticle solutions together with the photographs of corresponding NP solutions. Dash lines represent UV–Vis spectra of Au and Ag aqueous nanoparticle solutions stored for 3 months at laboratory conditions [31].

Targeted variation of metal nanoparticle size belongs to the great challenges of current science. Variation of capturing media temperature may provide a broad spectrum of nanoparticle diameters. The control of nanoparticle size through the capturing media temperature may be expressed upon a simple, generally known approach [86], which states that nanoparticle growth is governed by diffusion from the bulk of colloidal dispersion to the solution/particle interphase. When applied this approach, both AgNP4–6 and AuNP4–6 (indexes refer to average nanoparticle size) exhibit distinctive narrow peaks with maximum absorption at 400 and 520 nm, respectively (see Figure 3) [58]. These spectra are almost identical to those obtained in our former study on Ag and Au with average diameter of 3–5 nm, indicating stable solutions with minimal particle dispersion. Considerable broadening with pronounced red shift of absorption peek occurs at AuNP9–12. This phenomenon originates from both, increase in dimensions of individual Au particles and broader size distribution compared to AuNP4–6. Distinctive red coloration of AuNP4–6 turns into purple one, belonging to AuNP9–12, as the average size increases from ca 5 to 10 nm (see Figure 4).

intensity of LSPR absorption peak are affected by many factors, e.g., used solvents, type of stabilizers, and counter ions. The shape of the spectra exhibiting significant LSPR peaks corresponds to uniform nanoparticle colloidal solutions and also corresponds to the narrow size distribution [12, 84]. UV–Vis absorption spectrum of Au sample measured directly after deposition (solid line) shows diminishing LSPR peak with maximum at 512 nm. The shape and position of this peak suggest the particle sizes to be in interval from 2 to 10 nm [82], which corresponds with the sample's type (aqueous solution). Nevertheless, the influence of the AuNPs stabilization by –OH groups on the peak shift is not yet fully known. Mild red shift in spectrum after 3 month storage indicates both agglomeration effect and less narrow distribu‐ tioncomparedtooriginal sample.Narrowandhigh-intensitypeakoccurringat 400nmtogether with the pure yellow coloration of the colloid solution indicates the presence of zerovalent Ag. Oxidized AgNPs do not exhibit this peak in their UV–Vis spectrum [85]. Contrary to AuNPs, AgNPs solution is much more stable even after 3 months storage. The corresponding spec‐

trum exhibits almost the same peak position with pretty narrow distribution.

**Ag** 

for 3 months at laboratory conditions [31].

80 Nanoparticles Technology

**Figure 2.** UV–Vis absorption spectra of Au and Ag aqueous nanoparticle solutions together with the photographs of corresponding NP solutions. Dash lines represent UV–Vis spectra of Au and Ag aqueous nanoparticle solutions stored

Targeted variation of metal nanoparticle size belongs to the great challenges of current science. Variation of capturing media temperature may provide a broad spectrum of nanoparticle diameters. The control of nanoparticle size through the capturing media temperature may be expressed upon a simple, generally known approach [86], which states that nanoparticle growth is governed by diffusion from the bulk of colloidal dispersion to the solution/particle interphase. When applied this approach, both AgNP4–6 and AuNP4–6 (indexes refer to average nanoparticle size) exhibit distinctive narrow peaks with maximum absorption at 400 and 520 nm, respectively (see Figure 3) [58]. These spectra are almost identical to those obtained in our former study on Ag and Au with average diameter of 3–5 nm, indicating stable solutions with

**Au** 

**Figure 3.** UV–Vis absorption spectra of Ag (AgNP4–6) and Au (AuNP4–6, AuNP9–12) aqueous solutions of different nano‐ particle size together with the photographs of corresponding NP solutions (indexes refer to average NP diameter in nm) [58].

**Figure 4.** TEM images of nanoparticles of different composition and size AgNP4–6 (a), AuNP4–6 (b), and AuNP9–12 (c), together with detailed HRTEM image of corresponding area (d). Indexes refer to NP diameter in nm [58].

Figure 5

Generally, two major aspects determine whether the growth of NPs during metal sputtering into liquids will take a place rather than formation of thin film on the liquid surface [52]. Firstly, the kinetic energy of sputtered atoms or clusters must be sufficient to penetrate the surface of capturing media. Secondly, the first condition may be fulfilled only in the limited range of media viscosity; thus, the liquid media cannot be too viscous. When those conditions are met, the formation of NPs occurs. Figure 5 shows TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs. Image analysis of more than 500 particles from 10 different areas of TEM pictures proved that prepared PtNPs and PdNPs have average diameter of (1.7 ± 0.3) nm and (2.4 ± 0.4) nm, respectively. Apparent aggregation of particles is due to the preparation method of samples for TEM (HRTEM) analysis. Observed discrepancy in particle size is probably due to different sputtering yield of both metals (Pt ~ 1.27, Pd ~ 2.09) [87], which is in accordance with supposed growth mechanism causes considerably slower growth of Pt particles (lower concentration gradient). surface of capturing media. Secondly, the first condition may be fulfilled only in the limited range of media viscosity; thus, the liquid media cannot be too viscous. When those conditions are met, the formation of NPs occurs. Figure 5 shows TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs. Image analysis of more than 500 particles from 10 different areas of TEM pictures proved that prepared PtNPs and PdNPs have average diameter of (1.7 ± 0.3) nm and (2.4 ± 0.4) nm, respectively. Apparent aggregation of particles is due to the preparation method of samples for TEM (HRTEM) analysis. Observed discrepancy in particle size is probably due to different sputtering yield of both metals (Pt ~ 1.27, Pd ~ 2.09) [87], which is in accordance with supposed growth mechanism causes considerably slower growth of Pt particles (lower concentration gradient).

Generally, two major aspects determine whether the growth of NPs during metal sputtering

Firstly, the kinetic energy of sputtered atoms or clusters must be sufficient to penetrate the

TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs insets show **Figure 5.** TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs insets show histogram of corre‐ sponding particle size and distribution [21].

histogram of corresponding particle size and distribution [21].

Generally, two major aspects determine whether the growth of NPs during metal sputtering into liquids will take a place rather than formation of thin film on the liquid surface [52]. Firstly, the kinetic energy of sputtered atoms or clusters must be sufficient to penetrate the surface of capturing media. Secondly, the first condition may be fulfilled only in the limited range of media viscosity; thus, the liquid media cannot be too viscous. When those conditions are met, the formation of NPs occurs. Figure 5 shows TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs. Image analysis of more than 500 particles from 10 different areas of TEM pictures proved that prepared PtNPs and PdNPs have average diameter of (1.7 ± 0.3) nm and (2.4 ± 0.4) nm, respectively. Apparent aggregation of particles is due to the preparation method of samples for TEM (HRTEM) analysis. Observed discrepancy in particle size is probably due to different sputtering yield of both metals (Pt ~ 1.27, Pd ~ 2.09) [87], which is in accordance with supposed growth mechanism causes considerably slower growth of Pt

TEM

HRTEM

Pd

Pd

TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs insets show

**Figure 5.** TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs insets show histogram of corre‐

histogram of corresponding particle size and distribution [21].

sponding particle size and distribution [21].

Generally, two major aspects determine whether the growth of NPs during metal sputtering into liquids will take a place rather than formation of thin film on the liquid surface [52]. Firstly, the kinetic energy of sputtered atoms or clusters must be sufficient to penetrate the surface of capturing media. Secondly, the first condition may be fulfilled only in the limited range of media viscosity; thus, the liquid media cannot be too viscous. When those conditions are met, the formation of NPs occurs. Figure 5 shows TEM and HRTEM images of diluted aqueous solutions of Pt and PdNPs. Image analysis of more than 500 particles from 10 different areas of TEM pictures proved that prepared PtNPs and PdNPs have average diameter of (1.7 ± 0.3) nm and (2.4 ± 0.4) nm, respectively. Apparent aggregation of particles is due to the preparation method of samples for TEM (HRTEM) analysis. Observed discrepancy in particle size is probably due to different sputtering yield of both metals (Pt ~ 1.27, Pd ~ 2.09) [87], which is in accordance with supposed growth mechanism causes

particles (lower concentration gradient).

Pt

82 Nanoparticles Technology

Pt

Figure 5

considerably slower growth of Pt particles (lower concentration gradient).

**Figure 6.** Characterization of aqueous solutions with metal nanoparticles by dynamic light scattering. Green line refers to PtNPs and red one PdNPs [21].

