**Synthesis and Applications**

[17] Feng Z, Zhao J, Huggins FE, Huffman GP. Agglomeration and phase transition of a

[18] Zhao J, Huggins FE, Feng Z, Huffman GP. Ferrihydrite; surface structure and its effects

[19] Han KN. Fundamentals of Aqueous Metallurgy. Littleton: Society for Mining, Metal-

[20] Yamashita M, Maeda A, Uchida H, Kamimura T, Miyuki H. Crystalline rust compositions and weathering properties of steels exposed in nation-wide atmospheres for 17

nanophase iron oxide catalyst. J Catal. 1993;143:510–519.

lurgy, and Exploration; 2002. 197 p.

182 Magnetic Spinels- Synthesis, Properties and Applications

years. J Jpn Inst Metals. 2001;65:967–971.

on phase transformation. Clay Clay Miner. 1994;42:737–746.

Provisional chapter

## **Nanostructured Spinel Ferrites: Synthesis, Functionalization, Nanomagnetism and Environmental Applications** Nanostructured Spinel Ferrites: Synthesis, Functionalization, Nanomagnetism and Environmental Applications

Oscar F. Odio and Edilso Reguera Oscar F. Odio and Edilso Reguera

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

http://dx.doi.org/10.5772/67513

#### Abstract

Nanostructured spinel ferrites have gained a great deal of attention. It comes from the possibility of tuning their magnetic properties by careful manipulation of the synthetic conditions. At the same time, since the nanoparticle (NP) surface is reactive toward many chemical groups, it provides great versatility for further functionalization of the nanosystems. Such characteristics make ferrite nanoparticles excellent candidates for environmental applications. First, the chapter deals with the basics of the synthetic methodologies, functionalization strategies and magnetic properties of nanoparticles, with emphasis on how surface manipulation is reflected in the properties of the materials. Next, we review some of the applications of ferrites as magnetic sorbents for several hazardous substances in aqueous medium and try to systematize the adsorption mechanism as a function of the coating material. Finally, a short summary concerning the main uses of ferrites as magnetic catalysts in oxidation technologies is included.

Keywords: spinel ferrites, superparamagnetism, surface complexes, heavy metals, dyes

## 1. Introduction

Magnetic nanoparticles (NPs) have been the focus of intense studies, both at the fundamental and at the technological level. Among many promising materials, nanostructured spinel ferrites occupy a special place. These iron oxide-based materials are easy and cheap to synthesize, are stable under a wide range of conditions and some family members present low toxicity for living organisms. Besides, due to their high reactivity toward several organic groups, ferrite surface offers a great versatility for ligand functionalization, which in many cases defines the ultimate application. In addition, one of the most prominent properties of spinel ferrite NPs is

© The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons © 2017 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

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

the onset of superparamagnetism. This phenomenon is a crucial feature for several biomedical applications [1], catalytic processes [2, 3] and environmental remediation strategies [1, 4–7]. Currently, there are available in the literature several extensive reviews covering these issues in detail [8, 9]. In this chapter, we focus primarily on adsorption and oxidation technologies for water decontamination using nanostructured spinel ferrites where particle functionalization plays a major role. In particular, we focus on basic topics concerning spinel ferrite NPs with an emphasis on the surface manipulation by chemical methods and how it is reflected in the properties and performances of the ultimate nanomaterial. Also, attention is paid to the machinery that governs the adsorption process in order to try to systematize the available data. Every step in this direction is aimed to improve and design newer and better solutions for the great challenge of water remediation.

## 2. Structural and magnetic properties

Spinel ferrites are mixed valence oxides where oxygen anions form a close-packed cubic array, while metallic cations occupy randomly one-eighth of the tetrahedral (A) sites and one-half of the octahedral (B) interstices. This family is classified into the Fd3m space group with the general formula ½Mð1�iÞFei� <sup>A</sup>½MiFeð2�<sup>i</sup>Þ� <sup>B</sup>O4, where M (Fe, Ni, Mg, Mn, Zn, Co, etc.) is a divalent cation that shares with Fe(III) cations the occupancies of A and B sites, while i is defined as the inversion parameter. There are three possibilities [10]: (i) i = 0 yields a normal spinel such as ZnFe2O4; (ii) i = 1 yields an inverse spinel such as Fe3O4 and CoFe2O4 and (iii) for 0 < i < 1 cations are distributed on both sites yielding a partially inverted spinel such as MnFe2O4 in which i = 0.2. In addition, it is possible to synthesize mixed ferrites in the sense that different divalent cations could coexist in the same compound [11, 12]. All these possibilities open an amazing range for tailoring different properties [11, 13].

In the spinel structure (left panel of Figure 1), magnetic moments of sublattice A are coupled with magnetic moments of sublattice B in an antiferromagnetic fashion by superexchange interactions mediated by the oxygen anions [14]. Since spins in both lattices are generally uncompensated, the resulting net magnetic moment causes the material to display ferrimagnetic behavior [10]. As can be inferred, either the type of divalent ion or the average ion distribution plays a critical role in the magnetic properties of the material. This effect can be illustrated by varying the Zn content in the mixed ZnxFe3�<sup>x</sup>O4 ferrite [11] in which Zn2+ presents a zero magnetic moment (μ = 0) and a high tendency for tetrahedral A sites. In the interval 0 < x < 0.4, the saturation magnetization (MS) increases as x increases, since the antiferromagnetic A-B interaction is progressively weakened, thus enhancing the net magnetic moment in sublattice B (see the right panel of Figure 1). However, for values of x near 1, A-B interactions are no longer dominant, but the magnetism now depends on the very weak B-B interactions, thus leading to a marked decrease of MS, which becomes zero at x = 1.

#### 2.1. Nanomagnetism

Magnetic NPs differ from bulk magnetic materials mainly due to the finite size and surface effects. The reduction of size leads to a single magnetic domain at a particular size and the onset

Nanostructured Spinel Ferrites: Synthesis, Functionalization, Nanomagnetism and Environmental Applications http://dx.doi.org/10.5772/67513 187

Figure 1. Left: Representation of a partial spinel ferrite unit cell and the ferrimagnetic order of the structure (Reproduced from Ref. [10] with permission of the American Society of Chemistry). Right: Schematic view of the spin organization in sublattice A and B as a function of the content of Zn in ZnxFe3�<sup>x</sup>O4 (Adapted from Ref. [11] with permission of Wiley).

of superparamagnetism, while surface effects result in symmetry breaking of the crystal structure, which could also alter the magnetic properties. These new features are treated briefly below.

#### 2.1.1. Single-domain limit

the onset of superparamagnetism. This phenomenon is a crucial feature for several biomedical applications [1], catalytic processes [2, 3] and environmental remediation strategies [1, 4–7]. Currently, there are available in the literature several extensive reviews covering these issues in detail [8, 9]. In this chapter, we focus primarily on adsorption and oxidation technologies for water decontamination using nanostructured spinel ferrites where particle functionalization plays a major role. In particular, we focus on basic topics concerning spinel ferrite NPs with an emphasis on the surface manipulation by chemical methods and how it is reflected in the properties and performances of the ultimate nanomaterial. Also, attention is paid to the machinery that governs the adsorption process in order to try to systematize the available data. Every step in this direction is aimed to improve and design newer and better solutions

Spinel ferrites are mixed valence oxides where oxygen anions form a close-packed cubic array, while metallic cations occupy randomly one-eighth of the tetrahedral (A) sites and one-half of the octahedral (B) interstices. This family is classified into the Fd3m space group with the

cation that shares with Fe(III) cations the occupancies of A and B sites, while i is defined as the inversion parameter. There are three possibilities [10]: (i) i = 0 yields a normal spinel such as ZnFe2O4; (ii) i = 1 yields an inverse spinel such as Fe3O4 and CoFe2O4 and (iii) for 0 < i < 1 cations are distributed on both sites yielding a partially inverted spinel such as MnFe2O4 in which i = 0.2. In addition, it is possible to synthesize mixed ferrites in the sense that different divalent cations could coexist in the same compound [11, 12]. All these possibilities open an

In the spinel structure (left panel of Figure 1), magnetic moments of sublattice A are coupled with magnetic moments of sublattice B in an antiferromagnetic fashion by superexchange interactions mediated by the oxygen anions [14]. Since spins in both lattices are generally uncompensated, the resulting net magnetic moment causes the material to display ferrimagnetic behavior [10]. As can be inferred, either the type of divalent ion or the average ion distribution plays a critical role in the magnetic properties of the material. This effect can be illustrated by varying the Zn content in the mixed ZnxFe3�<sup>x</sup>O4 ferrite [11] in which Zn2+ presents a zero magnetic moment (μ = 0) and a high tendency for tetrahedral A sites. In the interval 0 < x < 0.4, the saturation magnetization (MS) increases as x increases, since the antiferromagnetic A-B interaction is progressively weakened, thus enhancing the net magnetic moment in sublattice B (see the right panel of Figure 1). However, for values of x near 1, A-B interactions are no longer dominant, but the magnetism now depends on the very weak B-B

interactions, thus leading to a marked decrease of MS, which becomes zero at x = 1.

Magnetic NPs differ from bulk magnetic materials mainly due to the finite size and surface effects. The reduction of size leads to a single magnetic domain at a particular size and the onset

<sup>B</sup>O4, where M (Fe, Ni, Mg, Mn, Zn, Co, etc.) is a divalent

for the great challenge of water remediation.

186 Magnetic Spinels- Synthesis, Properties and Applications

2. Structural and magnetic properties

<sup>A</sup>½MiFeð2�<sup>i</sup>Þ�

amazing range for tailoring different properties [11, 13].

general formula ½Mð1�iÞFei�

2.1. Nanomagnetism

Large magnetic particles usually have a multidomain structure, each domain separated from its neighbors by domain walls. As the particle diameter D or volume V is decreased, domain wall creation is no longer energy favorable for a specific size, leading to single-domain NPs with all atomic spins aligned in the same direction. This critical particle diameter is characteristic for each material and is of the order of tens of nanometers (128 nm for Fe3O4) [9]. Since the spins are parallel within the NPs, magnetic reversal does not depend on wall displacement but is only possible by the coherent rotation of spins, which depends entirely on the effective anisotropy (Keff). Given that coercitivity (HC) is proportional to Keff, it is higher for nanomaterials with respect to their bulk counterparts. For single-domain particles with uniaxial anisotropy, the anisotropy energy is defined as:

$$E(\theta) = K\_{\rm eff} V \sin^2 \theta \tag{1}$$

Here, θ is the angle between the net magnetization and the easy axis of magnetization. For spherical particles, Keff can be expressed as:

$$K\_{eff} = K\_V + \frac{6}{D} K\_S \tag{2}$$

Here, KV and KS are the volume and surface anisotropies, respectively. As can be seen, for ultrasmall NPs, the surface term may dominate the total anisotropy of the material.

#### 2.1.2. Superparamagnetism

The product Kef f V is the energy barrier for the coherent rotation of all atomic spins between the two equivalent easy axes of magnetization. As D is decreased, the thermal energy kBT eventually overcomes the energy barrier for a particular size, thus leading to thermal equilibrium of the total magnetic moment of the system. In this state, the resulting HC is zero and the system behaves like a paramagnet but with a huge magnetic moment. For an assembly of noninteracting single-domain NPs, each magnetic moment fluctuates with a relaxation time τ given by the Arrhenius-Néel law [15]:

$$
\pi = \pi\_0 \mathcal{e}^{(\mathbf{K}\_{\rm eff} V/\mathbf{k}\_{\rm B} \mathbf{T})} \tag{3}
$$

Here, τ<sup>0</sup> is a characteristic time of the system and the actual magnetic state at a given T depends on the measuring time τm. If τ < τm, the systems are in the superparamagnetic state, and for τ > τm, the spins appear in a blocked state. The temperature for which τ ¼ τ<sup>m</sup> is the socalled blocking temperature TB and is given by:

$$T\_B = \frac{K\_{eff}V}{k\_B} \ln \quad (\tau\_0/\tau\_m) \tag{4}$$

One of the main advantages of ferrites is the possibility for tuning the magnetic properties by varying simply either the divalent cation or the arrangement of the metals into the spinel structure. For instance, in a series of nanoparticle ferrites MFe2O4 (M = Mn2+, Fe2+, Co2+, Ni2+, Zn2+), Mohapatra et al. [16] could varyMS and TB by taking advantage of the differing magnetic moments and spin-orbit coupling strengths of M2+ cation. Similar conclusions can be extracted from other reports [17]. Likewise, for ultra-small CoFe2O4 NPs, the progressive variation in the inversion index (from a total to a partially inverted spinel) as D decreases is reflected in a decrease in the magnetocrystalline anisotropy and hence, in the HC of the material [18]. In other works, Taharet al. [13] found that as Zn2+ substitutes Co2+ cations in 5-nm-mixed ZnxCo1�xFe2O4 ferrite NPs both HC and TB decrease, in accordance with progressive reduction of magnetocrystalline anisotropy, which is higher for octahedral coordinated Co2+ cations.

#### 2.1.3. Surface effects

Progressive decrease in NPs' size makes the number of atoms on the surface comparable with the number of atoms in the bulk. For magnetic NPs, this trend lowers MS and increases Keff as D decreases [19, 20]. The decrease in MS is associated with the presence of a magnetic dead layer, occurrence of spin canting or spins glass-like behavior at the surface level. The increase in Keff is associated with the onset of surface anisotropy following Eq. (2).