Additionally to TEM analysis, DLS measurement was accomplished to determine average size and size distribution of prepared NPs (see Figure 6, previous page). DLS measurements indicates that both particles are about 7% bigger compared to TEM-based analysis, which is in a good agreement with the published results [88], since DLS technique provides hydrody‐ namic diameter. Moreover, observed size discrepancy is inherently caused by intensity weighted mean particle diameter in case of DLS contrary to number weighted diameter obtained by TEM analysis [11]. More importantly, DLS proves that prepared NPs are not agglomerated, which is of crucial importance for evaluation of their bactericidal effects.

**Figure 7.** UV–Vis absorption spectra of Pt and Pd aqueous nanoparticle solutions. Dash lines represent UV–Vis spectra of Pt and Pd aqueous nanoparticle solutions stored for 3 months at laboratory conditions [21].

Formation and stability of metal NPs in aqueous colloidal solution were confirmed by UV–Vis spectroscopy. Figure 7 shows the UV–Vis spectra recorded from the Pd and PtNPs immedi‐ ately after NPs preparation (solid lines) and after 3 months storage at laboratory conditions (dashed lines). It can be seen that apart from a distinct surface plasmon resonance (SPR) characteristic for other noble metal NPs (Au and Ag) [58], no distinctive SPR peaks occur in case of both Pd and PtNPs, which is in accordance with the earlier reports [89, 90]. Studied metal NPs exhibit increasing absorption toward shorter wavelengths [89, 91–93]. The course of dependence in case of both Pt and PdNPs indicates the presence of zerovalent metals in the solution [91]. Slight shift of the curves to lower absorbances at red region of visible spectrum (dashed lines) refers to mild agglomeration of individual particles in the solution after 3 months of storage. Bactericidal potency of various metal NPs has been previously reported [94–97]. It is known that their antimicrobial potency is governed by their size, composition, surface area, and charge [98, 99].

### **4. Antibacterial effects of noble metal nanoparticles**

Antibacterial properties of Ag, Au, Pd, and PtNPs were examined using two bacterial strains, *E. coli* and *S. epidermidis*, frequently involved in infections associated with a biofilm formation. Demonstration of inhibition effect of AgNP4–6 on *S. epidermidis* is shown in Figure 8. The growth of *E. coli* and *S. epidermidis* was completely inhibited in the presence of AgNP4–6 after 24 h (for both 6 and 24 h incubated samples) when compared to the control samples (bacteria incubated in glycerol or physiological saline solution [PBS]), see Table 2 and Figure 8 [58].

**Figure 8.** Photograph showing the inhibition effect of AgNP4–6 on *Staphylococcus epidermidis* (a) positive control of bac‐ terial colonies growing on agar plate, (b) bacterial sample treated with silver nanoparticles [58].

Noble Metal Nanoparticles Prepared by Metal Sputtering into Glycerol and their Grafting to Polymer Surface http://dx.doi.org/10.5772/61403 85


Formation and stability of metal NPs in aqueous colloidal solution were confirmed by UV–Vis spectroscopy. Figure 7 shows the UV–Vis spectra recorded from the Pd and PtNPs immedi‐ ately after NPs preparation (solid lines) and after 3 months storage at laboratory conditions (dashed lines). It can be seen that apart from a distinct surface plasmon resonance (SPR) characteristic for other noble metal NPs (Au and Ag) [58], no distinctive SPR peaks occur in case of both Pd and PtNPs, which is in accordance with the earlier reports [89, 90]. Studied metal NPs exhibit increasing absorption toward shorter wavelengths [89, 91–93]. The course of dependence in case of both Pt and PdNPs indicates the presence of zerovalent metals in the solution [91]. Slight shift of the curves to lower absorbances at red region of visible spectrum (dashed lines) refers to mild agglomeration of individual particles in the solution after 3 months of storage. Bactericidal potency of various metal NPs has been previously reported [94–97]. It is known that their antimicrobial potency is governed by their size, composition,

Antibacterial properties of Ag, Au, Pd, and PtNPs were examined using two bacterial strains, *E. coli* and *S. epidermidis*, frequently involved in infections associated with a biofilm formation. Demonstration of inhibition effect of AgNP4–6 on *S. epidermidis* is shown in Figure 8. The growth of *E. coli* and *S. epidermidis* was completely inhibited in the presence of AgNP4–6 after 24 h (for both 6 and 24 h incubated samples) when compared to the control samples (bacteria incubated

**Figure 8.** Photograph showing the inhibition effect of AgNP4–6 on *Staphylococcus epidermidis* (a) positive control of bac‐

terial colonies growing on agar plate, (b) bacterial sample treated with silver nanoparticles [58].

in glycerol or physiological saline solution [PBS]), see Table 2 and Figure 8 [58].

surface area, and charge [98, 99].

84 Nanoparticles Technology

**4. Antibacterial effects of noble metal nanoparticles**

**Table 2.** Inhibitory effect of silver and gold nanoparticles against Gram-negative (*E. coli*) and Gram-positive (*S. epidermidis*) bacteria (a time length of bacterial growth on LB plates after inoculation, b an empty circle indicates inhibition effect, c a full circle indicates positive growth, dphysiological saline solution, e a half full circle indicates 50 % growth) [58].

The growth inhibition of both bacterial strains was further maintained even after 30 and 48 h of growth, suggesting strong bactericidal activity of AgNP4–6. This is in contrary to generally accepted fact that Gram-negative bacteria are more susceptible to the inhibitory action of silver [100] caused probably by more facile penetration of silver through their thinner cell wall. On contrary, any growth inhibition was observed toward *E. coli* neither in the presence of AuNP4–6 nor in the presence of AuNP9–12 during the whole experiment. Surprisingly, AuNP4– 6, but not AuNP9–12, were able to inhibit the growth of *S. epidermidis* and the effect was preserved for whole tested period (48 h), when compared to control samples. The antibacterial action of Ag+ ions has been broadly reported so far [101], also AgNPs repeatedly showed their potency against bacteria [102]. Thus, our results (growth inhibition of both bacterial strains by AgNP4– 6) are in agreement with other groups results. Biocidal properties of AuNPs of similar size (5 nm) as prepared by our group (4–6 nm), were observed by Lima et al. [103]. Their AuNPs dispersed on zeolites were effective against Gram-negative *E. coli* and *Salmonella typhi* (90–95% growth inhibition). It has been reported that the antibacterial activity of AgNP is dependent on particle size and shape [104, 105]. Thus, it is very likely that also the size of AuNP plays significant role in the antimicrobial action.

Although there has been increased interest in studying bactericidal properties of noble metals over the past decades, only few have been reported on Pt and PdNPs. The inhibitory effect of one NPs concentration and four concentrations of bacteria was examined after 4 h in contact with NPs and 24 h of further post-incubation. From Figure 9, it is apparent that incubation of PdNPs with *E. coli* had pronounced effect on its growth up to 1 105 CFU per sample when compared to untreated control cells.

This remarkable antibacterial activity of Pd was diminished at higher bacterial concentrations ranging from 1 106 CFU per sample. Similar potency of PdNPs was observed against *S. epidermidis* (see Figure 9), even up to 1 106 CFU. Since PtNPs had very similar size as the first examined particles, the palladium ones, we expected comparable antimicrobial potential. Nevertheless, we detected only insignificant inhibition of bacterial growth induced by PtNPs at a concentration of 1 105 CFU of *E. coli* when compared to the control samples. Interestingly,

**Figure 9.** Inhibition effect of PtNPs and PdNPs toward bacterial strains of *E. coli* and *S. epidermidis* with different bacte‐ ria concentrations [21].

we did not observe any growth inhibition of *S. epidermidis* in the presence of PtNPs. More importantly, we even recorded slight stimulation of their growth.

### **5. Grafting of noble metal nanoparticles to plasma-activated polymer carriers**

Polymer surface can be considered as the phase boundary between the bulk polymer and the ambient environment. The application of polymeric materials depends significantly on the boundaries' conditions. The most of polymeric materials, nevertheless, have chemically inert surface, mostly hydrophobic. Non-treated polymer surfaces exhibit very often poor perform‐ ance in adhesion, coating, packaging, or colloid stabilization. [106–109].

During the past few decades, the modifications of polymer surface were under intensive research for applications in various fields of industry. Different improved techniques including chemical or physical enhancements were applied. Physical processes possess the advantage of surface interaction with radiation of electromagnetic waves, and oxidation with gases. On the contrary, chemical modifications use wet-treatment, such as blending. This part of manuscript is focused on the review of recent advances in surface grafting of polymer substrates which can be performed by the combination of physical and chemical processes [107, 110]. The advantages of grafting procedures in comparison to other procedures may be summarized to several points, such as controllable introduction of graft chains with a high density, the precise localization of graft chains to the surface with almost no change in bulk's surface properties. In contrast to physically coated polymer chains, covalent anchoring of graft chains onto a polymer surface avoids their delamination and assures their long-term chemical stability [111].