The effect of surface coating in order to tune the magnetic properties of NPs is another area of active investigations. The adsorption of organic ligands could alter the particle size distribution, the interparticle interactions and the spin canting at the surface [21]. The overall effect seems to be the result of a complex interplay between the coordination mode, the capping density and the surface disorder in the synthesized sample [22]. In a careful study of adsorption of stilbene carboxylates and phosphonates as capping agents of 39 nm Fe3O4 NPs, Daou et al. [23] reported that carboxyl ligands tend to enhance spin canting at the surface of the oxide, leading to MS reduction; on the other hand, phosphonate ligands seem to mimic the iron coordination in the bulk, hence, MS of the NPs is unaltered as for uncovered Fe3O4. However, a direct correlation between magnetic measurements and the nature of the coordination bonding at the organic-inorganic interface is still needed. Regarding the important case of carboxylates, Odio et al. [24] found that spin disorder is larger in chelating than in bridging complexes; it suggests that occurrence of this last geometry makes possible the partial reconstruction of the crystal field of iron ions in the bulk phase. In another report, Aslibeiki et al. [25] noted that tetraethylene glycol ligands attached to Fe3O4 surface also contribute to decrease the surface spin disorder. Similarly, Jia et al. [26] noted that in Co- and Ni-mixed ferrite with SO4 2 attached to the surface, MS values increase with the content of superficial anions. In contrast, adsorption of carbonyl groups belonging to poly(vinyl pyrrolidone) (PVP) capping chains of Mn ferrites seems to decrease MS values [27].

In an interesting paper, Vestal and Zhang [28] performed a systematic study of the correlation between the nature of the capping ligand (substituted benzenes) and the magnetic properties of MnFe2O4 ferrites with different particles' diameter. They found that for very small NPs (4 nm), HC is reduced by ligand interaction with respect to the uncovered ferrite, while MS is enhanced. Such behavior is consistent with the fact that metal-ligand interactions at the surface reduce KS, leading to a reduction of HC, and at the same time, they induce spin order at the surface. Furthermore, the authors showed that the higher the crystal field splitting energy of the surface complex, the smaller the magnitude of the spin-orbit interaction parameter and hence, the smaller the K<sup>S</sup> value. The fact that such trends are less pronounced for larger particles (12 nm and 25 nm) reveals the importance of surface effects in HC and MS of singledomain NPs. This point has been verified in other reports [16, 29]. Besides, the contribution of KS to Keff is revealed when comparing the HC variation of MnFe2O4 and CoFe2O4 particles with the same diameter and organic ligands: in the case of Co ferrite, the larger magnetocrystalline anisotropy masks K<sup>S</sup> variation, leading to nearly unchanged values of HC [30].

#### 2.1.4. Magnetic interparticle interactions

Kef f ¼ KV þ

small NPs, the surface term may dominate the total anisotropy of the material.

2.1.2. Superparamagnetism

given by the Arrhenius-Néel law [15]:

188 Magnetic Spinels- Synthesis, Properties and Applications

called blocking temperature TB and is given by:

which is higher for octahedral coordinated Co2+ cations.

2.1.3. Surface effects

Here, KV and KS are the volume and surface anisotropies, respectively. As can be seen, for ultra-

The product Kef f V is the energy barrier for the coherent rotation of all atomic spins between the two equivalent easy axes of magnetization. As D is decreased, the thermal energy kBT eventually overcomes the energy barrier for a particular size, thus leading to thermal equilibrium of the total magnetic moment of the system. In this state, the resulting HC is zero and the system behaves like a paramagnet but with a huge magnetic moment. For an assembly of noninteracting single-domain NPs, each magnetic moment fluctuates with a relaxation time τ

τ ¼ τ0e

TB <sup>¼</sup> Kef f <sup>V</sup> kB

Here, τ<sup>0</sup> is a characteristic time of the system and the actual magnetic state at a given T depends on the measuring time τm. If τ < τm, the systems are in the superparamagnetic state, and for τ > τm, the spins appear in a blocked state. The temperature for which τ ¼ τ<sup>m</sup> is the so-

One of the main advantages of ferrites is the possibility for tuning the magnetic properties by varying simply either the divalent cation or the arrangement of the metals into the spinel structure. For instance, in a series of nanoparticle ferrites MFe2O4 (M = Mn2+, Fe2+, Co2+, Ni2+, Zn2+), Mohapatra et al. [16] could varyMS and TB by taking advantage of the differing magnetic moments and spin-orbit coupling strengths of M2+ cation. Similar conclusions can be extracted from other reports [17]. Likewise, for ultra-small CoFe2O4 NPs, the progressive variation in the inversion index (from a total to a partially inverted spinel) as D decreases is reflected in a decrease in the magnetocrystalline anisotropy and hence, in the HC of the material [18]. In other works, Taharet al. [13] found that as Zn2+ substitutes Co2+ cations in 5-nm-mixed ZnxCo1�xFe2O4 ferrite NPs both HC and TB decrease, in accordance with progressive reduction of magnetocrystalline anisotropy,

Progressive decrease in NPs' size makes the number of atoms on the surface comparable with the number of atoms in the bulk. For magnetic NPs, this trend lowers MS and increases Keff as D decreases [19, 20]. The decrease in MS is associated with the presence of a magnetic dead layer, occurrence of spin canting or spins glass-like behavior at the surface level. The increase

in Keff is associated with the onset of surface anisotropy following Eq. (2).

6

<sup>D</sup> KS (2)

<sup>ð</sup>Kef f <sup>V</sup>=kBT<sup>Þ</sup> (3)

ln ðτ0=τmÞ (4)

The presence of magnetic interactions between particles has a great influence on superparamagnetism [31–33]. This effect alters the energy barrier for coherent rotation, which is no longer governed by only anisotropic contributions. The system becomes very complex and results in the difficulty to separate the contributions of different factors [34]. For ferrite NPs, the ordinary kinds of magnetic interactions are dipolar-dipolar and direct exchange interactions between spins at the interface of particles in close contact [35]. The first contribution is almost ubiquitous in any system, given its anisotropic and long-range nature, which could favor either ferro or antiferromagnetic alignment of the spins. The minimization of such an effect can only be achieved in samples where individual particles are well separated from each other, either by steric or by coulombic repulsions [25, 36]. The dipolar magnetic field generated by a single spherical particle is proportional to its volume; hence, the effect is more pronounced for large particles. In that case, the magnetic energy between two spheres decays with d�<sup>3</sup> , but smaller systems in the superparamagnetic state (where moments fluctuate) partially destroy the order induced by dipolar interactions, and hence, the energy associated decays with d�<sup>6</sup> . In the limit of strong interactions, particles do not relax according to their own energy barrier, but the magnetic evolution depends on the energy of the whole ensemble [22, 34, 36, 37].

In the presence of interparticle interactions (IPI), Eq. (4) for TB is modified to [38–40]:

$$T\_B = To + \frac{K\_{eff}V}{k\_B} \ln\left(\tau\_0/\tau\_m\right) \tag{5}$$

Here, To < TB is a measure of the strength of the IPI in the system. To determine To, TB needs to be measured at several different measuring frequencies fm = 1/τm. The presence of IPI is also determined by the quantity Φ ¼ ΔTB=½TBΔlog10f <sup>m</sup>� with ΔTB ¼ TB(2)�TB(1) being the difference in TB determined at two sufficiently different measuring frequencies fm(2) > fm(1) [41]. The magnitude of Φ < 0.13 signifies the presence of IPI with its strength increasing with decreasing magnitude of Φ: In a recent paper [42], Φ and To have been shown to be related by the equation:

$$\Phi = \Phi\_o \left[ 1 - \left( T\_o / T\_B(1) \right) \right] \tag{6}$$

In Eq. (6), Φ<sup>0</sup> ¼ 2:3026=fln½f <sup>0</sup>=f <sup>m</sup>ð2Þ�g. As noted above, measurements of TB at several different frequencies are essential in order to determine Keff since fo and To must also be determined simultaneously [40].

#### 3. Synthesis

Synthesis of spinel ferrite NPs is a challenging task owing to their colloidal nature. A good methodology must yield well-dispersed particles with uniform size and good crystallinity; besides, it is desirable that the synthetic setup allows for the tuning of structure and properties of the materials by simple modification of the conditions. Other important features entail the use of nontoxic reagents, low-temperature processes and the requirement of simple scalable operations. The procedures for the synthesis of ferrite NPs are given in several reviews [2, 8]. Here, we outline some of the most common examples.

#### 3.1. Co-precipitation

Co-precipitation is straightforward and efficient and can be extended for a wide variety of simple and mixed ferrites [12, 43, 44]. This method, developed by Massart, consists of the joint precipitation of an aqueous solution containing inorganic salts in the proper stoichiometry by increasing the solution pH. Ageing of the resulting particles can be assessed at room or higher temperatures. By changing experimental conditions (e.g., concentration of metal precursors, pH of the final solution, anion of initial salts, reaction time, reaction temperature and ionic strength), it is possible to obtain a wide variety of particle sizes and shapes [45, 46]. The main drawback relies on the difficulties for a proper separation of nucleation and growth stages, which leads to relative broad size distribution [45]. Besides, in some instances, the resulting powder is subjected to thermal annealing to enhance crystallinity [12, 43].

Nucleation and growth of NPs can be affected by the addition of surfactant molecules like sodium dodecyl sulfate [47], poly(acrylic) acid (PAA) chains [48] and hexadecyl trimethylammonium bromide (CTAB) [49]. Variations in the surfactant content give rise to different particle sizes and morphologies of the as-synthesized material. Other employed additives aiming to decrease the particle size dispersion are polymeric matrices like cellulose [50] and chitosan [51, 52]. A similar approach entailed the use of graphene oxide (GO) during the coprecipitation step. After the formation of the ferrite/GO composite, GO is reduced to yield porous nanocomposites containing superparamagnetic ferrite NPs and reduced GO (rGO) is used as a functional material. The resulting material possesses high surface area since rGO avoids ferrite particle agglomeration. Ni, Co and Mn ferrites/rGO nanocomposites have been synthesized with this strategy [53–55]. Alternatively, some investigations have reported the use of organic amines, which can act as precipitating and stabilizing agents [17, 56, 57]. Alkanolamines limit and control the particle growth by forming surface complexes with M2+ cations resulting in a marked reduction in Co ferrite size as compared when using NaOH [56].

#### 3.2. Thermal decomposition

effect can only be achieved in samples where individual particles are well separated from each other, either by steric or by coulombic repulsions [25, 36]. The dipolar magnetic field generated by a single spherical particle is proportional to its volume; hence, the effect is more pronounced for large particles. In that case, the magnetic energy between two spheres decays with d�<sup>3</sup>

smaller systems in the superparamagnetic state (where moments fluctuate) partially destroy the order induced by dipolar interactions, and hence, the energy associated decays with d�<sup>6</sup>

the limit of strong interactions, particles do not relax according to their own energy barrier, but

Kef f V kB

Here, To < TB is a measure of the strength of the IPI in the system. To determine To, TB needs to be measured at several different measuring frequencies fm = 1/τm. The presence of IPI is also determined by the quantity Φ ¼ ΔTB=½TBΔlog10f <sup>m</sup>� with ΔTB ¼ TB(2)�TB(1) being the difference in TB determined at two sufficiently different measuring frequencies fm(2) > fm(1) [41]. The magnitude of Φ < 0.13 signifies the presence of IPI with its strength increasing with decreasing magnitude of Φ: In a recent paper [42], Φ and To have been shown to be related by the equation:

In Eq. (6), Φ<sup>0</sup> ¼ 2:3026=fln½f <sup>0</sup>=f <sup>m</sup>ð2Þ�g. As noted above, measurements of TB at several different frequencies are essential in order to determine Keff since fo and To must also be determined

Synthesis of spinel ferrite NPs is a challenging task owing to their colloidal nature. A good methodology must yield well-dispersed particles with uniform size and good crystallinity; besides, it is desirable that the synthetic setup allows for the tuning of structure and properties of the materials by simple modification of the conditions. Other important features entail the use of nontoxic reagents, low-temperature processes and the requirement of simple scalable operations. The procedures for the synthesis of ferrite NPs are given in several reviews [2, 8].

Co-precipitation is straightforward and efficient and can be extended for a wide variety of simple and mixed ferrites [12, 43, 44]. This method, developed by Massart, consists of the joint precipitation of an aqueous solution containing inorganic salts in the proper stoichiometry by increasing the solution pH. Ageing of the resulting particles can be assessed at room or higher temperatures. By changing experimental conditions (e.g., concentration of metal precursors, pH of the final solution, anion of initial salts, reaction time, reaction temperature and ionic

To=TBð1Þ

�i

the magnetic evolution depends on the energy of the whole ensemble [22, 34, 36, 37]. In the presence of interparticle interactions (IPI), Eq. (4) for TB is modified to [38–40]:

TB ¼ To þ

Φ ¼ Φ<sup>o</sup> h 1� �

Here, we outline some of the most common examples.

simultaneously [40].