Minimally four strategies can be used for creation of functionalized polymer surfaces with gold nanostructures: (i) "grafting from procedure" is based on polymer chain growth from small initiators connected to AuNPs [112, 113]. This procedure leads to a very dense polymer brush construction. Other methods are based on radical polymerization (LRP) [114] and atomtransfer radical polymerization (SI-ATRP) [115, 116] (ii) "grafting to" enables one-pot synthesis of AuNPs stabilized by sulphur-containing polymers [10, 117], and generally produces a sparser coverage [118] (iii) physisorption using block copolymer micelles (nanoreactors), water-soluble polymers, or star block copolymers [119, 120] (iv) "post-modification of preformed AuNPs". In this method, AuNPs are generated in the first stage through conventional methods such as Brust−Schiffrin synthesis, followed by the exchange or modification with polymers [121, 122].

This section is dedicated to description of two procedures of Au and AgNPs grafting on plasma-modified PET developed by our group. In the first procedure, Au and AgNPs were deposited on PET, beforehand grafted with biphenyl-4,4'-dithiol (BPD). In the second proce‐ dure, the Au and AgNPs (both denoted with\* ), beforehand grafted with BPD, were coated onto plasma-modified PET (see Figure 10).

we did not observe any growth inhibition of *S. epidermidis* in the presence of PtNPs. More

**Figure 9.** Inhibition effect of PtNPs and PdNPs toward bacterial strains of *E. coli* and *S. epidermidis* with different bacte‐

Polymer surface can be considered as the phase boundary between the bulk polymer and the ambient environment. The application of polymeric materials depends significantly on the boundaries' conditions. The most of polymeric materials, nevertheless, have chemically inert surface, mostly hydrophobic. Non-treated polymer surfaces exhibit very often poor perform‐

During the past few decades, the modifications of polymer surface were under intensive research for applications in various fields of industry. Different improved techniques including chemical or physical enhancements were applied. Physical processes possess the advantage of surface interaction with radiation of electromagnetic waves, and oxidation with gases. On the contrary, chemical modifications use wet-treatment, such as blending. This part of manuscript is focused on the review of recent advances in surface grafting of polymer substrates which can be performed by the combination of physical and chemical processes [107, 110]. The advantages of grafting procedures in comparison to other procedures may be summarized to several points, such as controllable introduction of graft chains with a high density, the precise localization of graft chains to the surface with almost no change in bulk's surface properties. In contrast to physically coated polymer chains, covalent anchoring of graft

**5. Grafting of noble metal nanoparticles to plasma-activated polymer**

importantly, we even recorded slight stimulation of their growth.

ance in adhesion, coating, packaging, or colloid stabilization. [106–109].

**carriers**

ria concentrations [21].

86 Nanoparticles Technology

**Figure 10.** Scheme of PET modification: (A) plasma treatment, grafting with biphenyl-4,4'-dithiol (BPD) and then with AuNP or AgNP. (B) Plasma treatment, grafting with AuNP\* or AgNP\* covered with BPD [126].

In Figure 11, one can see the images of AuNPs obtained by TEM and HRTEM. It is evident that the AuNPs size is the same before and after BPD coating and their average diameter is 14.7 nm. Grafting with BPD does not lead to AuNPs aggregation, thanks to the presence of hydrophilic groups on the NP surface. It is known that the gold grows in a square *fcc* crystal structure with dominant plane (111) [123]. At higher HRTEM resolution (Figure 11C,E), it is seen that the AuNPs are in the form of decahedral particles. As a result of deficiency, real AuNPs should contain defects or be intrinsically strained [124]. In Figure 10E, gold atoms arranged in (111) plane are clearly seen. From Figure 11D, it is seen that some of the form of BPD-coated AuNPs differs from that of uncoated AuNPs and they are not in the form of decahedral particles due to the decrease of the surface tension [125]. It is apparent that the behavior of pristine AgNPs (AgNP in Figure 11F) and AgNPs modified with BPD (AgNP\* in Figure 11G) is significantly changed. AgNPs form uniform aggregates with nonspherical shape. On the contrary, AgNPs\* form structures with spherical character which are well dispersed. Surface modification technique with BPD application does not result in AgNPs aggregation due to the presence of hydrophilic (-SH) and hydrophobic groups the nanoparti‐ cle's surface. The calculated diameters of AgNP and AgNP\* were (55±10) and (45±10) nm, respectively [126]. BPD‐coated AuNPs differs from that of uncoated AuNPs and they are not in the form of decahedral particles due to the decrease of the surface tension [125]. It is apparent that the behavior of pristine AgNPs (AgNP in Figure 11F) and AgNPs modified with BPD (AgNP\* in Figure 11G) is significantly changed. AgNPs form uniform aggregates with nonspherical shape. On the contrary, AgNPs\* form structures with spherical character which are well dispersed. Surface modification technique with BPD application does not result in AgNPs aggregation due to the presence of hydrophilic (‐SH) and hydrophobic groups the

nanoparticle's surface. The calculated diameters of AgNP and AgNP\* were (55±10) and

In Figure 11, one can see the images of AuNPs obtained by TEM and HRTEM. It is evident that the AuNPs size is the same before and after BPD coating and their average diameter is 14.7 nm. Grafting with BPD does not lead to AuNPs aggregation, thanks to the presence of hydrophilic groups on the NP surface. It is known that the gold grows in a square *fcc* crystal structure with dominant plane (111) [123]. At higher HRTEM resolution (Figure 11C,E), it is

AuNPs should contain defects or be intrinsically strained [124]. In Figure 10E, gold atoms arranged in (111) plane are clearly seen. From Figure 11D, it is seen that some of the form of

**Figure 11.** Images of the Au nanoparticles. TEM: (A) AuNPs and (B) AuNPs coated with dithiol (AuNP\*). HRTEM: (C) AuNPs, (D) AuNPs coated with BPD (AuNP\*), and (E) detail of AuNP; TEM images of the silver nanoparticles: (F) AgNP and (G) AgNPs coated with BPD (AgNP\*) [125, 126].

Both (Au, Ag)NPs and (Au, Ag)NPs\* in solution were characterized by UV–Vis spectroscopy (see Figure 12). The absorption spectrum of the AuNPs solution (see Figure 12A) shows a maximum around 521 nm, corresponding to the transverse plasmon oscillation band. The wavelength of the surface plasmon resonance (SPR) [127] corresponds well with the average diameter estimated from TEM images. SPR absorption wavelength is known to increase with nanoparticle size [124]. From Figure 12, it is obvious that the absorption maximum of AuNP\* is shifted toward larger wave lengths. This finding is in accordance with the above mentioned fact that the BPD-coated AuNPs (AuNPs\*) exhibit enlarged surface [128]. The UV–Vis absorption spectra of AgNPs and AgNPs\* suspensions (Figure 12B) exhibit well-defined plasmon bands at about 455 and 413 nm, which is characteristic of nano-sized silver [31, 97, 129]. Observed wavelengths correspond well with average diameters of AgNPs estimated from TEM images (Figure 11F,G).

characteristic of nano‐sized silver [31, 97, 129]. Observed wavelengths correspond well with average diameters of AgNPs estimated from TEM images (Figure 11F,G). Noble Metal Nanoparticles Prepared by Metal Sputtering into Glycerol and their Grafting to Polymer Surface http://dx.doi.org/10.5772/61403 89

Images of the Au nanoparticles. TEM: (A) AuNPs and (B) AuNPs coated with dithiol (AuNP\*). HRTEM: (C) AuNPs, (D) AuNPs coated with BPD (AuNP\*), and (E) detail of AuNP; TEM images of the silver nanoparticles: (F) AgNP and (G) AgNPs coated with BPD

Both (Au, Ag)NPs and (Au, Ag)NPs\* in solution were characterized by UV–Vis spectroscopy (see Figure 12). The absorption spectrum of the AuNPs solution (see Figure 12A) shows a maximum around 521 nm, corresponding to the transverse plasmon oscillation band. The wavelength of the surface plasmon resonance (SPR) [127] corresponds well with the average diameter estimated from TEM images. SPR absorption wavelength is known to increase with nanoparticle size [124]. From Figure 12, it is obvious that the absorption maximum of AuNP\* is shifted toward larger wave lengths. This finding is in accordance with the above mentioned fact that the BPD‐coated AuNPs (AuNPs\*) exhibit enlarged surface [128]. The UV–Vis absorption spectra of AgNPs and AgNPs\* suspensions (Figure 12B) exhibit well‐defined plasmon bands at about 455 and 413 nm, which is

Figure 11

(AgNP\*) [125, 126].