190 Magnetic Spinels- Synthesis, Properties and Applications

3.1. Co-precipitation

3. Synthesis

, but

. In

(6)

ln ðτ0=τmÞ (5)

This method utilizes thermal decomposition of organic metal complexes in a high boiling point solvent and in the presence of a surfactant. This approach yields monodisperse highly crystalline NPs and allows for the fine-tuning of NP size and morphology by controlling several parameters like the solvent nature, kind and concentration of surfactant, aging temperature and reaction time. The typical setup with oleic acid (OA)/oleylamine (OAm) as surfactants can be used to obtain (ZnxM1<sup>x</sup>)Fe2O4 (M = Fe2+, Mn2+) mixed ferrites with different Zn contents as the doping cation [11]. OAm is believed to assist OA deprotonation, which promotes the formation of iron carboxylates at the NP surface [58]. In a systematic study, Mohapatra et al. [16] have reported the synthesis of MFe2O4 (M = Mn2+, Co2+, Fe2+, Ni2+, Zn2+) NPs using chloride salts as precursors and OAm acting as a solvent, reducing agent and stabilizing surface capping agent. By decreasing the amine content, it was possible to obtain uniform NPs with D values between 2 and 9 nm. OAm chains seem to control the growth process: high concentration enables an extended coverage of the initial nuclei, which hampers a fast growth and leads to uniform small NPs.

In an important report on the synthesis of Fe3O4 and other ferrite NPs, Hyeon et al. [59] used iron(III) oleate as an organometallic precursor with the purpose to avoid environmental harmful reagents like Fe(CO)5 [60] and Fe[acac]3 [61]. Also, there is no need for external reducing agents and extra surfactant stabilizers [36, 61–63]. The authors obtained high yields of wellcrystallized monodispersed NPs with D values ranging from 9–22 nm by varying the solvent boiling point. Moriya et al. [64] introduced an interesting approach in which a pre-synthesized trimetallic complex containing two Fe3+ cations and one divalent cation (Fe, Co, Mn) is decomposed in dibenzyl ether in the presence of benzylic acid (BA). The resulting nanocrystals showed uniform size with variable shapes (from truncated octahedrons to cubes) that can be tuned as BA concentration is increased; the nature of the weak intermolecular interactions between adsorbed BA molecules seems to play a key role to bring about the final morphology.

One drawback of the thermal decomposition method is that the as-synthesized ferrite NPs do not disperse in water due to the hydrophobic surfactant adsorbed onto the surface, which leads to further phase transfer steps. To obtain directly water-soluble NPs, the groups of Li [29] and Verma [58] introduced a variation in which 4–5 nm Fe and Co ferrites are obtained through the decomposition of metal acetylacetonates in the presence of pyrrolidones that act either as solvents or as hydrophilic stabilizing agents.

### 3.3. Polyol method

This is a variation of the thermal decomposition in which a given polyol acts as a high-boiling solvent, reducing and stabilizer agent. Metal precursors are generally organic complexes like acetylacetonates and other carboxylate complexes [65]. Given that the reaction mixture is refluxed at the boiling point of the polyol, changing either the kind or concentration of the polyol leads to different particle D values [66–68], generally between 4 and 15 nm. The obtained NPs have a narrow size distribution and high crystallinity, although particle agglomeration could occur. There is question about the nature of the molecules adsorbed at the particle surface; some authors have stressed that polyol anchoring to the surface occurs through R-O interactions [69, 70], but others have claimed that at high temperatures, hydroxyl groups are oxidized to carboxylic acids, which are further adsorbed onto the oxide surface by forming carboxylate complexes [71]. In many applications, there is no need for further treatment, since the NPs are stable in polar solvents [71]; however, in order to increase water stability and avoid particle aggregation, several polymers can be added to the reaction mixture like PVP and poly(ethylene imine) [67]. Many approaches also include functional materials like carbon nanotubes (CNT) in order to yield magnetic composites with enhanced properties [72]. This method has been used for Fe3O4 as well as for other nanospinels such as Co and Zn ferrites [13, 18, 70, 73].

#### 3.4. Hydrothermal and solvothermal synthesis

Hydrothermal methodology consists of the formation of an aqueous (or aqueous-alcoholic) solution of the metal salts followed by the addition of a base until basic pH is reached. The resulting mixture is then transferred to a pressurized autoclave and subjected to T > 180 C for many hours. The mechanism involves the initial formation of metal hydroxides, which are oxidized and converted into the crystalline spinel ferrite due to the thermal treatment at high pressures. The resulting NPs have high crystallinity and an acceptable narrow size distribution. Particle size and shape can be effectively tuned by varying the metal concentration, solvent composition, temperature and reaction time. The addition of surfactants like CTAB [74] and poly(ethylene glycol) (PEG) [75] can change the shape of NPs and aid to control the growth and avoid agglomeration. This method can be adapted for the in situ synthesis of ferrite composites with functional materials like rGO [76, 77]; after base addition, metal hydroxides are adsorbed onto GO, and both spinel crystallization and GO reduction (by the action of supercritical water) occur during hydrothermal treatment. As a result, the size of near-spherical Zn ferrite NPs decreases as the GO content increases. On the contrary, the absence of GO sheets yielded NPs with a rod-like shape [78]. An alternative procedure by Komarneniet al. [79] involves hydrothermal treatment under microwave radiation, leading to a drastic reduction in reaction times to just few minutes; this approach was useful for many ferrites.

Solvothermal synthesis can be understood as a modified hydrothermal process where water is replaced by an organic solvent. For instance, n-octanol along with sodium dodecylbenzenesulfonate has been employed for the preparation of mixed ferrite NPs of Ni and Co with several compositions and varying sizes (7–16 nm), which was tuned as a function of the reaction time [26]. OA can also be used as a steric stabilizer in the reaction mixture using n-pentanol as a solvent [22]; increasing OA content decreases D from 19 to 5 nm and changes the particle morphology from nanoplatelets to well-dispersed nanospheres. Other approaches reported the use of diol molecules as solvents [25, 80, 81]. Bastami et al. [80] introduced PEG and PVP as polymeric surfactants, which bind preferentially at the surface of the near normal (i = 0.2) MnFe2O4 spinel ferrite compared to the inverse spinel Fe3O4 NPs synthesized similarly. As a consequence, an increase in D was noted for magnetite relative to Mn ferrite. Such behavior can be rationalized by taking into account the larger content of highly reactive Fe(III) species in octahedral coordination in the normal ferrite. The same procedure has been used for including GO in the reaction mixture [81]; since polyols can act as reducing agents under these conditions, the simultaneous reduction of GO and the formation of ferrite NPs are verified. Other methods used for synthesizing ferrite NPs include sol-gel [28, 82, 83], micro-emulsion [10, 36, 84], biogenic [85, 86], auto-ignited combustion [87–89], electrochemical [90] and mechanical [91] methods.

## 4. Functionalization

tuned as BA concentration is increased; the nature of the weak intermolecular interactions between adsorbed BA molecules seems to play a key role to bring about the final morphology. One drawback of the thermal decomposition method is that the as-synthesized ferrite NPs do not disperse in water due to the hydrophobic surfactant adsorbed onto the surface, which leads to further phase transfer steps. To obtain directly water-soluble NPs, the groups of Li [29] and Verma [58] introduced a variation in which 4–5 nm Fe and Co ferrites are obtained through the decomposition of metal acetylacetonates in the presence of pyrrolidones that act

This is a variation of the thermal decomposition in which a given polyol acts as a high-boiling solvent, reducing and stabilizer agent. Metal precursors are generally organic complexes like acetylacetonates and other carboxylate complexes [65]. Given that the reaction mixture is refluxed at the boiling point of the polyol, changing either the kind or concentration of the polyol leads to different particle D values [66–68], generally between 4 and 15 nm. The obtained NPs have a narrow size distribution and high crystallinity, although particle agglomeration could occur. There is question about the nature of the molecules adsorbed at the particle surface; some authors have stressed that polyol anchoring to the surface occurs through R-O interactions [69, 70], but others have claimed that at high temperatures, hydroxyl groups are oxidized to carboxylic acids, which are further adsorbed onto the oxide surface by forming carboxylate complexes [71]. In many applications, there is no need for further treatment, since the NPs are stable in polar solvents [71]; however, in order to increase water stability and avoid particle aggregation, several polymers can be added to the reaction mixture like PVP and poly(ethylene imine) [67]. Many approaches also include functional materials like carbon nanotubes (CNT) in order to yield magnetic composites with enhanced properties [72]. This method has been used

for Fe3O4 as well as for other nanospinels such as Co and Zn ferrites [13, 18, 70, 73].

Hydrothermal methodology consists of the formation of an aqueous (or aqueous-alcoholic) solution of the metal salts followed by the addition of a base until basic pH is reached. The resulting mixture is then transferred to a pressurized autoclave and subjected to T > 180 C for many hours. The mechanism involves the initial formation of metal hydroxides, which are oxidized and converted into the crystalline spinel ferrite due to the thermal treatment at high pressures. The resulting NPs have high crystallinity and an acceptable narrow size distribution. Particle size and shape can be effectively tuned by varying the metal concentration, solvent composition, temperature and reaction time. The addition of surfactants like CTAB [74] and poly(ethylene glycol) (PEG) [75] can change the shape of NPs and aid to control the growth and avoid agglomeration. This method can be adapted for the in situ synthesis of ferrite composites with functional materials like rGO [76, 77]; after base addition, metal hydroxides are adsorbed onto GO, and both spinel crystallization and GO reduction (by the action of supercritical water) occur during hydrothermal treatment. As a result, the size of near-spherical Zn ferrite NPs decreases as the GO content increases. On the contrary, the absence of GO sheets yielded NPs with a rod-like shape [78]. An alternative procedure by Komarneniet al. [79] involves hydrothermal treatment under microwave

either as solvents or as hydrophilic stabilizing agents.

192 Magnetic Spinels- Synthesis, Properties and Applications

3.4. Hydrothermal and solvothermal synthesis

3.3. Polyol method

Surface functionalization of nanostructured ferrites is a crucial step in the design of nanodevices for many applications since proper functionalization determines the final use and allows control over the physico-chemical processes at the surface, thus tuning several magnetic, optical and electrical properties in the desired direction. Although several synthetic methods allow in situ functionalization of obtained ferrite NPs, this approach is not always enough, and postsynthetic surface functionalization becomes necessary. For example, biomedical and environmental applications require hydrophilic NPs with definite chemical groups. The crucial feature that allows for surface functionalization is the availability of superficial transition metal d orbitals acting as Lewis acids in the presence of donor ligands. Fortunately, spinel ferrite surface is reactive toward many chemical groups, which provide room for multiple combinations. Ligands comprise many low and high molecular weight compounds [92]; functional groups available for surface complexation include carboxylic acids [93–95], amides [27, 29], hydroxyl [70, 96], phosphonic [73, 97] and hydroxamic acids [73].

There are mainly three approaches to make hydrophilic functional NPs: (i) ligand exchange reaction, (ii) silica coating and (iii) polymer coating. Ligand exchange reactions effectively transfer hydrophobic particles to aqueous medium by the replacement of hydrophobic ligands with hydrophilic ones, without affecting the magnetic core considerably. However, for some applications, magnetic NPs can also be transferred from polar to nonpolar mediums [98, 99]. Small ligands stabilize the NPs mainly by coulombic repulsion of ionized groups, like quaternary ammonium cations and carboxylates [100]; charged groups not only stabilize the magnetic suspension but also reinforce water affinity by facile solvation. Conversely, macromolecular ligands stabilize NPs by interparticle steric repulsions due to extended conformations that they can adopt in contact with good solvents [101]. In cases when the polymer carries ionisable groups, as PAA [34, 93, 101], coulombic repulsions enhance their capabilities as a stabilizer.

Silica coating has the advantage that provides excellent chemical stability to the magnetic core while preventing magnetic interactions, which is traduced into colloidal stability. Following the hydrolysis-condensation method established by Stöber [102], it is possible to achieve silica shells with controlled thickness by careful addition of tetraethyl orthosilicate (TEOS) to the NP dispersion without the appearance of individual silica particles, which in turn allows for a finetuning of magnetic interactions [103]. Furthermore, silica coating can be functionalized with several organosilanes containing suitable groups like -SH [104–106] and -NH [105], as depicted in Figure 2.

Figure 2. Scheme of the synthesis of amine and thiol functionalization of core magnetite NPs protected with a silica shell (Adapted from Ref. [105] with permission of the Royal Society of Chemistry).

Two main routes for polymer coating of NPs [107, 108] are: (i) functionalization of the NP surface with a molecule that acts as an initiator for further interfacial-controlled polymerization [109, 110] and (ii) synthesis of the polymer as the first step followed by surface anchoring [111–114]. The latter is simpler and allows for a wide variety of macromolecules, provided they bear suitable functional groups for surface binding. The former, although more laborious, has the advantage that it is possible to control the surface density of the grafted polymer and the length of the growing chains [115]. A shortcoming concerning macromolecular coating of magnetic NPs emerges when high mass magnetizations are required. Since polymers do not contribute to magnetization, mass magnetization of highly functionalized NPs drops noticeably, and so they might disable the whole system.