Figure 11G) is significantly changed. AgNPs form uniform aggregates with nonspherical shape. On the contrary, AgNPs\* form structures with spherical character which are well dispersed. Surface modification technique with BPD application does not result in AgNPs aggregation due to the presence of hydrophilic (-SH) and hydrophobic groups the nanoparti‐

**Figure 11.** Images of the Au nanoparticles. TEM: (A) AuNPs and (B) AuNPs coated with dithiol (AuNP\*). HRTEM: (C) AuNPs, (D) AuNPs coated with BPD (AuNP\*), and (E) detail of AuNP; TEM images of the silver nanoparticles: (F)

Both (Au, Ag)NPs and (Au, Ag)NPs\* in solution were characterized by UV–Vis spectroscopy (see Figure 12). The absorption spectrum of the AuNPs solution (see Figure 12A) shows a maximum around 521 nm, corresponding to the transverse plasmon oscillation band. The wavelength of the surface plasmon resonance (SPR) [127] corresponds well with the average diameter estimated from TEM images. SPR absorption wavelength is known to increase with nanoparticle size [124]. From Figure 12, it is obvious that the absorption maximum of AuNP\* is shifted toward larger wave lengths. This finding is in accordance with the above mentioned fact that the BPD-coated AuNPs (AuNPs\*) exhibit enlarged surface [128]. The UV–Vis absorption spectra of AgNPs and AgNPs\* suspensions (Figure 12B) exhibit well-defined plasmon bands at about 455 and 413 nm, which is characteristic of nano-sized silver [31, 97, 129]. Observed wavelengths correspond well with average diameters of AgNPs estimated

In Figure 11, one can see the images of AuNPs obtained by TEM and HRTEM. It is evident that the AuNPs size is the same before and after BPD coating and their average diameter is 14.7 nm. Grafting with BPD does not lead to AuNPs aggregation, thanks to the presence of hydrophilic groups on the NP surface. It is known that the gold grows in a square *fcc* crystal structure with dominant plane (111) [123]. At higher HRTEM resolution (Figure 11C,E), it is seen that the AuNPs are in the form of decahedral particles. As a result of deficiency, real AuNPs should contain defects or be intrinsically strained [124]. In Figure 10E, gold atoms arranged in (111) plane are clearly seen. From Figure 11D, it is seen that some of the form of BPD‐coated AuNPs differs from that of uncoated AuNPs and they are not in the form of decahedral particles due to the decrease of the surface tension [125]. It is apparent that the behavior of pristine AgNPs (AgNP in Figure 11F) and AgNPs modified with BPD (AgNP\* in Figure 11G) is significantly changed. AgNPs form uniform aggregates with nonspherical shape. On the contrary, AgNPs\* form structures with spherical character which are well dispersed. Surface modification technique with BPD application does not result in AgNPs aggregation due to the presence of hydrophilic (‐SH) and hydrophobic groups the nanoparticle's surface. The calculated diameters of AgNP and AgNP\* were (55±10) and

were (55±10) and (45±10) nm,

cle's surface. The calculated diameters of AgNP and AgNP\*

(45±10) nm, respectively [126].

AgNP and (G) AgNPs coated with BPD (AgNP\*) [125, 126].

from TEM images (Figure 11F,G).

respectively [126].

88 Nanoparticles Technology

**Figure 12.** UV–Vis spectra of water solutions of (A) AuNPs (red) and AuNPs covered with BPD (AuNP\*, blue); (B) AgNPs (black) and AgNPs covered with BPD (AgNP\*, blue) [125, 126].

**Figure 13.** Differential FTIR spectra of plasma-treated PET grafted with BPD and then with AuNPs (PET/plasma/BPD/ AuNP, red) or plasma-treated PET and grafted with AuNPs covered with BPD (PET/plasma/AuNP\*, blue) [125].

Differential FTIR spectroscopy was used for the characterization of pristine PET, plasmatreated PET, and PET grafted with BPD and AuNPs using both techniques described above. Typical difference FTIR spectra are shown in Figure 13. The band at 2634 cm–1 corresponds to residual dithiol group. The band at 715 cm–1 corresponds to the –S–Au group. So that, the FTIR measurement proves that the AuNPs, deposited using both procedures examined, are chemically bound to the plasma-activated PET surface.

XPS analysis was used to monitor the change in the surface chemical composition after subsequent preparation steps. Element concentration of C, O, S, and Au on the surface of all samples (using both grafting techniques) is summarized in Table 3. After the plasma treatment, the PET surface is oxidized dramatically. The polymer oxidation after the plasma treatment and creation of oxygen-containing groups (C–O ~ 286.0 eV, O–C=O ~ 288.8 eV) at the polymer surface is well known [130]. Polymer chain is characterized by the C1s component centered at 284.7 eV due to C–C group. There is a change in the contribution of the shake-up satellite 290.9 eV due to the attachment of aromatic species (BPD) to the plasma-treated PET, see references [125, 131]. After grafting of plasma-treated PET with BPD, the oxygen concentration decreases dramatically. The attachment of BPD to the surface of PET (PET/BPD) was evidenced by the detection of the sulphur with concentration of 5.7 at. %. The concentration of bonded AuNP and AuNP\* varies from 0.1 to 0.02 at. %, respectively. Atomic concentrations of C(1s), O(1s), S(2p), and Ag(3d) in pristine, plasma-modified PET and after grafting with BPD and silver nanoparticles are summarized in Table 3. After the next grafting procedure realized with AgNP and AgNP\* particles, the concentration of sulphur decreased and silver was confirmed for the PET/plasma/BPD/AgNP samples, which indicates the presence of AgNPs on the surface of treated PET. In the case of PET/plasma/AgNP\* samples, the silver concentration is presum‐ ably below the detection limit of XPS method. However, the evidence of successful AgNP\* attachment was confirmed by the presence of sulphur on the treated surface.


**Table 3.** Atomic concentration of C, O, S, and Au measured with XPS in surface polymer layer for: pristine (PET), plasma exposed (PET/plasma), plasma exposed and grafted with BPD (PET/plasma/BPD) and then grafted with AuNP (PET/plasma/BPD/AuNP) and grafted with AuNP\* (PET/plasma/AuNP\*) 4 days after the exposure [125].

It is known that the thiol groups bond to double bonds created on surface by plasma treatment by free radical chain mechanism [132]. Radicals created after plasma activation of PET are located mainly at heteroatoms (e.g., oxygen of ester or hydroxyl group). Therefore, these radicals are stable and do not stabilize by creation of double bonds on surface [133]. In our previous study [134], it has been proved the amount of free radicals created on PET surface modified by plasma is very low. That is the reason why more double bonds are not available on PET, and therefore the amount of bonded thiol groups is not much higher [125, 126].

### **6. Conclusions**

This chapter attempts to provide a comprehensive insight into the problems of new techniques of preparation of noble metal NPs. It summarizes the basic information about the elementary characteristics and technology of preparation of those man-made materials. The introduction gives some basic information on the history of development in this area, especially in terms of dimensionality of metal nanostructures and their possible applications. General overview is supported by the specific research of the author team in this field. Presented findings support the hypothesis that Au, Ag, Pd, and PtNPs can be prepared in a simple and cost-effective manner and are suitable for formulation of new types of bactericidal materials. Presented approach is based on temperature control of capturing media during the sputtering deposition process into propane-1,2,3-triol, allowing targeted variation of NPs size. Bactericidal action of such synthesized NPs was examined against two common pollutants (*E. coli* and *S. epidermi‐ dis*). While AgNP4–6 exhibit strong potential to completely inhibit both bacterial strains after 24 h, AuNPs showed pronounced inhibition selectivity regarding the specific NP size. Regardless of AuNP size, any growth inhibition of *E. coli* was observed. Contrary to that, AuNP4–6 were able to inhibit the growth of *S. epidermidis*. We have also illustrated a significant difference in biological activities of Pt and PdNPs. More specifically, PdNPs exhibited considerable inhibitory potential against both *E. coli* and *S. epidermidis*, which was in contrast to ineffective PtNPs. Those findings indicate that palladium has high potential to combat both Grampositive and Gram-negative bacterial strains. Understanding of nanoparticle–bacteria inter‐ actions is a key issue in their potential applications in clinics on everyday basis. Observed discrepancy in bactericidal action between Pt and Pd nanoparticles could be attributed to different size of individual particles. Finally, we have introduced procedures for PET surface coating with noble metal nanoparticles. The techniques were based on activation of PET surface with Ar plasma discharge and consequent application of dithiol as binding reagent between Au and Ag nanoparticles and plasma-treated PET. The introduction of gold and silver nanoparticle–thiol layer on the activated PET surface was confirmed with XPS analysis. Higher concentration of Au and AgNPs was achieved by deposition on PET grafted beforehand with dithiol.