Conjugation after primary NP synthesis and water stabilization constitutes the final step prior to environmental and biomedical applications. It affords the ultimate precise chemical functions. Several strategies have been reported to achieve this goal entailing many known organic reactions [107]. For example, Zhao et al. [116] obtained hydroxamic acid-decorated Fe3O4/poly(acrylamide) (PAM) nanocomposites by treating the amide bonds with hydroxylamine solution and Zhao et al. [117] introduced amine groups by reacting ethylenediamine with a polymer containing epoxy moieties previously attached to Fe3O4 NPs. A very similar approach was employed in order to obtain thiol groups by adding NaSH to episulfide moieties [118]. Amide and ester formation are nice strategies to achieve NPs conjugation; in these reactions, carbodiimides are used as effective coupling agents. Following this strategy, Ren et al. [119] incorporated EDTA ligands to Fe3O4@SiO2@chitosan particles through amide bonds between EDTA –COO moieties and –NH2 groups in chitosan shell (see Figure 3), while Ge et al. [120] incorporated a polycarboxylic acid into amine-decorated magnetite. Rare sulfur-containing functional groups have been incorporated from chitosan modification, as is the case of xanthate-decorated magnetite NPs [121]. Thiol-ene reactions have also employed to include phosphonic acid moieties in a thioldecorated nanoplatform [122]; the kinetics and efficiency of this reaction prevent phosphonic acid depletion due to surface binding. The reaction is depicted in Figure 4.

ligands stabilize NPs by interparticle steric repulsions due to extended conformations that they can adopt in contact with good solvents [101]. In cases when the polymer carries ionisable groups, as PAA [34, 93, 101], coulombic repulsions enhance their capabilities as a stabilizer.

Silica coating has the advantage that provides excellent chemical stability to the magnetic core while preventing magnetic interactions, which is traduced into colloidal stability. Following the hydrolysis-condensation method established by Stöber [102], it is possible to achieve silica shells with controlled thickness by careful addition of tetraethyl orthosilicate (TEOS) to the NP dispersion without the appearance of individual silica particles, which in turn allows for a finetuning of magnetic interactions [103]. Furthermore, silica coating can be functionalized with several organosilanes containing suitable groups like -SH [104–106] and -NH [105], as depicted

Two main routes for polymer coating of NPs [107, 108] are: (i) functionalization of the NP surface with a molecule that acts as an initiator for further interfacial-controlled polymerization [109, 110] and (ii) synthesis of the polymer as the first step followed by surface anchoring [111–114]. The latter is simpler and allows for a wide variety of macromolecules, provided they bear suitable functional groups for surface binding. The former, although more laborious, has the advantage that it is possible to control the surface density of the grafted polymer and the length of the growing chains [115]. A shortcoming concerning macromolecular coating of magnetic NPs emerges when high mass magnetizations are required. Since polymers do not contribute to magnetization, mass magnetization of highly functionalized NPs drops notice-

Figure 2. Scheme of the synthesis of amine and thiol functionalization of core magnetite NPs protected with a silica shell

Conjugation after primary NP synthesis and water stabilization constitutes the final step prior to environmental and biomedical applications. It affords the ultimate precise chemical functions. Several strategies have been reported to achieve this goal entailing many known organic reactions [107]. For example, Zhao et al. [116] obtained hydroxamic acid-decorated Fe3O4/poly(acrylamide) (PAM) nanocomposites by treating the amide bonds with hydroxylamine solution and Zhao et al. [117] introduced amine groups by reacting ethylenediamine with a polymer containing epoxy moieties previously attached to Fe3O4 NPs. A very similar approach was employed in order to obtain thiol groups by adding NaSH to episulfide moieties [118]. Amide and ester formation are nice strategies to achieve NPs conjugation; in these reactions, carbodiimides are

ably, and so they might disable the whole system.

(Adapted from Ref. [105] with permission of the Royal Society of Chemistry).

in Figure 2.

194 Magnetic Spinels- Synthesis, Properties and Applications

Figure 3. Synthesis of EDTA-containing magnetic NPs by amidation of chitosan (Adapted from Ref. [119] with permission of Elsevier).

Figure 4. Synthesis of diphosphonic acid-containing magnetic NPs by a facile thiol-ene reaction (Adapted from Ref. [122] with permission of the American Chemical Society).

Coordination reactions like MOF construction have also been developed at ferrite surface. Fe3O4 NPs decorated with carboxyl groups were conjugated with a zeolitic imidazolate framework (ZIF-8) for the adsorption of contaminants [123, 124]. Such MOF was grown in a step-bystep assembly, initiated by the Zn2+ chelation to the oxide surface through carboxyl groups, resulting in a magnetic core surrounded by the ZIF shell. By varying the number of growth cycles, it is possible to tune the thickness of the MOF shell and hence, the interparticle distance between the magnetic cores. Another inorganic reaction at the interface of magnetic NPs reported recently [125] consists of the deposition of hydrous lanthanum oxide over Fe3O4@SiO2

Figure 5. Two methods for thiol protection. Top: Steric hindrance by surface PEG grafting (Adapted from Ref. [130] with permission of the American Chemical Society). Bottom: Reversible thiol oxidation to disulfide bridges (Reproduced from Ref. [34] with permission of Elsevier).

core-shell NPs simply by adding LaCl3 at basic pH in the presence of the magnetic material. Nanostructures composed of ferrites and noble metals have interesting and promising optical and magnetic properties; the synthesis of such materials can be easily performed by reduction of the corresponding metal salt in the presence of ferrite NPs [126].

In the case of surface thiol-decorated nanostructures, special care must be taken, since free thiol groups are prone to be oxidized during the synthetic procedures. For example, several papers have reported the oxidation of DMSA and cysteine to disulfide and sulfoxide compounds in the presence of Fe3O4 NPs [24, 127–129]; these undesirable processes not only reduce the effective amount of –SH groups but also could alter magnetite phase. To overcome this drawback, Maurizi et al. [130] grafted PEG chains onto the oxide surface as steric barriers in order to avoid the formation of intermolecular disulfide bridges between adjacent DMSA molecules. More recently, Odio et al. [34] designed a novel PAA copolymer containing disulfide bridges; after Fe3O4 functionalization with this polymer, resulting NPs were treated with tributyl phosphine in order to reduce disulfide bridges to free –SH groups. Both strategies are depicted in Figure 5.

## 5. Coordination chemistry at the surface

Coordination reactions like MOF construction have also been developed at ferrite surface. Fe3O4 NPs decorated with carboxyl groups were conjugated with a zeolitic imidazolate framework (ZIF-8) for the adsorption of contaminants [123, 124]. Such MOF was grown in a step-bystep assembly, initiated by the Zn2+ chelation to the oxide surface through carboxyl groups, resulting in a magnetic core surrounded by the ZIF shell. By varying the number of growth cycles, it is possible to tune the thickness of the MOF shell and hence, the interparticle distance between the magnetic cores. Another inorganic reaction at the interface of magnetic NPs reported recently [125] consists of the deposition of hydrous lanthanum oxide over Fe3O4@SiO2

196 Magnetic Spinels- Synthesis, Properties and Applications

Figure 5. Two methods for thiol protection. Top: Steric hindrance by surface PEG grafting (Adapted from Ref. [130] with permission of the American Chemical Society). Bottom: Reversible thiol oxidation to disulfide bridges (Reproduced from

Ref. [34] with permission of Elsevier).

Since the nature of the metal-ligand interactions at the interface of the ferrites plays a key role in the properties of NPs, efforts have been devoted to unravel the structure and implications of the surface complexes occurring for different types of ligands. For this purpose, spectroscopic techniques like FTIR, XPS, EXAF and XANES are usually employed [22, 24, 58, 122, 131–133]. Mössbauer spectroscopy has also been used since iron spectra are sensitive to spin reorganization after ligand binding and to the kind of iron site that participates in the surface complexes [23, 58, 131].

Specifically, Daou et al. [131] showed that phosphate ligands bind to magnetite surface by forming bidentate binuclear complexes with octahedral Fe(III) cations. In contrast, Costo et al. [21] suggested that phosphonates bind to the surface through mono-dentate ligands. For carboxylates, it is reported that in solvothermal and thermal decomposition methods, OA is anchored to the surface by bridging bidentate complexes [22, 24, 134], a conclusion supported by a theoretical DFT study showing that the bidentate mode was the most stable configuration of iron-oleate complex no matter the surface plane exposed to the ligand [135]. However, after ligand exchange reactions with other carboxylic acids, carboxylate complexes could form chelates [24, 34, 136]. The use of amine functional groups showed that long-chain amine molecules adsorb at the surface of ferrites by N-metal interactions [16, 137]. In contrast, if both OA and OAm are used in the ferrite synthesis, the mode of coordination of each functional group could depend on the concentration and molar ratio of the surfactants. Thus, if the surfactants are diluted in N-methyl 2-pyrrolidone with a molar ratio 1:1, Verma and Pravarthana [58] suggested by means of IR analysis that OA complexes retain the bidentate mode while OAm molecules appear protonated and associated to the surface through coulombic interactions. However, XPS studies of Wilson and Langell [133] indicated that if the reaction is performed without a solvent and a higher OAm proportion is employed, OAm is anchored to the surface by Fe-N coordination, while OA binds to the surface though a monodentate complex. Using IR studies, alcohols are reported to be anchored to the oxide surface by metal-OH interactions [70]. Besides, ligands with thiol moieties anchor to magnetite surface by Fe(II)-S interactions, as suggested by IR and XPS studies [24, 138]. Finally, cyclic amides like PVP and 2-pyrrolidone seem to interact with ferrite surfaces through the carbonyl groups [27, 29]. Since this subject is not yet completely understood, computational methods are likely necessary to unravel the structural and electronic properties of surface complexes.

## 6. Environmental applications

In this section, we focus on two applications of nanostructured spinel ferrites for environmental remediation technologies in connection with water decontamination: adsorption and oxidation technologies.

#### 6.1. Adsorption technologies for removal of inorganic and organic contaminants

Adsorption is often the most suitable choice for removal of toxic substances in drinking or waste waters, mainly due to its simplicity and high efficiency; the main disadvantage is the sorbent separation after the adsorption process, which can become tedious and energy consuming. However, the use of magnetic materials for adsorption makes the task of sorbent separation easier by allowing magnetic decantation with a permanent magnet. The high surface area of ferrite NPs along with their room temperature superparamagnetism and the great versatility for binding specific functional groups on their surfaces for specific contaminants makes them ideal candidates for the design and development of innovative adsorption strategies. Although several recent reviews covering this subject are available [6, 7, 139, 140], these have generally focused on the thermodynamics and kinetics of the adsorption process and relatively less attention has been paid to unravel the atomic and molecular nature of the interactions occurring at the interface. Although this is a difficult task, this information is crucial for the improvement and optimization of the nanoadsorbent.

#### 6.1.1. Heavy metal cations

Heavy metal cations, found in natural and waste waters resulting from industrial activities, comprise a wide family of hazardous substances with a high impact on human health [141]. Here, we concentrate on those reported studies with a focus on two directions: (i) improving the adsorption capacity and/or selectivity toward a given contaminant by surface functionalization of ferrite NPs and (ii) shedding light on the adsorption mechanism at a molecular and atomic level.

#### 6.1.1.1. Amine-functionalized nanosystems

Fe3O4 NPs functionalized with several amino-containing polymers were tested as Cr(VI) and Cu(II) sorbents in aqueous medium [142], showing the increase of adsorption capacity for both cations with the number of –NH moieties in the ligand incorporated to the magnetic nanoplatform. Adsorption and spectroscopic data suggested that metal removal involves coulombic interactions, ion exchange processes and formation of complexes between amine groups and metal ions, although the structure of such complexes was not revealed. Similar results were reported by Huang and Chen [113], in which Fe3O4@PAA NPs decorated with amine groups were proved as a good adsorbent for several heavy metals with positive and negative charges; based on pH studies, authors suggested that cations are adsorbed through chelate complexes while anions are incorporated after ion exchange mechanisms. New insights about Cr(VI) adsorption with an amino-decorated magnetic sorbent were reported by Zhao et al. [143]. This approach comprises nanocomposites of Fe3O4 and amino-functionalized GO sheets. Based on XPS measurements, authors suggested that after attractive coulombic interactions between chromate species and protonated amino groups, a fraction of Cr(VI) is reduced to Cr(III), which further forms amino-complexes; the source of electron for Cr(VI) reduction is provided by the GO sheets. Without extra evidence, these results should be taken with care because XPS deconvolution was not rigorous.

Amino-functionalized Fe3O4 NPs were tested as a sorbent for Cu(II), Cd(II) and Pb(II) [144]. Adsorption decreased at acid pH values and adsorption capacity for Cu(II) was higher than that for softer Lewis acids Cd(II) and Pb(II). Both results, along with thermodynamic and kinetic data, could indicate the prevalence of coulombic and complexing reactions between surface –NH moieties and the cations. Similar results were presented in other reports [52, 145]. Co-ferrite NPs coated with a polystyrene shell modified with amino and thioether groups were tested for Hg(II) adsorption [146]. Authors proposed the Hg(II) complexation by these functional groups, followed by partial reduction of Hg(II) to Hg(I), although no proof for this mechanism was presented.

#### 6.1.1.2. Carboxyl-functionalized nanosystems

anchored to the surface by Fe-N coordination, while OA binds to the surface though a monodentate complex. Using IR studies, alcohols are reported to be anchored to the oxide surface by metal-OH interactions [70]. Besides, ligands with thiol moieties anchor to magnetite surface by Fe(II)-S interactions, as suggested by IR and XPS studies [24, 138]. Finally, cyclic amides like PVP and 2-pyrrolidone seem to interact with ferrite surfaces through the carbonyl groups [27, 29]. Since this subject is not yet completely understood, computational methods are likely

In this section, we focus on two applications of nanostructured spinel ferrites for environmental remediation technologies in connection with water decontamination: adsorption and oxi-

Adsorption is often the most suitable choice for removal of toxic substances in drinking or waste waters, mainly due to its simplicity and high efficiency; the main disadvantage is the sorbent separation after the adsorption process, which can become tedious and energy consuming. However, the use of magnetic materials for adsorption makes the task of sorbent separation easier by allowing magnetic decantation with a permanent magnet. The high surface area of ferrite NPs along with their room temperature superparamagnetism and the great versatility for binding specific functional groups on their surfaces for specific contaminants makes them ideal candidates for the design and development of innovative adsorption strategies. Although several recent reviews covering this subject are available [6, 7, 139, 140], these have generally focused on the thermodynamics and kinetics of the adsorption process and relatively less attention has been paid to unravel the atomic and molecular nature of the interactions occurring at the interface. Although this is a difficult task, this information is

Heavy metal cations, found in natural and waste waters resulting from industrial activities, comprise a wide family of hazardous substances with a high impact on human health [141]. Here, we concentrate on those reported studies with a focus on two directions: (i) improving the adsorption capacity and/or selectivity toward a given contaminant by surface functionalization of ferrite

Fe3O4 NPs functionalized with several amino-containing polymers were tested as Cr(VI) and Cu(II) sorbents in aqueous medium [142], showing the increase of adsorption capacity for both cations with the number of –NH moieties in the ligand incorporated to the magnetic nanoplatform. Adsorption and spectroscopic data suggested that metal removal involves

NPs and (ii) shedding light on the adsorption mechanism at a molecular and atomic level.

necessary to unravel the structural and electronic properties of surface complexes.

6.1. Adsorption technologies for removal of inorganic and organic contaminants

crucial for the improvement and optimization of the nanoadsorbent.

6. Environmental applications

198 Magnetic Spinels- Synthesis, Properties and Applications

dation technologies.

6.1.1. Heavy metal cations

6.1.1.1. Amine-functionalized nanosystems

Several reports have focused on the adsorption of heavy metal cations by EDTA-modified magnetic nanosystems in order to take advantage of the high chelating ability of this multifunctional ligand. The key role of EDTA has been confirmed since the adsorption capacity decreases when no EDTA was used in the preparation of the sorbents. Indeed, Ren et al. [119] noted that the adsorption capacity follows the same order of the metal-EDTA complex stability: Cu(II) > Pb(II) > Cd(II). Similar conclusions are drawn from recent reports [147, 148] in which the formation of metal-carboxylates was verified by means of FTIR measurements. Carboxyl-decorated NPs from acrylic and crotonic acid copolymer were tested for Cu(II), Pb(II), Zn(II) and Cd(II) [120]. Although authors did not show spectroscopic evidences of metal-carboxylates interactions, the maximum adsorption capacity increases with the increase in Lewis acid hardness of the ion tested (Cu(II) > Zn(II) ≈ Pb(II) > Cd(II)). Likewise, Mahdavian et al. [149] studied the adsorption behavior of Pb(II), Cu(II), Ni(II) and Cd(II) with a nanoplatform consisting of PAA chains grown at the surface of magnetite NPs and found that the metal uptake increases with pH, suggesting chelate formation. Other carboxyl-based magnetite NPs for Pb(II) removal can be found elsewhere [150].

#### 6.1.1.3. Thiol and other sulfur-containing compounds functionalized nanosystems

Fe3O4 NPs functionalized with a polythiolated ligand was probed as Hg(II) adsorbent [118]. XPS studies supported the occurrence of Hg(II)-S interactions and the simultaneous reduction of Hg(II) to Hg(I) likely at the expense of Fe(II) cations at the surface of the magnetic core. Curiously, no sign of thiol oxidation was encountered. Recently, Wang et al. [151] tested a simpler Fe3O4@SiO2-RSH nanoplatform toward Hg(II), Pb(II) and Cd(II) ions. Although no mechanism was elucidated, the fact that the adsorption capacities followed the order Hg(II) >> Cd(II) > Pb(II) might be indicative of two phenomena: (i) metal uptake is governed by softsoft interactions between the cations and the thiolate groups and (ii) for the case of Hg(II) ions, adsorption involves reduction. A similar CoFe2O4@SiO2-RSH nanosystem for Pb(II) removal has been employed, but no mechanism was proposed [152].

Zhu et al. [121] designed a Fe3O4/chitosan-OC(=S)SH platform for Pb(II), Cu(II) and Zn(II) adsorption. FTIR and XPS suggested that metal-ligand interactions comprise both the N atom of the residual chitosan amine groups and the sulfur groups of xanthate moieties. However, neither the specific role of each functional group nor the structure of such complexes was elucidated. Based on the adsorption capacity order Cu(II) >> Pb(II) ≈ Zn(II), it is likely that relative hard chitosan N groups play an important role in the overall performance of the adsorbent. Alternatively, Zhang et al. [72] chose nanocomposites of Fe3O4 and thiolated CNTs as sorbents for Pb(II) and Hg(II) ions. Thiol grafting material proved to be a better adsorbent than Fe3O4/CNT composites. Adsorption capacity is larger for softer Lewis acid Hg(II), which could imply the occurrence of metal-thiol complexes. A similar trend was found in another report [106].

Fe3O4 NPs functionalized with a copolymer obtained by the partial modification of PAA with thio-salicyl-hydrazide were tested for several divalent cations [153]. This system contains both soft (thiol) and hard (carboxyl and amine) moieties, which might explain the good adsorption properties toward soft Cd(II) and hard Co(II) cations. Regarding Pb(II) uptake, XPS studies confirmed the presence of Pb-S interactions; it is interesting that only one contribution was proposed for the deconvolution of the Pb 4f spectrum, which implies that there is only one coordination environment for Pb(II) cations. The prevalence of Pb-S interactions is coherent with the small interference effect produced by alkaline/earth metals, since these hard cations largely prefer hard ligands.

Surface ion imprinting techniques can also be used for efficient and selective sequestration of heavy metal cations. Guo et al. [154] added Fe3O4@SiO2 NPs to a solution containing the Pb-MPTS complex as a template. Condensation of silane groups followed by Pb(II) removal with HCl results in imprinted cavities with the proper thiol configuration (Figure 6). This nanosystem was shown to be a good adsorbent toward Pb(II) ions with excellent selectivity over other heavy metals like Cu(II), Zn(II) and Co(II). In this case, selectivity is not only ruled by the chemical affinity between cation and thiols but it also depends on the ionic radius, coordination number and coordination geometry.

Yantasee et al. [155] employed Fe3O4@DMSA NPs for the removal of several cations, suggesting that metal adsorption was driven by free –SH groups, while –COOH groups were anchored to the magnetic core. However, no evidence regarding either the state of the sulfur before and after metal incorporation or the nature of metal-S complex was presented. Afterward, Odio et al. [24] tested the adsorption capabilities of the same nanoplatform aiming at Au (III). By means of XPS and UV-vis analysis, they found that before adsorption, DMSA ligands are mostly oxidized to disulfide bridges and Au(III) could adsorb onto Fe3O4@DMSA NPs by three possible ways: (i) chelation with free –COOH moieties; (ii) reduction to Au<sup>0</sup> sub-nanometer clusters triggered by surface Fe(II) oxidation in the bare sectors of the NPs and (iii) extensive reduction to Au<sup>0</sup> nanoclusters in the region covered by the organic ligand, which is caused likely due to oxidation of disulfide bridges to –SOx species. The process is depicted in the upper part of Figure 7. On the contrary, when Pb(II) was tested with the same material, the results indicated that neither Pb reduction nor Pb-S interactions contribute to the adsorption process, but it is solely caused by the occurrence of chelating carboxylates and oxo-complex at bare sectors [34]. In order to unravel the actual role of both –SH and –COOH functions in the adsorption of hardly reducible divalent cations, the same group studied the adsorption of Pb (II) and Cd(II) by a novel nanoplatform based on Fe3O4 NPs capped with a copolymer with pendant free –COOH and –SH groups [34]; thiol moieties were protected during the synthesis to avoid oxidation and were regenerated just before the adsorption experiments. A detailed XPS analysis indicated that both metal-carboxylate and metal-thiolate interactions are verified during adsorption. In addition, it was shown that although both cations showed higher affinity to thiols, this tendency was more pronounced for Cd(II), that is, Pb(II) is less selective, in agreement with its borderline softness characteristics. Nevertheless, the actual structure of the surface metal complexes is still elusive.

of Hg(II) to Hg(I) likely at the expense of Fe(II) cations at the surface of the magnetic core. Curiously, no sign of thiol oxidation was encountered. Recently, Wang et al. [151] tested a simpler Fe3O4@SiO2-RSH nanoplatform toward Hg(II), Pb(II) and Cd(II) ions. Although no mechanism was elucidated, the fact that the adsorption capacities followed the order Hg(II) >> Cd(II) > Pb(II) might be indicative of two phenomena: (i) metal uptake is governed by softsoft interactions between the cations and the thiolate groups and (ii) for the case of Hg(II) ions, adsorption involves reduction. A similar CoFe2O4@SiO2-RSH nanosystem for Pb(II) removal

Zhu et al. [121] designed a Fe3O4/chitosan-OC(=S)SH platform for Pb(II), Cu(II) and Zn(II) adsorption. FTIR and XPS suggested that metal-ligand interactions comprise both the N atom of the residual chitosan amine groups and the sulfur groups of xanthate moieties. However, neither the specific role of each functional group nor the structure of such complexes was elucidated. Based on the adsorption capacity order Cu(II) >> Pb(II) ≈ Zn(II), it is likely that relative hard chitosan N groups play an important role in the overall performance of the adsorbent. Alternatively, Zhang et al. [72] chose nanocomposites of Fe3O4 and thiolated CNTs as sorbents for Pb(II) and Hg(II) ions. Thiol grafting material proved to be a better adsorbent than Fe3O4/CNT composites. Adsorption capacity is larger for softer Lewis acid Hg(II), which could imply the occurrence of metal-thiol complexes. A similar trend was found in another

Fe3O4 NPs functionalized with a copolymer obtained by the partial modification of PAA with thio-salicyl-hydrazide were tested for several divalent cations [153]. This system contains both soft (thiol) and hard (carboxyl and amine) moieties, which might explain the good adsorption properties toward soft Cd(II) and hard Co(II) cations. Regarding Pb(II) uptake, XPS studies confirmed the presence of Pb-S interactions; it is interesting that only one contribution was proposed for the deconvolution of the Pb 4f spectrum, which implies that there is only one coordination environment for Pb(II) cations. The prevalence of Pb-S interactions is coherent with the small interference effect produced by alkaline/earth metals, since these hard cations

Surface ion imprinting techniques can also be used for efficient and selective sequestration of heavy metal cations. Guo et al. [154] added Fe3O4@SiO2 NPs to a solution containing the Pb-MPTS complex as a template. Condensation of silane groups followed by Pb(II) removal with HCl results in imprinted cavities with the proper thiol configuration (Figure 6). This nanosystem was shown to be a good adsorbent toward Pb(II) ions with excellent selectivity over other heavy metals like Cu(II), Zn(II) and Co(II). In this case, selectivity is not only ruled by the chemical affinity between cation and thiols but it also depends on the ionic radius,

Yantasee et al. [155] employed Fe3O4@DMSA NPs for the removal of several cations, suggesting that metal adsorption was driven by free –SH groups, while –COOH groups were anchored to the magnetic core. However, no evidence regarding either the state of the sulfur before and after metal incorporation or the nature of metal-S complex was presented. Afterward, Odio et al. [24] tested the adsorption capabilities of the same nanoplatform aiming at Au (III). By means of XPS and UV-vis analysis, they found that before adsorption, DMSA ligands

has been employed, but no mechanism was proposed [152].

200 Magnetic Spinels- Synthesis, Properties and Applications

report [106].

largely prefer hard ligands.

coordination number and coordination geometry.

Figure 6. Surface ion imprinting technique for selective adsorption of Pb(II) ions onto magnetic NPs (Adapted from Ref. [154] with permission of Elsevier).

6.1.1.4. Other functional groups and particles with a bare surface

Fe3O4/PAM nanocomposites functionalized with hydroxamic acid moieties were shown to adsorb Pb(II), Cd(II), Co(II) and Ni(II) ions by forming bidentate chelating complexes [116]. The key role of hydroxamic groups was demonstrated, which agrees with the fact that the stability constant of metal-hydroxamic complexes follows the same order as the maximum adsorption capacity. The structure of these surface complexes was determined from IR and DFT studies and the system was selective toward Pb(II) uptake.