### **Acknowledgements**

XPS analysis was used to monitor the change in the surface chemical composition after subsequent preparation steps. Element concentration of C, O, S, and Au on the surface of all samples (using both grafting techniques) is summarized in Table 3. After the plasma treatment, the PET surface is oxidized dramatically. The polymer oxidation after the plasma treatment and creation of oxygen-containing groups (C–O ~ 286.0 eV, O–C=O ~ 288.8 eV) at the polymer surface is well known [130]. Polymer chain is characterized by the C1s component centered at 284.7 eV due to C–C group. There is a change in the contribution of the shake-up satellite 290.9 eV due to the attachment of aromatic species (BPD) to the plasma-treated PET, see references [125, 131]. After grafting of plasma-treated PET with BPD, the oxygen concentration decreases dramatically. The attachment of BPD to the surface of PET (PET/BPD) was evidenced by the detection of the sulphur with concentration of 5.7 at. %. The concentration of bonded AuNP and AuNP\* varies from 0.1 to 0.02 at. %, respectively. Atomic concentrations of C(1s), O(1s), S(2p), and Ag(3d) in pristine, plasma-modified PET and after grafting with BPD and silver nanoparticles are summarized in Table 3. After the next grafting procedure realized with AgNP and AgNP\* particles, the concentration of sulphur decreased and silver was confirmed for the PET/plasma/BPD/AgNP samples, which indicates the presence of AgNPs on the surface of treated PET. In the case of PET/plasma/AgNP\* samples, the silver concentration is presum‐ ably below the detection limit of XPS method. However, the evidence of successful AgNP\*

attachment was confirmed by the presence of sulphur on the treated surface.

PET 72.5 27.5 - - PET/plasma 29.0 71.0 - - PET/plasma/BPD 75.4 18.9 5.7 - PET/plasma/BPD/AuNP 72.6 24.6 2.7 0.1 PET/plasma/AuNP\* 72.2 26.4 1.4 0.02

**Table 3.** Atomic concentration of C, O, S, and Au measured with XPS in surface polymer layer for: pristine (PET), plasma exposed (PET/plasma), plasma exposed and grafted with BPD (PET/plasma/BPD) and then grafted with AuNP

It is known that the thiol groups bond to double bonds created on surface by plasma treatment by free radical chain mechanism [132]. Radicals created after plasma activation of PET are located mainly at heteroatoms (e.g., oxygen of ester or hydroxyl group). Therefore, these radicals are stable and do not stabilize by creation of double bonds on surface [133]. In our previous study [134], it has been proved the amount of free radicals created on PET surface modified by plasma is very low. That is the reason why more double bonds are not available on PET, and therefore the amount of bonded thiol groups is not much higher [125, 126].

This chapter attempts to provide a comprehensive insight into the problems of new techniques of preparation of noble metal NPs. It summarizes the basic information about the elementary

(PET/plasma/BPD/AuNP) and grafted with AuNP\* (PET/plasma/AuNP\*) 4 days after the exposure [125].

C(1s) O(1s) S(2p) Au(4f)

**Sample Element concentration (at. %)**

**6. Conclusions**

90 Nanoparticles Technology

Financial support of this work under the GACR projects Nos. 14-18131S, 14-18149P, and 13-06609S is gratefully acknowledged.

### **Author details**

Jakub Siegel\* , Alena Řezníčková, Petr Slepička and Václav Švorčík

\*Address all correspondence to: jakub.siegel@vscht.cz

Department of Solid State Engineering, University of Chemistry and Technology Prague, Prague, Czech Republic

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### **The Development of Smart, Multi-Responsive Core@Shell Composite Nanoparticles**

Bo Sang Kim and T. Randall Lee

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/61262

### **Abstract**

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The unique optical, magnetic, and electronic properties of metal nanoparticles (NPs) give rise to photothermal, therapeutic, and electronic device applications, correspond‐ ingly. On the other hand, the limited range of the properties of simple spherical metal NPs has complicated their ability to function in many of these applications. Therefore, this chapter starts by reviewing a specific type of NP that can be classified into three main groups: silica coated with metal (silica@metal) NPs, metal coated with silica (metal@silica) NPs, and other similar forms of core@shell structures. The objective of this review is to introduce the concept of multi-responsive core@shell nanoparticles. More specifically, this chapter highlights "smart" core@shell composite NPs having multiple response mechanisms (e.g., temperature, light, and/or an applied magnetic field) due to the ability of these systems to perform a task by remotely responding to stimuli. Additionally, hydrogel-coated metal@silica NPs, with the ability to store drugs in a mesoporous silica (m-silica) interlayer, are examined because these nanomaterials potentially provide substantial advantages for carrying cargos to targeted sites. To demonstrate this capability, we examine recent research that provided initial tests of composite NPs with a pH- and temperature-responsive hydrogel coating, in‐ cluding the application of an underlying m-silica interlayer to improve the capacity of these NPs to load and release small molecules.

**Keywords:** Nanoshells, core@shell nanoparticles, surface plasmon resonance, smart multi-responsive nanoparticles

### **1. Introduction**

Metal nanoparticles (NPs) provide optical, magnetic, and electronic properties which are different from the corresponding bulk metal materials, leading to photothermal, therapeutic, and electronic device applications, respectively. Recent research efforts have explored NPs

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having complex compositions with controllable sizes and morphologies. Special attention has been given to metal NPs for use as optical materials (i.e., using palladium [1, 2], silver [3-12], copper [13-18], and particularly gold [19-31]). This is due to the ease of fabrication of these metal nanoparticles and their unique physical properties [32-34], such as the ability to absorb or scatter light [25, 35, 36]. However, the narrow scope of simple spherical metal NPs' optical properties has limited their use in practical applications. Thus, investigators have pursued new classes of nanostructured materials such as triangular prisms [37-40], disks [41, 42], nanorods [43-46], nanocubes [47-51], and NPs coated with a shell (core@shell NPs) [52-58], to overcome the restrictions on the optical properties associated with simple spherical metal NPs. Among these, core@shell NPs represent a promising nanoscale tool for biomedical research [59-63]. A number of metals have been used as the shell for these NPs (as detailed in Table 1), but silica has proven to be the favorite core for such structures [64-76].


**Table 1.** References for core@shell nanoparticles with silica cores and metal shells.

Recent research has also centered on the capacity of core@shell NPs containing metal/metal oxide cores to respond to an external stimulus (e.g., a magnetic field or near IR light) and to affect their local surroundings, resulting in their use in diagnostic, therapeutic, and drug delivery applications utilizing either a polymer coating or a more complex shell design [77-80]. The development of these more sophisticated NP structures has been presaged by a number of simpler particle architectures.

### **2. Types of Core@Shell nanoparticles**

Core@shell NPs can be categorized according to their material properties (e.g., dielectric, semiconductor, etc.) and form (e.g., single-layered shell, multilayered shell, etc.). In this chapter, the core or shell materials in a core@shell NP are considered in terms of their consti‐ tution or shape. For this report, the core@shell NPs of interest are classified into three main groups: (i) silica@metal NPs; (ii) metal@silica NPs; and (iii) other forms of core@shell NPs.

*Silica@Metal Nanoparticles.* Since the publication of the "Stöber Method" of preparing silica particles, various metals and metal oxide shell materials have been deposited on silica spheres [81]. Because of the unique qualities of gold, silica-based core@shell NPs with gold as the shell material became the focus of efforts to produce the first of these structures in the late 1990's [64, 68, 82]. A gold coating on the silica core provides improved biocompatibility, photonic energy absorption, catalytic properties, chemical stability, bio-affinity (through greater diversity in the functionalization of the surface), and tunable optical properties [64]. Similar to the sizedependent color of pure gold NPs, the optical response of gold nanoshells (GNSs) depends dramatically on the relative size of the core NP as well as the thickness of the gold shell [67, 83]. Importantly, the surface-enhanced Raman scattering (SERS) activity of GNSs can be increased or decreased by decreasing or increasing the thickness of the gold shell material [84, 85]. By adjusting the relative core size and shell thickness, an intense light absorption associ‐ ated with the gold in the GNSs can be varied across a broad range of the optical spectrum, spanning the visible and the near-infrared spectral regions. This phenomenon is known as surface plasmon resonance (SPR) and its impact upon the extinction spectra of GNSs is illustrated in Figure 1 with regards to changes in shell thickness.

having complex compositions with controllable sizes and morphologies. Special attention has been given to metal NPs for use as optical materials (i.e., using palladium [1, 2], silver [3-12], copper [13-18], and particularly gold [19-31]). This is due to the ease of fabrication of these metal nanoparticles and their unique physical properties [32-34], such as the ability to absorb or scatter light [25, 35, 36]. However, the narrow scope of simple spherical metal NPs' optical properties has limited their use in practical applications. Thus, investigators have pursued new classes of nanostructured materials such as triangular prisms [37-40], disks [41, 42], nanorods [43-46], nanocubes [47-51], and NPs coated with a shell (core@shell NPs) [52-58], to overcome the restrictions on the optical properties associated with simple spherical metal NPs. Among these, core@shell NPs represent a promising nanoscale tool for biomedical research [59-63]. A number of metals have been used as the shell for these NPs (as detailed in Table 1),

Recent research has also centered on the capacity of core@shell NPs containing metal/metal oxide cores to respond to an external stimulus (e.g., a magnetic field or near IR light) and to affect their local surroundings, resulting in their use in diagnostic, therapeutic, and drug delivery applications utilizing either a polymer coating or a more complex shell design [77-80]. The development of these more sophisticated NP structures has been presaged by a number

Core@shell NPs can be categorized according to their material properties (e.g., dielectric, semiconductor, etc.) and form (e.g., single-layered shell, multilayered shell, etc.). In this chapter, the core or shell materials in a core@shell NP are considered in terms of their consti‐ tution or shape. For this report, the core@shell NPs of interest are classified into three main groups: (i) silica@metal NPs; (ii) metal@silica NPs; and (iii) other forms of core@shell NPs.