Rutledge et al. [122] designed a nanoplatform decorated with diphosphonic acid and thiol moieties for Pb(II), Hg(II) and f-block elements like La(III) and Eu(III) adsorption. Comparison between different adsorption sites suggested that both hard Lewis acid cations La(III) and Eu (III) are efficiently adsorbed by a hard Lewis base as diphosphonate groups, but they are inert toward thiol-mediated binding, in accordance with the soft nature of this Lewis base. However, Hg(II) uptake showed the inverse tendency, being very sensitive to –SH moieties, but a kind of unreactive toward diphosphonate basic functions. On the contrary, Pb(II) did not display such selectivity toward a particular functional group, but it showed a synergic effect upon ion uptake. The same behavior for Pb(II) was mentioned above: the fact that Pb(II) cation

Figure 7. Top: Likely distribution of Au atoms during adsorption onto Fe3O4@DMSA NPs; different colors correspond to distinctive Au 4f XPS signals. Bottom: Likely distribution of Pb(II) and Cd(II) cations during adsorption onto the thiol- and carboxyl-containing Fe3O4 NPs; note the preponderance of metal-thiol interactions over metal-carboxylate ones (Adapted from Ref. [34] with permission of Elsevier).

can bind a wide variety of hard and soft ligands is related with unique electronic properties steaming from relativistic effects [156]. Further theoretical and experimental investigations on this issue are still needed.

Given that surface magnetite NPs biosynthesized by microorganisms are richer in Fe(II) content with respect to stoichiometric Fe3O4, this biomaterial has been tested for the adsorption and reduction of toxic oxyanions containing Cr(VI) and m99Tc(VII) [85]. Results confirm that biomagnetite is a better absorber compared to a commercial magnetite of similar size, and the removal capacity changes with the particular iron substrate that was used for bacteria culture. The adsorption-reduction mechanism of chromate anions was studied by means of XPS and Xray magnetic circular dichroism (XMCD). Authors suggested that after the fast electron transfer reactions between Cr(VI) and surface Fe(II), Cr(III) ions are incorporated into the spinel structure and occupy octahedral interstices, thus forming a layer of ferrimagnetic CrFe2O4 spinel.

#### 6.1.2. Arsenic, phosphorous and fluoride

The key role of hydroxamic groups was demonstrated, which agrees with the fact that the stability constant of metal-hydroxamic complexes follows the same order as the maximum adsorption capacity. The structure of these surface complexes was determined from IR and

Rutledge et al. [122] designed a nanoplatform decorated with diphosphonic acid and thiol moieties for Pb(II), Hg(II) and f-block elements like La(III) and Eu(III) adsorption. Comparison between different adsorption sites suggested that both hard Lewis acid cations La(III) and Eu (III) are efficiently adsorbed by a hard Lewis base as diphosphonate groups, but they are inert toward thiol-mediated binding, in accordance with the soft nature of this Lewis base. However, Hg(II) uptake showed the inverse tendency, being very sensitive to –SH moieties, but a kind of unreactive toward diphosphonate basic functions. On the contrary, Pb(II) did not display such selectivity toward a particular functional group, but it showed a synergic effect upon ion uptake. The same behavior for Pb(II) was mentioned above: the fact that Pb(II) cation

Figure 7. Top: Likely distribution of Au atoms during adsorption onto Fe3O4@DMSA NPs; different colors correspond to distinctive Au 4f XPS signals. Bottom: Likely distribution of Pb(II) and Cd(II) cations during adsorption onto the thiol- and carboxyl-containing Fe3O4 NPs; note the preponderance of metal-thiol interactions over metal-carboxylate ones (Adapted

from Ref. [34] with permission of Elsevier).

DFT studies and the system was selective toward Pb(II) uptake.

202 Magnetic Spinels- Synthesis, Properties and Applications

In a careful spectroscopic study, Liu et al. [157] studied As(V) and As(III) adsorption onto 34 nm magnetite NPs, avoiding the presence of oxygen during the adsorption procedure. After EXAF and XANES analysis, they confirmed that arsenate is adsorbed as a bidentate binuclear corner-sharing complex, while arsenite binds to the surface through a tridentate hexanuclear corner-sharing complex (see the left panel of Figure 8). Also important, they noted, based on XPS analysis, that if anoxic conditions are fulfilled, no redox reactions involving arsenic species are verified. This result disagrees with other reports [158, 159] that claimed for redox reactions

Figure 8. Left: Geometry of surface complexes during As(V) and As(III) adsorption onto magnetite surface in anaerobic aqueous medium. Right: possible redox reactions occurring at the As-Fe3O4 interface when exposed to air (Reproduced from Ref. [157] with permission of the American Chemical Society).

during adsorption and generates a reasonable doubt about the role of the magnetite surface in the arsenic redox processes. On the contrary, after exposure of the adsorbed material to aerobic conditions, XPS analysis showed substantial amounts of As(V) in the As(III)-treated NPs and As (III) in the As(V)-treated NPs. In both samples, the Fe(II)/Fe(III) molar ratio was less than 0.5, denoting Fe3O4 oxidation to γ-Fe2O3. As(III) oxidation to As(V) can be caused by oxygen contact. On the other hand, As(V) reduction to As(III) was explained assuming that during magnetite oxidation, an electron flow from the core to the surface takes place; eventually, these electrons could cause As(V) reduction. Both processes are depicted in the right panel of Figure 8.

Zhang et al. [160] employed Fe3O4/activated carbon fiber nanocomposites for As(V) removal. Based on XPS studies, authors claimed the occurrence of inner-sphere bidentate complexes between the mono-protonated anion and surface oxygen anions from either the magnetite phase or the carbon fibers. In contrast, no redox reactions were claimed. Another carbonaceous material in conjunction with ferrites has also been tested for arsenic removal. Lingamdinne et al. [55] chose 30 nm porous NiFe2O4/rGO nanocomposites and concluded that the high removal efficiency is due to the extended porous structure of the composite, which favors adsorption. The authors claimed that adsorption involves both electrostatic attraction and surface-complexation reactions, but to our understanding, this issue was not totally clarified.

Another recent report uses Fe3O4@ZIF-8 as a sorbent [123]. In this case, arsenic adsorption is entirely caused by the ZIF-8 shell, while magnetite core only acts as a magnetic device to remove the contaminant in a facile and efficient way. Alternatively, magnetite particles encapsulated with calcium alginate were tested as an adsorbent for inorganic and organic As(V) species [159]. The authors found that inorganic species are better adsorbed than monomethyl arsenate. Based on IR and XPS measurements, they suggested that arsenic incorporation likely occurs through the partial reduction of As(V) to As(III) species and the oxidation of both alginate and magnetite. However, spectroscopic studies were not conclusive.

A recent report from Penke et al. [161] deals with As(III) and As(V) adsorption onto 20–30 nm Al-substituted NiFe2O4 NPs. Based on Raman, FTIR and XPS studies, the authors proposed that both species are adsorbed onto the ternary oxide surface through inner-sphere complexes. In addition, redox reactions between the spinel and arsenic species were claimed. However, this report does not clarify either the geometry of the formed surface complexes or the nature of the implicated redox reactions, although it showed that the replacement of Fe(III) cations by Al(III) cations enhances the adsorption properties of the oxide due to an increase in the number of surface hydroxyl groups. Similar conclusions were drawn from the work of Peng et al. [162]. In this case, Fe3O4@Cu(OH)2 core-shell nanostructures were tested as adsorbents for As(V). They stressed the key role of surface hydroxyl groups of the copper shell and suggested that arsenate complexes are formed at the surface. However, the proposed complex structure is not rigorously justified.

The use of metal hydroxides can be extended to other elements of group V like phosphorus. Thus, Lai et al. [125] employed a shell of hydrous lanthanum oxide incorporated into Fe3O4@SiO2 NPs to drive the removal of phosphate anions from the water medium. The adsorption was fast and efficient given the high affinity of La(III) species toward phosphate ligands. This methodology overcomes the use of bare Fe3O4. Fe3O4/polypyrrole nanocomposites were employed for removing fluoride anions without interference effects [163]. Based on thermodynamic and kinetic data, authors postulated an ion exchange mechanism.

#### 6.1.3. Dyes

during adsorption and generates a reasonable doubt about the role of the magnetite surface in the arsenic redox processes. On the contrary, after exposure of the adsorbed material to aerobic conditions, XPS analysis showed substantial amounts of As(V) in the As(III)-treated NPs and As (III) in the As(V)-treated NPs. In both samples, the Fe(II)/Fe(III) molar ratio was less than 0.5, denoting Fe3O4 oxidation to γ-Fe2O3. As(III) oxidation to As(V) can be caused by oxygen contact. On the other hand, As(V) reduction to As(III) was explained assuming that during magnetite oxidation, an electron flow from the core to the surface takes place; eventually, these electrons

204 Magnetic Spinels- Synthesis, Properties and Applications

could cause As(V) reduction. Both processes are depicted in the right panel of Figure 8.

Zhang et al. [160] employed Fe3O4/activated carbon fiber nanocomposites for As(V) removal. Based on XPS studies, authors claimed the occurrence of inner-sphere bidentate complexes between the mono-protonated anion and surface oxygen anions from either the magnetite phase or the carbon fibers. In contrast, no redox reactions were claimed. Another carbonaceous material in conjunction with ferrites has also been tested for arsenic removal. Lingamdinne et al. [55] chose 30 nm porous NiFe2O4/rGO nanocomposites and concluded that the high removal efficiency is due to the extended porous structure of the composite, which favors adsorption. The authors claimed that adsorption involves both electrostatic attraction and surface-complexation reactions, but to our understanding, this issue was not totally clarified.

Another recent report uses Fe3O4@ZIF-8 as a sorbent [123]. In this case, arsenic adsorption is entirely caused by the ZIF-8 shell, while magnetite core only acts as a magnetic device to remove the contaminant in a facile and efficient way. Alternatively, magnetite particles encapsulated with calcium alginate were tested as an adsorbent for inorganic and organic As(V) species [159]. The authors found that inorganic species are better adsorbed than monomethyl arsenate. Based on IR and XPS measurements, they suggested that arsenic incorporation likely occurs through the partial reduction of As(V) to As(III) species and the oxidation of both

A recent report from Penke et al. [161] deals with As(III) and As(V) adsorption onto 20–30 nm Al-substituted NiFe2O4 NPs. Based on Raman, FTIR and XPS studies, the authors proposed that both species are adsorbed onto the ternary oxide surface through inner-sphere complexes. In addition, redox reactions between the spinel and arsenic species were claimed. However, this report does not clarify either the geometry of the formed surface complexes or the nature of the implicated redox reactions, although it showed that the replacement of Fe(III) cations by Al(III) cations enhances the adsorption properties of the oxide due to an increase in the number of surface hydroxyl groups. Similar conclusions were drawn from the work of Peng et al. [162]. In this case, Fe3O4@Cu(OH)2 core-shell nanostructures were tested as adsorbents for As(V). They stressed the key role of surface hydroxyl groups of the copper shell and suggested that arsenate complexes are formed at the surface. However, the proposed complex structure is

The use of metal hydroxides can be extended to other elements of group V like phosphorus. Thus, Lai et al. [125] employed a shell of hydrous lanthanum oxide incorporated into Fe3O4@SiO2 NPs to drive the removal of phosphate anions from the water medium. The adsorption was fast and efficient given the high affinity of La(III) species toward phosphate ligands. This methodology overcomes the use of bare Fe3O4. Fe3O4/polypyrrole nanocomposites were employed for

alginate and magnetite. However, spectroscopic studies were not conclusive.

not rigorously justified.

Extensive use of organic dyes has become a serious environmental problem since this family of organic compounds is difficult to decompose and transforms to carcinogenic amines. A series of ferrite MFe2O4/rGO (M = Mn2+, Ni2+, Zn2+, Co2+) nanocomposites were tested as combined magnetic materials for adsorption and photocatalytic degradation of Methylene Blue (MB) and Rhodamine B (RhB) under visible light [81] (see Section 6.2). Authors devoted the high adsorption capacity and fast removal rate to the large surface area of the material. For this system, though electrostatic interactions cannot be ruled out, dye retention is mainly caused by the rGO sheets, comprising π-π stacking interactions between the aromatic moieties of the dyes and the extended π-conjugated regions in the graphene structure. The same mechanism was claimed earlier using Fe3O4/rGO nanocomposites for MB adsorption [164]; this report also tested other materials like activated carbon and multi-walled carbon nanotubes (MWCN).

Cobalt ferrites covered by PEG chains were shown to be good adsorbents for several dyes such as methyl orange (MO), MB and Congo red (CR) [75]. Adsorption data indicate that electrostatic interactions are not the prominent cause for adsorption; instead, H-bonding interactions between –OH groups of PEG and functional groups in the dyes seem to be the responsible cause. The interactions are depicted in Figure 9. H-bonding has also been claimed as the main interaction of several dyes with naked MnFe2O4 NPs [43].

In a recent work, Dolatkhah and Wilson [114] functionalized Fe3O4 NPs with chitosan grafted with PAA and poly(itaconic) acid (PIA) chains. This polymeric material displays reversible pH-responsive behavior, which was tested for MB adsorption. As the pH increases, the ionization of the chitosan-grafted acid groups also increases, favoring the expansion of the grafted chains and the ionic interactions with MB, since this dye is cationic. Hence, adsorption is favored. Afterward, the desorption of MB is accomplished simply by acidification until dyesorbent interactions become very weak and the polymeric chains no longer stabilize the colloid, leading to the collapse of the dye-free NPs. The process is represented in Figure 10.

#### 6.1.4. Aromatic compounds and other organic pollutants

Rodovalho et al. [98] employed mixed Mn and Co ferrite functionalized with carboxyl-terminated polydimethylsiloxane brushes for the adsorption of toluene in water. Significant hydrophobic interactions between toluene and polymeric chains lead to high adsorption capacity of the NPs. Moreover, authors exploited the suitable magnetic characteristic of this mixed ferrite for fast and efficient toluene desorption through hyperthermia treatment of the toluene-loaded nanoadsorbent. McCormick and Adriaens [86] employed ultra-small biogenic magnetite in the reductive transformation of CCl4, stressing the key role of octahedral Fe(II) cations and the importance of the electron hopping between Fe(II) and Fe(III) cations in the B sites, which enables a good electron mobility for the surface reduction of CCl4. Efficient adsorption of tetracycline and diclofenac by proper functionalization of magnetic nanostructures has also been reported [165, 166].

Figure 9. Proposed mechanism for CR adsorption onto PEG-functionalized MFe2O4 NPs (Adapted from Ref. [75] with permission of Elsevier).

Figure 10. Mechanism of reversible pH-responsive behavior (Adapted from Ref. [114] with permission of the American Chemical Society).

#### 6.2. Advanced oxidation technology

Advanced oxidation technologies consist of the assisted degradation of a given pollutant by using a source of highly oxidizing transient species. Such species are generally activated by the action of another substance that acts as a catalyst. Since the removal, reuse and toxicity of catalysts are major concerns, investigations have focused on the development of heterogeneous magnetic materials that can activate efficiently the oxidative degradation of the pollutants and at the same time minimize secondary contamination events. Taking into account these requirements, it is not surprising that the growing interest in ferrite NPs is due to the following reasons: (i) large surface area enhances the catalytic activity; (ii) the onset of superparamagnetism enables facile removal of the catalyst; (iii) versatility of ferrite compositions makes feasible the tuning of the optical band gap of the material enabling photo-degradation approaches and (iv) chemical stability of the ferrite structure avoids metal leaking to the environment.

Nanostructured CoFe2O4 is shown to be a promising material for heterogeneous peroximonosulfate (HSO5 �) activation in order to generate sulfate radicals (SO4 �• ) that promote the oxidative decomposition of organic pollutants like 2,4-dichlorophenol [167]. The convenience of this ferrite over other cobalt oxides stems from the fact that Fe(III) cations easily cause the hydrolysis of water, leading to surface Fe(III)-OH species that can be further converted into surface Co(II)-OH complexes. In turn, such complexes are the key for the reaction of HSO5 � to yield the SO4 �• radicals that promote pollutant decomposition. Besides, CoFe2O4 NPs present other advantages such as no Co(II) leaching and suitable magnetic properties for the easy recovery of the catalyst. CuFe2O4 NPs have also been employed for the HSO5 � activation in the catalytic degradation of atrazine [168] and tetrabromobisphenol [82]. In both cases, the HSO5 � decomposition was claimed to be triggered by the cycle Cu2+/Cu+ [168]. This is in disagreement with the report of Zhang et al. [169], which assigned the main role to the redox pair Cu2+/Cu3+. Other oxidation technologies with the aid of ferrite NPs entail the persulfate (S2O8 <sup>2</sup>�) [170] and H2O2 [12, 44, 81] heterogeneous activation for the decomposition of a wide range of organic pollutants.

An improvement in the catalytic properties of ferrites can be assessed by using composites with rGO [53, 54, 81] and MWCNs [83]. Such a synergic effect is attributed to the large surface area of the composites and to the electronic properties of these carbon-based functional materials. The proposed mechanism is outlined below [171]. It comprises the initial formation of the electron-hole pair in the ferrite phase by photon absorption (I), followed by the rapid electron transfer reaction from the ferrite conduction band to the rGO sheets (II). H2O2 is then decomposed in the vicinity of the rGO producing highly oxidative •OH radicals (III), which are also formed from the remaining holes in the ferrite (IV). As can be seen, step (II) is crucial for the efficient separation of photo-generated carriers, which is facilitated by the high electron conductivity of the conjugated π structure of the rGO sheets, which inhibits electron-hole recombination [172]. Moreover, •OH radicals are generated close to the rGO-adsorbed target organic pollutants, thus enhancing the decomposition rate.

MFe2O4 + hν ! MFe2O4 (h + e) (I)

$$\text{MFe}\_2\text{O}\_4\text{ (e)} + \text{rGO} \rightarrow \text{MFe}\_2\text{O}\_4 + \text{rGO} \text{ (e)}\tag{II}$$

6.2. Advanced oxidation technology

206 Magnetic Spinels- Synthesis, Properties and Applications

permission of Elsevier).

Chemical Society).

Advanced oxidation technologies consist of the assisted degradation of a given pollutant by using a source of highly oxidizing transient species. Such species are generally activated by the action of another substance that acts as a catalyst. Since the removal, reuse and toxicity of catalysts are major concerns, investigations have focused on the development of heterogeneous magnetic materials that can activate efficiently the oxidative degradation of the pollutants and at the same time minimize secondary contamination events. Taking into account these requirements, it is not surprising that the growing interest in ferrite NPs is due to the

Figure 10. Mechanism of reversible pH-responsive behavior (Adapted from Ref. [114] with permission of the American

Figure 9. Proposed mechanism for CR adsorption onto PEG-functionalized MFe2O4 NPs (Adapted from Ref. [75] with

$$\text{rGO } (\text{e}) + \text{H}\_2\text{O}\_2 \rightarrow \text{rGO} + \bullet\text{OH} + \bullet\text{OH}^- \tag{11}$$

$$\text{MFe}\_2\text{O}\_4\text{ (}h\text{)} + \text{OH}^- \rightarrow \text{MFe}\_2\text{O}\_4 + \text{\textdegree OH} \tag{1V}$$

Along the same lines, Fu et al. [172] noted that in the case of CoFe2O4/rGO nanocomposites with 40% of GO, there is no need for H2O2 to achieve efficient catalytic degradation of several dyes. Other materials for dye degradation involving ferrites are CoFe2O4/TiO2 nanocomposites [173].

## 7. Concluding remarks and perspectives


## Acknowledgements

The preparation of this chapter was partially supported by the CONACyT (Mexico) Projects 2013-05-231461, CB-2014-01-235840 and 2015-270810.

## Author details

7. Concluding remarks and perspectives

208 Magnetic Spinels- Synthesis, Properties and Applications

properties.

desired binding affinity.

paramagnetic behavior is required.

calculations could help in this regard.

2013-05-231461, CB-2014-01-235840 and 2015-270810.

lines of development.

Acknowledgements

ligands like sulfur groups, the tendency is inverted.

• Synthesis techniques for nanostructured spinel ferrites are available to tune their magnetic

• The nanoparticle surface is able to bind a wide variety of molecules with distinct functional groups that not only contribute to colloidal stabilization but also serve as the starting point for further conjugation steps. Many organic and inorganic reactions can be driven at the surface of ferrites, which allow for the tailoring of specific ligands with the

• The combination of these two advantages—tuning of magnetic properties and surface versatility—makes ferrites useful and promising materials for applications where super-

• Functionalized ferrite NPs, especially Fe3O4, are useful for removing a wide variety of heavy metals. In the case of cations, amino, carboxyl and thiol functional groups prevail as preferred candidates for metal uptake, although phosphonic and hydroxamic acids constitute promising ligands. Multifunctional ligands (synthetic and natural polymers) contribute to increase the stability and the adsorption capacity of the sorbent. At intermediate pH values, the tendency between metal-ligand affinities shows that for carboxyl and amino groups, the NPs are more selective toward hard Lewis acids, while for softer

• For removing arsenic, additional studies are warranted since controversy exists about the structure of the inner-sphere complexes and the nature of redox reactions at the interface. Also, the use of organic ligands to drive arsenic removal has not been exhausted yet. • Most adsorption studies are limited to thermodynamic and kinetic analysis and the investigations of metal-binding interactions are supported by phenomenological models. But the mode of coordination and the geometry of the surface complexes are not clear and so detailed spectroscopic studies are still needed. Since this is a tough task due to the inherent difficulties for the achievement of a rigorous surface picture, the use of theoretical

• Organic dyes are preferentially adsorbed by ligand-decorated magnetic NPs. Composites with functional carbonaceous materials and grafting of smart polymers are promising

• Spinel ferrites are useful materials for different advanced oxidation technologies, especially as composites with graphene-based materials due to the electronic and adsorptive properties of these carbon-based functional materials, which enhance the overall efficiency of the process.

The preparation of this chapter was partially supported by the CONACyT (Mexico) Projects

Oscar F. Odio1,2 and Edilso Reguera<sup>2</sup> \*

\*Address all correspondence to: edilso.reguera@gmail.com

1 Universidad de La Habana, Instituto de Ciencia y Tecnología de Materiales, La Habana, Cuba

2 Instituto Politécnico Nacional, Centro de Investigación en Ciencia Aplicada y Tecnología Avanzada-Unidad Legaria, México

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## **CVD‐Made Spinels: Synthesis, Characterization and Applications for Clean Energy CVD**‐**Made Spinels: Synthesis, Characterization and Applications for Clean Energy**

Patrick Mountapmbeme Kouotou, Guan‐Fu Pan and Zhen‐Yu Tian Patrick Mountapmbeme Kouotou, Guan‐Fu Pan and Zhen‐Yu Tian

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

http://dx.doi.org/10.5772/66285

#### **Abstract**

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> To reduce emissions and protect environment from pollution caused by volatile organic compounds (VOCs) and CO, catalytic oxidation can be applied as an efficient and promising technique. This review provides a novel and facile strategy to synthesize spinel-type and non-spinel-type transition metal oxides (TMOs). Specifically, single (Co3O4, α-Fe2O3, Mn3O4, CuO, Cu2O and Cr2O3) and binary (Co3-xCuxO4, Co3-xMnxO4 and Co3-xFexO4) TMOs have been prepared using pulsed spray evaporation chemical vapor deposition approach (PSE-CVD). PSE-CVD offers several advantages over conventional methods, such as relatively low cost, simplicity and high throughput, which makes it a promising strategy. Moreover, the PSE delivery system allows using less stable precursors and permits improving the reproducibility of the film properties with tailored compositions. The above listed TMOs prepared by PSE-CVD were successfully tested as catalysts toward the complete oxidation of some real fuels such as CO, C2H2, C3H6, *n*-C4H8 and C2H6O as representatives of VOCs and industrial exhaust streams. The active TMOs explored in this review could be potential catalysts candidates in one of the research areas that are currently under scrutiny, as the battle for the future of energy and environment involves the generation and application of clean energy.

> **Keywords:** clean energy, VOCs, biofuels, catalytic oxidation, PSE-CVD, TMOs, in-situ diagnostic; mechanism

© 2017 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. © 2017 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.

## **1. Introduction**

Volatile organic compounds (VOCs) are widely recognized as the major contributors to the global air pollution [1]. VOCs are composed of a variety of substances, which may be either natural or of anthropogenic origin from different human activities such as transportation and many factories or industrial processes including chemical, power and pharmaceutical plants, gas stations, petroleum refining, printing, food processing, automobile, as well as textile manufacturing [2]. The volatility of the emission from the above-listed sources enables them to diffuse more or less away from their place of issue, thus causing direct and indirect impacts on human health, animals and nature. In recent years, with the rapid increase of population, industrialization, transportation and urbanization, extremely severe and persistent haze pollution has been frequently observed in developing countries. For example, by 2020, VOCs emissions are predicted to increase by 49% relative to 2005 levels in China [3]. Therefore, in addition to the increasingly stringent controls for VOCs emissions level, it is urgent to develop and apply approaches to accelerate the reduction in VOCs emissions which is also vital in the context of climate change.

VOCs include a wide range of compounds such as aromatic and aliphatic hydrocarbons, alcohols, ketones, aldehydes, which are not easy to be oxidized. Abatement by catalytic oxidation appears to be a preferable technique compared to the thermal incineration in reducing VOCs. In fact, thermal oxidation is known to be expensive since it requires a substantial energy input to destroy dilute gas phase containing VOCs at relatively high temperature (750°C) which favors the formation of toxic by-products. In contrast, catalytic oxidation allows operating at much lower temperatures (200–500°C) and leads to none or negligible NOx formation in the combustion chamber. More importantly, catalytic oxidation can destroy VOCs and convert them into harmless CO2 and H2O [4]. In addition, the selectivity of catalytic oxidation could be well controlled. However, to achieve deep oxidation at mild temperature, highly active, nonselective and stable catalysts for extended periods of time are required.

Nowadays, the selection of catalyst for various organic pollutants abatement has been the subject of many studies, although the optimization of catalyst formulation does not appear to be an easy task. Noble metals and transition metal oxides (TMOs) have been widely explored in most commercial applications [5–11]. Noble metals are very active at low temperature, but their use is limited due to the high price, low thermal stability and tendency to poisoning [12]. In contrary, TMOs are considered as suitable alternatives because of higher thermal stability and lower price [13]. Among TMOs, single and mixed oxides, such as manganese and cobalt oxides, perovskites, zirconia-based catalysts, have been claimed for their effectiveness in VOCs oxidation [14–19]. In particular, Co3O4-based catalysts, which have been studied several decades ago regarding the high activity for CO and VOCs oxidation [20–22], have received again considerable attention in the recent years [23–28]. However, the physico-chemical and catalytic properties of TMOs thin films can be modulated with respect to the morphology, surface and bulk composition as well as metallic ratio, which are strongly dependent upon the preparation approaches and experimental conditions [29–32]. Thus, suitable synthesis route for the deposition of thin films of high purity and crystallinity is urgently needed.