*Silica@Metal Nanoparticles.* Since the publication of the "Stöber Method" of preparing silica particles, various metals and metal oxide shell materials have been deposited on silica spheres [81]. Because of the unique qualities of gold, silica-based core@shell NPs with gold as the shell material became the focus of efforts to produce the first of these structures in the late 1990's [64, 68, 82]. A gold coating on the silica core provides improved biocompatibility, photonic energy absorption, catalytic properties, chemical stability, bio-affinity (through greater diversity in

[57, 64-68] [54, 69-73] [53, 74] [53, 54, 75, 76]

but silica has proven to be the favorite core for such structures [64-76].

**Core@Shell Nanoparticles References**

**Table 1.** References for core@shell nanoparticles with silica cores and metal shells.

gold@silica silver@silica palladium@silica gold alloy@silica

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of simpler particle architectures.

**2. Types of Core@Shell nanoparticles**

**Figure 1.** Illustration of the optical resonances of various sizes of silica@gold composite particles (gold nanoshells) as a function of their shell thickness. Reproduced with permission from reference [67].

The development of these gold shells has been followed by a number of alternative metal and metal oxide coatings, broadening the number of applications for such core@shell structures [53, 54, 69, 74]. These data indirectly indicate that GNSs with a silica core (SiO2@Au NPs) may be useful for biomedical imaging, cancer treatment, and optical materials in future applications owing to the ability of certain wavelengths of light to penetrate human tissue.

*Metal@ Silica Nanoparticles.* Using a silica coating as the shell on a metal core to form NPs with core@shell architectures offers several benefits. First, a silica shell diminishes the bulk metal conductivity and improves the chemical and physical stability of the core component [86]. Second, silica particles produced via the Stöber Method have been shown to be chemically biocompatible [87, 88]. And silica can be used to block the core surface from making contact with a biological environment without disrupting key phenomena involving the core surface, such as light absorption. This is an important feature for metals/metal oxides that have been shown to be cytotoxic. Third, the silica coating can be used to modulate the position and intensity of the surface plasmon absorption band due to both the optical transparency of silica at technologically important wavelengths of light and the shift in these bands associated with a small change in the metal's optical properties because of the material in contact with their surfaces [89]. And the fourth point is that an outer silica shell on a metal nanoshell can help stabilize the metal shell when it is generating intense localized heat in response to intense light absorption [90, 91]. Therefore, scientists have recently focused more on silica coatings as a shell for a variety of core materials such as metals [92-99] and metal oxides [100-103], than on any other material. Many research groups have investigated the use of silica on coinage metals such as gold [92, 93] and silver [98]. Silica-coated gold or silver NPs are synthesized by using a slightly modified Stöber (or sol-gel) Method. The resulting coating does not interfere with the intensity of the light absorption for targeted wavelengths, and only produces a minor shift in the absorption band toward the higher wavelength region, as compared to an uncoated NP, as shown in Figure 2 [104]. This coating method has been adjusted further to control the uniformity of the thickness of the silica layer on gold NP cores [105].

**Figure 2.** UV-visible absorption spectra of aqueous dispersions that contained gold@silica NPs with different shell thicknesses (t). The gold cores were 50 nm in diameter for all samples. Reproduced with permission from reference [104].

Figure 2 provides the absorption spectra of aqueous dispersions of gold@silica (Au@SiO2) NP colloids with different shell thicknesses. The characteristic surface plasmon peak of the core of these NPs is ~528 nm before adding the silica coating to the gold cores. After coating with silica, the metal's plasmon peak red-shifted to ~540 nm. This is because the refractive index of silica (n = 1.52) is slightly higher than that of water (n = 1.31), the solvent used to suspend the nanoparticles for collecting spectroscopic data [106]. The optical intensity of these Au@SiO2 NPs was increased correspondingly when thicker silica shells were created. But the specific positioning of this peak was not sensitive to the change in the silica coating thickness.

such as light absorption. This is an important feature for metals/metal oxides that have been shown to be cytotoxic. Third, the silica coating can be used to modulate the position and intensity of the surface plasmon absorption band due to both the optical transparency of silica at technologically important wavelengths of light and the shift in these bands associated with a small change in the metal's optical properties because of the material in contact with their surfaces [89]. And the fourth point is that an outer silica shell on a metal nanoshell can help stabilize the metal shell when it is generating intense localized heat in response to intense light absorption [90, 91]. Therefore, scientists have recently focused more on silica coatings as a shell for a variety of core materials such as metals [92-99] and metal oxides [100-103], than on any other material. Many research groups have investigated the use of silica on coinage metals such as gold [92, 93] and silver [98]. Silica-coated gold or silver NPs are synthesized by using a slightly modified Stöber (or sol-gel) Method. The resulting coating does not interfere with the intensity of the light absorption for targeted wavelengths, and only produces a minor shift in the absorption band toward the higher wavelength region, as compared to an uncoated NP, as shown in Figure 2 [104]. This coating method has been adjusted further to control the

**Figure 2.** UV-visible absorption spectra of aqueous dispersions that contained gold@silica NPs with different shell thicknesses (t). The gold cores were 50 nm in diameter for all samples. Reproduced with permission from reference

Figure 2 provides the absorption spectra of aqueous dispersions of gold@silica (Au@SiO2) NP colloids with different shell thicknesses. The characteristic surface plasmon peak of the core of these NPs is ~528 nm before adding the silica coating to the gold cores. After coating with silica, the metal's plasmon peak red-shifted to ~540 nm. This is because the refractive index of silica (n = 1.52) is slightly higher than that of water (n = 1.31), the solvent used to suspend the

[104].

106 Nanoparticles Technology

uniformity of the thickness of the silica layer on gold NP cores [105].

In the report by Lu *et al.* referenced above, the silica-coated gold NPs were assembled into ordered arrays [104]. The transmission spectra collected from these lattices of Au@SiO2 NPs are shown in Figure 3. All samples were wet with the hollow spaces between NPs being completely filled with water, when these spectra were measured. The incident light was kept vertical to the (111) planes of these face-center-cubic lattices (Au).

**Figure 3.** Transmission spectra taken from the arrays produced from Au@SiO2 NPs with different shell thicknesses (t). The incident light was perpendicular to the (111) planes of these face-center-cubic crystalline lattices for all measure‐ ments. The gold cores were 50 nm in diameter for all samples. Reproduced with permission from reference [104].

These transmission spectra show two peaks resulted from the surface plasmon resonance of the gold core NPs around ~540 nm and the Bragg diffraction of each opaline lattice. Depending on the changes in the silica shell's thickness, the position of the Bragg diffraction peak varied. For example, the Bragg diffraction peak overlapped with the surface plasmon resonance band when the silica shell was 70 nm in thickness, as shown Figure 3. In this case, only one broad absorption peak was observed at ~540 nm.

Figure 4 demonstrates the reflection spectra acquired from the surface array of Au@SiO2 NPs. All samples were wet with the hollow spaces between NPs being completely filled with water when these spectra were measured. The incident light was kept vertical to the (111) planes of these face-center-cubic lattices (Au).

For the reflectance spectra, the Bragg diffraction features were the only peaks detected, and they aligned with the features in the transmission spectra. Additionally, the peaks in Figure

**Figure 4.** Reflectance spectra taken from the arrays produced from Au@SiO2 NPs with different shell thicknesses (t). The incident light was perpendicular to the (111) planes of these face-center-cubic crystalline lattices for all measure‐ ments. The gold cores were 50 nm in diameter for all samples. Reproduced with permission from reference [104].

4 associated with particles that possessed various shell thicknesses, were narrow and wellresolved. This shows that such core@shell nanoparticle arrays provide two pathways for modulating their interaction with light. Liz-Marzan *et al.* also recently illustrated the fabrica‐ tion of 3D crystalline lattices from Au@SiO2 NP colloids; however, these authors failed to provide similar optical characterization [95]. Above all, the spectra displayed in Figures 2, 3, and 4 provide perspective regarding the potential for utilizing the optical properties of Au@SiO2 NP colloids, and their crystalline lattices.

To tune the shell thickness from 20 to 100 nm, the experimental parameters (e.g., coating time and concentration of reactants, catalyst, or other precursors) can be precisely and systemati‐ cally controlled. Li *et al.* demonstrated how the shell thickness for Ag@SiO2 NPs can be tuned by controlling certain parameters, the molar volume ratio of water to surfactant, R (R = [water]/ [surfactant]), and the molar volume ratio of water to TEOS, H (H = [water]/[TEOS]) [98]. The manipulation of these parameters provides control over the availability of water molecules for the hydrolysis of TEOS (tetraethyl othosilicate).

Other researchers have developed the magnetic properties of alternative silica-coated cores (e.g., Fe, Ni, Co, and alloyed metal compounds) to be used in the presence of external magnetic fields for the improvement of bio-imaging, biological labeling, information storage, catalysis, etc. [101, 107-109]. Magnetic NPs can be easily synthesized by using wet chemical processes in aqueous systems, but there is a disadvantage; the difficulty of making a stable dispersion of these NPs for use in aqueous environments or biological systems. To resolve this restriction, a silica-coating on the magnetite core NPs offers excellent dispersion and improved biocom‐ patibility [110]. More recently, magnetic NPs formed from different core and shell magnetic materials have been reported [111]. These unique core@shell structures allow the magnetic properties of the resulting assembly to be more precisely tuned through the choice of magnetic materials and the dimensions of the component parts.

*Other Forms of Core@Shell Nanoparticles.* Various shaped core@shell NPs have proven to be similarly important in research because of their potential applications in the fields of catalysis [112], nanoelectronics [113], information storage [114], and sensors [115]. To synthesize these variously shaped core@shell NPs, researchers have employed a soft or hard core NP template particle to establish the physical shape [116, 117]. The most familiar examples are the use of a firm core NP of a specific shape as the template. Therefore, a soft shell material on a rigid core NP is deposited evenly to present a core@shell NP resembling the shape of the template core. Researchers synthesize core NPs through careful control of the reaction parameters, a process that relies upon controlled crystal growth using surfactants to manipulate the resulting structure [113, 116]. Examples of the types of shapes that can be made (cubic, cuboctahedral, and octahedral) are shown in Figure 5 [113]. These specific examples are Cu@Cu2O NPs and such shaped NPs can be synthesized on a similarly formed core by electro-deposition with a free capping agent.

4 associated with particles that possessed various shell thicknesses, were narrow and wellresolved. This shows that such core@shell nanoparticle arrays provide two pathways for modulating their interaction with light. Liz-Marzan *et al.* also recently illustrated the fabrica‐ tion of 3D crystalline lattices from Au@SiO2 NP colloids; however, these authors failed to provide similar optical characterization [95]. Above all, the spectra displayed in Figures 2, 3, and 4 provide perspective regarding the potential for utilizing the optical properties of

**Figure 4.** Reflectance spectra taken from the arrays produced from Au@SiO2 NPs with different shell thicknesses (t). The incident light was perpendicular to the (111) planes of these face-center-cubic crystalline lattices for all measure‐ ments. The gold cores were 50 nm in diameter for all samples. Reproduced with permission from reference [104].

To tune the shell thickness from 20 to 100 nm, the experimental parameters (e.g., coating time and concentration of reactants, catalyst, or other precursors) can be precisely and systemati‐ cally controlled. Li *et al.* demonstrated how the shell thickness for Ag@SiO2 NPs can be tuned by controlling certain parameters, the molar volume ratio of water to surfactant, R (R = [water]/ [surfactant]), and the molar volume ratio of water to TEOS, H (H = [water]/[TEOS]) [98]. The manipulation of these parameters provides control over the availability of water molecules for

Other researchers have developed the magnetic properties of alternative silica-coated cores (e.g., Fe, Ni, Co, and alloyed metal compounds) to be used in the presence of external magnetic fields for the improvement of bio-imaging, biological labeling, information storage, catalysis, etc. [101, 107-109]. Magnetic NPs can be easily synthesized by using wet chemical processes in aqueous systems, but there is a disadvantage; the difficulty of making a stable dispersion of these NPs for use in aqueous environments or biological systems. To resolve this restriction, a silica-coating on the magnetite core NPs offers excellent dispersion and improved biocom‐ patibility [110]. More recently, magnetic NPs formed from different core and shell magnetic

Au@SiO2 NP colloids, and their crystalline lattices.

108 Nanoparticles Technology

the hydrolysis of TEOS (tetraethyl othosilicate).

**Figure 5.** SEM images of Cu@Cu2O NPs in the form of a) cubic, b) cuboctahedral, and c) octahedral shapes. Repro‐ duced with permission from reference [114].

The fitness of the coating of the shell material can be impacted by the shape of the core NP or the nature of the materials used. This means that the production of a uniform coating might be reduced with shape distortions from spherical or when the material used to form the shell reacts with the core material [118]. In the case of the octahedral gold@platinum NPs, the shell material is incompletely coated on the octahedral core because there is an incomplete reduction of the salt by the core material. Additionally, in some bimetallic "core@shell" structures, the shell is formed by a reduction-transmetalation process which fails to produce a distinguishable shell [119]. On the other hand, with the example given above, a spherical core can be more completely coated with a shell material due to a more perfect reduction of the salt and a more uniform exposure of the metal core surface in solution. These examples highlight the impor‐ tance of the choice of materials used to prepare core@shell composite NPs.

### **3. Smart, multi-responsive Core@Shell nanoparticles**

The development of a stimuli-responsive composite nanoparticle requires an efficient trigger mechanism. One such trigger is the heat generated by the light absorption of metal nanostruc‐ tures. The optical properties of GNSs can be specifically controlled to maximize the absorption/ scattering of light in the wavelength range of 700 to 1000 nm [64, 83]. This is advantageous because light at wavelengths between 800 nm and 1200 nm, a range called the "water window", can penetrate human tissue, enabling its use in biomedical applications [59]. Several studies have pursued the use of this technology in combination with mesoporous silica shells that can carry model drugs. This drug carrying capacity exists in part because of the large pore volume in the etched silica surface, a feature that is tunable to achieve a specific mesopore diameter (2-50 nm). However, mesoporous silica nanoparticles by themselves are not "smart" materials because these NPs cannot release drugs in a precise and controlled manner at a specific location (i.e., they have irreversible pore openings). To overcome this drawback, our research group has explored the growth of a stimuli-responsive hydrogel polymer coating on gold nanoshells that can release a model drug upon the collapse of the hydrogel matrix. These hydrogel polymers are very useful materials in a variety of applications such as drug delivery, chemical separations, and catalysis. But such polymers need an appropriate stimulus to initiate the release of a drug remotely. For many hydrogel applications, heat is used to initiate the collapse of the polymer hydrogel. Since specific frequencies of light can be used to generate heat at the surface of GNSs that are responsive to such light, the combination of a well-tuned nanoshell for specific light absorption and a thermally-responsive hydrogel provide the potential for remotely controlled drug delivery, a smart, multi-responsive core@shell nanoparticle system.

Additionally, core@shell composite particles that respond to magnetic stimuli can also be used to perform useful tasks in controlled drug delivery [77, 78], bio-separation [120], chemical catalysis [121, 122], and electronics [1, 3]. By integrating the application of both their physical and chemical properties, these magnetic core@shell materials can become multifunctional devices that enable a variety of advanced applications that cannot be accomplished by simple magnetic NPs alone. Recent examples of the application of magnetic NPs in research have demonstrated their usefulness because of their capacity to produce heat under an external oscillating magnetic field or to be manipulated remotely, allowing for their use as an antitumor treatment, cell tracking tag, or drug delivery vehicle [123-125]. These core@shell composite particles that respond to both a magnet and other external stimuli, typically consist of a magnetic core encapsulated in a stimuli-responsive hydrogel copolymer layer that responds rapidly to changes in temperature slightly above that of the body. And such a coating heightens the particles' biocompatibility and chemical stability in an aqueous medium [126-138].

Core@shell composite particles that respond to optical, magnetic, and other external stimuli (e.g., temperature, ionic strength, and pH), can be used to perform more useful tasks in research. For this project, we employed a biocompatible mesoporous silica interlayer between the magnetic core and the hydrogel copolymer outer layer to improve the composite particle's loading capacity and payload release effectiveness. The advantages that such porous struc‐ tures provide is their high surface area (ca. 1000 m2 /g), large pore volume (ca. 1 cm3 /g), tunable mesopore diameter (2-50 nm) and biocompatibility [139, 140]. Thus, impregnation of these mesoporous silica and poly(N-isopropylacrylamide-co-acrylic acid; NIPAM-co-AAc)-coated magnetic NPs (or gold GNSs as the core) with drugs produces a nanoscale drug-delivery system that can be specifically targeted and magnetically (or phothermally) activated. We call them "smart" core@shell NPs.

### **4. Synthesis of smart multi-responsive Core@Shell nanoparticles**

To accomplish the research goals described above, hydrogel-based core@shell composite NPs were fabricated by encapsulating a mesoporous silica-coated GNS (or Fe3O4 NP) as the core with a PNIPAM-*co*-AAc copolymer coating [141]. The oleylamine-functionalized mesoporous silica-coated GNS (or Fe3O4 NP) was used as a nano-template for the shell layer growth of a hydrogel copolymer. Ammonium persulfate (APS) was used as a polymerization initiator to produce a hydrogel-encapsulated composite NP. The amount of NIPAM monomer was optimized for the hydrogel-encapsulated mesoporous silica-coated composite NPs [142]. The shell layer thickness was increased with an increase in polymerization time until no further increase in the shell layer thickness was clearly observed [143]. Hydrogel-encapsulated mesoporous silica-coated composite NPs exhibited systematic changes in particle size corresponding to the variation of temperature, which originates from hydrogen-bonding interactions between PNIPAM amide groups and water, as well as electrostatic forces attrib‐ uted to the ionization of carboxylic groups in the acrylic acid.

### **5. Long term research objectives**

**3. Smart, multi-responsive Core@Shell nanoparticles**

110 Nanoparticles Technology

The development of a stimuli-responsive composite nanoparticle requires an efficient trigger mechanism. One such trigger is the heat generated by the light absorption of metal nanostruc‐ tures. The optical properties of GNSs can be specifically controlled to maximize the absorption/ scattering of light in the wavelength range of 700 to 1000 nm [64, 83]. This is advantageous because light at wavelengths between 800 nm and 1200 nm, a range called the "water window", can penetrate human tissue, enabling its use in biomedical applications [59]. Several studies have pursued the use of this technology in combination with mesoporous silica shells that can carry model drugs. This drug carrying capacity exists in part because of the large pore volume in the etched silica surface, a feature that is tunable to achieve a specific mesopore diameter (2-50 nm). However, mesoporous silica nanoparticles by themselves are not "smart" materials because these NPs cannot release drugs in a precise and controlled manner at a specific location (i.e., they have irreversible pore openings). To overcome this drawback, our research group has explored the growth of a stimuli-responsive hydrogel polymer coating on gold nanoshells that can release a model drug upon the collapse of the hydrogel matrix. These hydrogel polymers are very useful materials in a variety of applications such as drug delivery, chemical separations, and catalysis. But such polymers need an appropriate stimulus to initiate the release of a drug remotely. For many hydrogel applications, heat is used to initiate the collapse of the polymer hydrogel. Since specific frequencies of light can be used to generate heat at the surface of GNSs that are responsive to such light, the combination of a well-tuned nanoshell for specific light absorption and a thermally-responsive hydrogel provide the potential for remotely controlled drug delivery, a smart, multi-responsive core@shell nanoparticle system.

Additionally, core@shell composite particles that respond to magnetic stimuli can also be used to perform useful tasks in controlled drug delivery [77, 78], bio-separation [120], chemical catalysis [121, 122], and electronics [1, 3]. By integrating the application of both their physical and chemical properties, these magnetic core@shell materials can become multifunctional devices that enable a variety of advanced applications that cannot be accomplished by simple magnetic NPs alone. Recent examples of the application of magnetic NPs in research have demonstrated their usefulness because of their capacity to produce heat under an external oscillating magnetic field or to be manipulated remotely, allowing for their use as an antitumor treatment, cell tracking tag, or drug delivery vehicle [123-125]. These core@shell composite particles that respond to both a magnet and other external stimuli, typically consist of a magnetic core encapsulated in a stimuli-responsive hydrogel copolymer layer that responds rapidly to changes in temperature slightly above that of the body. And such a coating heightens the particles' biocompatibility and chemical stability in an aqueous medium

Core@shell composite particles that respond to optical, magnetic, and other external stimuli (e.g., temperature, ionic strength, and pH), can be used to perform more useful tasks in research. For this project, we employed a biocompatible mesoporous silica interlayer between the magnetic core and the hydrogel copolymer outer layer to improve the composite particle's loading capacity and payload release effectiveness. The advantages that such porous struc‐

[126-138].

Lee and co-workers recently reported the initial methodology to precisely control drug delivery by employing gold NPs and GNSs coated with a pH- and temperature-responsive hydrogel originating from the co-polymerization of NIPAM and acrylic acid [142-144]. These nontoxic composite NPs were designed to be loaded with drug molecules, providing the ability for the NP cores to be photothermally activated, initiating collapse of the hydrogel coating and releasing the drug molecules, as illustrated in Figure 6.

Furthermore, we have been working to employ a *m*esoporous silica interlayer between a goldcoated silica core and a hydrogel outer layer to prevent unwanted structural changes to the gold shell during photomodulation and to assist in the carrying of hydrophobic or hydrophilic drugs to targeted sites [141]. The advantages that such porous structures provide is their high surface area (ca. 1000 m2 /g), large pore volume (ca. 1 cm3 /g), tunable mesopore diameter (2-50 nm) and biocompatibility [139, 140]. By using smart hydrogel technology as an outer layer and a mesoporous silica coating as an interlayer, GNSs that are activated by tissue-transparent near-IR light can be more effectively used for advanced medicinal applications. With our initial investigation, once these nontoxic composite NPs were loaded with methylene blue (MB; a

**Figure 6.** Illustration of drug release by increasing the environmental temperature or by generating heat through an SPR response in these optically active composite particles to a near IR laser.

dye used as a model drug), and the NPs thermally activated, the hydrogel coating collapsed and released the test molecules. The potential for such smart, multi-responsive core@shell nanoparticles for use as drug delivery carriers, contrast agents, and therapeutic entities will clearly encourage the application of this new technology in future research projects.

### **Acknowledgements**

The Robert A. Welch Foundation (Grant No. E-1320) and the Texas Center for Superconduc‐ tivity at the University of Houston provided generous support for this research.

### **Author details**

Bo Sang Kim and T. Randall Lee\*

\*Address all correspondence to: trlee@uh.edu

Department of Chemistry and the Texas Center for Superconductivity, University of Hous‐ ton, Houston, Texas, USA

### **References**

[1] Teranishi T, Miyake M. Size control of palladium nanoparticles and their crystal structures. Chem Mater. 1998;10:594-600. DOI: 10.1021/cm9705808.


dye used as a model drug), and the NPs thermally activated, the hydrogel coating collapsed and released the test molecules. The potential for such smart, multi-responsive core@shell nanoparticles for use as drug delivery carriers, contrast agents, and therapeutic entities will

**Figure 6.** Illustration of drug release by increasing the environmental temperature or by generating heat through an

The Robert A. Welch Foundation (Grant No. E-1320) and the Texas Center for Superconduc‐

Department of Chemistry and the Texas Center for Superconductivity, University of Hous‐

[1] Teranishi T, Miyake M. Size control of palladium nanoparticles and their crystal

structures. Chem Mater. 1998;10:594-600. DOI: 10.1021/cm9705808.

clearly encourage the application of this new technology in future research projects.

SPR response in these optically active composite particles to a near IR laser.

tivity at the University of Houston provided generous support for this research.

**Acknowledgements**

112 Nanoparticles Technology

**Author details**

Bo Sang Kim and T. Randall Lee\*

ton, Houston, Texas, USA

**References**

\*Address all correspondence to: trlee@uh.edu


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*Edited by Mahmood Aliofkhazraei*

During the past years, scientists have achieved significant success in the nanoscience and nanotechnology. Nanotechnology is a field of applied sciences which is focused on design, production, detection, and employing the nano-size materials, pieces, and equipment. Advances in nanotechnology lead to improvement of tools and equipment as well as their application in everyday life. In the chemistry this size involves the range of colloids, micelles, polymer molecules, and structures such as very large molecules or dense accumulation of the molecules. In physics of electrical engineering, the nanoscience is strongly related to quantum behaviour or electrons behaviour in structures with nano sizes. Also, in biology and biochemistry, there are interesting cellular components and molecular structures. This book collects new developments about nanoparticles.

Photo by pavel77731 / DollarPhotoClub

Nanoparticles Technology

Nanoparticles Technology

*Edited by Mahmood Aliofkhazraei*