In recent years, great efforts have been made to the development of efficient TMOs synthesis methods, including sol-gel [33], thermal decomposition [34], hydrothermal synthesis, electrodeposition [35] and pulsed spray evaporation chemical vapor deposition (PSE-CVD) [23–28]. Among these techniques, PSE-CVD shows the benefit of being a reliable one-pot method for the growth of complex oxides with controlled composition, since most of the functional oxides contain more than two elements in their structures and further tuning of their properties requires controlled doping ratio. The strategy of using multiple precursors in a single liquid feedstock and its combination with PSE-CVD have been proved to be a rationally controllable route for the growth of functionally mixed oxides such as spinels and perovskites [36]. In addition to complex oxides structures, PSE-CVD synthesis route offers also the potential to produce nano-scale layers of pure metals, metal carbides as well as alloys, which presents a large variety of potential applications.

In this review, we mainly focus on the progress made in the deposition of single and binary metal oxides-containing thin films using gas-phase processes namely PSE-CVD for complete catalytic oxidation of CO and VOCs operating at low temperature generally below 500°C or even at much lower temperature. Following a general introduction, a brief recall of the mechanisms involved in the catalytic oxidation over TMOs is described. The main sections deal with catalytic oxidation of VOCs over single and mixed TMOs followed by remarks and perspectives. We examined several typical metal oxides that are widely studied as the essential components for catalytic oxidation of CO and VOCs and explored the effect of some important influencing factors such as the redox properties, composition, doping, film morphology and the particle size of the metals oxides. The specific mechanisms involved in the catalytic activity process toward low-temperature VOC oxidation are discussed.

## **2. Reaction mechanisms with TMOs**

**1. Introduction**

218 Magnetic Spinels- Synthesis, Properties and Applications

context of climate change.

required.

Volatile organic compounds (VOCs) are widely recognized as the major contributors to the global air pollution [1]. VOCs are composed of a variety of substances, which may be either natural or of anthropogenic origin from different human activities such as transportation and many factories or industrial processes including chemical, power and pharmaceutical plants, gas stations, petroleum refining, printing, food processing, automobile, as well as textile manufacturing [2]. The volatility of the emission from the above-listed sources enables them to diffuse more or less away from their place of issue, thus causing direct and indirect impacts on human health, animals and nature. In recent years, with the rapid increase of population, industrialization, transportation and urbanization, extremely severe and persistent haze pollution has been frequently observed in developing countries. For example, by 2020, VOCs emissions are predicted to increase by 49% relative to 2005 levels in China [3]. Therefore, in addition to the increasingly stringent controls for VOCs emissions level, it is urgent to develop and apply approaches to accelerate the reduction in VOCs emissions which is also vital in the

VOCs include a wide range of compounds such as aromatic and aliphatic hydrocarbons, alcohols, ketones, aldehydes, which are not easy to be oxidized. Abatement by catalytic oxidation appears to be a preferable technique compared to the thermal incineration in reducing VOCs. In fact, thermal oxidation is known to be expensive since it requires a substantial energy input to destroy dilute gas phase containing VOCs at relatively high temperature (750°C) which favors the formation of toxic by-products. In contrast, catalytic oxidation allows operating at much lower temperatures (200–500°C) and leads to none or negligible NOx formation in the combustion chamber. More importantly, catalytic oxidation can destroy VOCs and convert them into harmless CO2 and H2O [4]. In addition, the selectivity of catalytic oxidation could be well controlled. However, to achieve deep oxidation at mild temperature, highly active, nonselective and stable catalysts for extended periods of time are

Nowadays, the selection of catalyst for various organic pollutants abatement has been the subject of many studies, although the optimization of catalyst formulation does not appear to be an easy task. Noble metals and transition metal oxides (TMOs) have been widely explored in most commercial applications [5–11]. Noble metals are very active at low temperature, but their use is limited due to the high price, low thermal stability and tendency to poisoning [12]. In contrary, TMOs are considered as suitable alternatives because of higher thermal stability and lower price [13]. Among TMOs, single and mixed oxides, such as manganese and cobalt oxides, perovskites, zirconia-based catalysts, have been claimed for their effectiveness in VOCs oxidation [14–19]. In particular, Co3O4-based catalysts, which have been studied several decades ago regarding the high activity for CO and VOCs oxidation [20–22], have received again considerable attention in the recent years [23–28]. However, the physico-chemical and catalytic properties of TMOs thin films can be modulated with respect to the morphology, surface and bulk composition as well as metallic ratio, which are strongly dependent upon the

Oxides-type catalysts made of transition metals are well known to selectively catalyze a large number of chemical processes. Most often, these oxides are used in the form of powder or supported thin film for oxidation reactions in the chemical industry or in automotive emission control. Although low-temperature catalytic oxidation of CO was intensively studied and the mechanism has been well addressed, it is still difficult to extend the results obtained from this reaction to catalytic oxidation of VOCs due to the different properties of pollutants and reaction conditions [37]. Depending on the partial reaction order, different reaction mechanisms have been proposed for CO and VOCs oxidation. First, the Langmuir mechanism states that the reaction occurs via the so-called Eley-Rideal (ER) process in which the reaction proceeds via a collision between an impinging gas-phase molecule and an adsorbed species, the controlling step being the reaction between an adsorbed molecule and a molecule from the gas phase [38]. The ER process for simple molecule such as CO can be schematically represented in **Scheme 1**.

**Scheme 1.** Schematic illustration of the ER process for CO oxidation.

Second, the Langmuir-Hinshelwood (LH) mechanism indicates the reaction happens through interactions among the adsorbed molecules, radicals or fragments of the reactant molecules [39] (see **Scheme 2**). According to LH, the controlling step is the surface reaction between two adsorbed molecules on analogous active sites.

**Scheme 2.** Schematic illustration of the LH process for CO oxidation.

Finally, the Mars-van Krevelen (MvK) mechanism [40] points out the lattice oxygen enters the reaction sequences, caused an oxidation-reduction sequence in the reaction of the reactant molecules and oxygen on different redox sites as schematically represented in **Scheme 3**. This mechanism has been widely accepted and used for the oxidation of a series of organic compounds.

**Scheme 3.** Illustration of the MvK mechanism for CO oxidation.

For the transition metal catalysts either in single oxide (Co3O4, CuO/Cu2O, Fe2O3, Cr2O3 as well as Mn3O4) or in binary oxide phase (Co-Cu, Co-Fe, Co-Mn and so on), it has been widely demonstrated and recognized that the determinant factors affecting their activities and performances toward complete oxidation of CO and VOCs are the formation of highly active oxygen species activated by the oxygen vacancies in addition to the close relationship between redox properties, oxygen vacancies and the bulk oxygen mobility through the Mars-van Krevelen (MvK) mechanism. In general, the MvK mechanism can proceed in two successive steps in terms of the cyclic reaction, namely the transfer of the bulk oxygen ions to the surface sites being the first step for the decomposition of bulk oxides and then the recovering of the bulk oxide. In this first step, oxygen vacancies on the catalyst surface are reduced as they react with the organic molecules. Therefore, the presence of surface oxygen vacancies plays a crucial role in the decomposition of bulk oxide. In the second step, the preformed reduced site is immediately regenerated through the consumption of gaseous oxygen or the transfer of oxygen atoms from the bulk to the surface, that is, the oxygen molecules compete with the bulk lattice oxygen for the surface oxygen vacancies, resulting in the inhibition of the bulk oxide decomposition process [41]. Since the catalyst is reduced in the first step and then reoxidized in the second step, this mechanism is also known as the redox mechanism.

## **3. Tailored synthesis, characterization and application of single and binary oxides**

The most active oxides frequently used are made of Ag, V, Cr, Mn, Fe, Co, Ni and Cu [42]. In fact, these *n‐*type and *p‐*type metal oxides are generally active catalysts (particularly *p‐*type) for deep oxidation since they are electron-deficient in the lattice and conduct electrons by means of positive "holes" [37]. In addition, the adsorbed oxygen species generally observed at their surfaces might participate in the reaction sequences together with the lattice oxide ions to improve the catalysts' performances. Therefore, numerous investigations on the development of TMOs catalysts are mainly focused on this type of oxides. Several *n‐*type and *p‐*type single oxides such as Co3O4 [23], *α*-Fe2O3, Mn3O4 [43], CuO [44], Cu2O [45] and Cr2O3 [46] have been prepared via PSE-CVD and successfully tested as active catalysts toward total oxidation of some real fuels such as CO, C2H2, C3H6, *n*-C4H8 and C2H6O as representatives of industrial exhaust stream or VOCs. Among these oxides mentioned above, a few of them seem to be particularly promising.

#### **3.1. CO and VOCs oxidation over PSE‐CVD made single oxide catalysts**

#### *3.1.1. Cobalt oxide with spinel structure*

**Scheme 1.** Schematic illustration of the ER process for CO oxidation.

**Scheme 2.** Schematic illustration of the LH process for CO oxidation.

**Scheme 3.** Illustration of the MvK mechanism for CO oxidation.

compounds.

adsorbed molecules on analogous active sites.

220 Magnetic Spinels- Synthesis, Properties and Applications

Second, the Langmuir-Hinshelwood (LH) mechanism indicates the reaction happens through interactions among the adsorbed molecules, radicals or fragments of the reactant molecules [39] (see **Scheme 2**). According to LH, the controlling step is the surface reaction between two

Finally, the Mars-van Krevelen (MvK) mechanism [40] points out the lattice oxygen enters the reaction sequences, caused an oxidation-reduction sequence in the reaction of the reactant molecules and oxygen on different redox sites as schematically represented in **Scheme 3**. This mechanism has been widely accepted and used for the oxidation of a series of organic

For the transition metal catalysts either in single oxide (Co3O4, CuO/Cu2O, Fe2O3, Cr2O3 as well as Mn3O4) or in binary oxide phase (Co-Cu, Co-Fe, Co-Mn and so on), it has been widely Cobalt oxide has a spinel structure and its formula can be written as CoO-Co2O3, or Co3O4. It presents an ideal spinel structure in which Co2+ cations occupy one-eighth of the tetrahedral sites, Co3+ cations occupy half of the octahedral sites and 32 sites are occupied by O2- ions [47]. Spinel Co3O4 has been used broadly and successfully for the destruction of CO and hydrocarbons compounds [48] and is claimed to be one of the most active catalysts in destruction of these compounds. The performances of some Co-based catalysts seem to be comparable to some noble metal catalysts. The high activity of Co3O4 is likely to be related to the relatively low enthalpy of vaporization (ΔHvap) of O2 [47]. Therefore, the Co–O bond strength of Co3O4 can affect desorption of lattice oxygen [49]. For example, CO frequently seems to react with pre-adsorbed or lattice oxygen to give CO2, which may further react to form surface carbonate species.

Kouotou et al. have studied the total oxidation of C3H6 and CO over PSE-CVD made spinel Co3O4 deposited on stainless steel grid mesh [23] and the as-deposited Co3O4 catalysts exhibited good activity compared with the reaction over non-coated mesh (NCM) as blank experiment (see **Figure 1**). The total conversion of the investigated compounds to CO2 was obtained respectively at around 380°C for C3H6 and 350°C for CO, which gives an obvious temperature shift relative to the NCM. This result shows that Co3O4 prepared by PSE-CVD was active for total oxidation of CO and C3H6. The catalytic performance of Co3O4 toward the oxidation of CO and C3H6 was attributed to the abundance of active Co3+/Co2+ cations and oxygen vacancies generally present at the surface of such materials. Co3+/Co2+ cations and oxygen vacancies were suggested to act as active sites for the oxidation process and are key parameters governing catalytic process during the total conversion rate of CO and C3H6.

**Figure 1.** Light-off curves of C3H6 and CO catalytic conversion over grid-mesh of stainless steel coated with Co3O4 and non-coated mesh as reference. Reproduced from [23] with permission. Copyright 2013, the Royal Society of Chemistry.

#### *3.1.2. Hematite*

Hematite, known as iron oxide (α-Fe2O3), has been extensively studied because of its excellent chemical stability, natural abundance, low cost, relatively nontoxic and environmentally benign [50]. As an important precursor, α-Fe2O3 is the most stable iron oxide phase, featuring some unique properties such as *n*-type semiconductor and magnetic as well as corrosionresistant properties [51]. *α‐*Fe2O3 can be converted into other functional materials such as maghemite (*γ‐*Fe2O3) and magnetite (Fe3O4) [52]. These properties have driven *α‐*Fe2O3 in numerous promising applications. Among them, application of α-Fe2O3 as catalysts caught wide attention. A systematic study was carried out by Walker et al. [53] who evaluated the possible application of iron catalysts for automotive emission control. Despite the promising results, *α‐*Fe2O3 has been scarcely tested as catalyst for the abatement of CO and VOCs. Recently, α-Fe2O3 thin films have been selectively prepared by PSE-CVD approach and successfully tested against the low-temperature catalytic oxidation of CO and propene. As a reference, the reaction was first performed on NCM in the temperature range of 180–950°C. An identical mesh coated with 20 mg *α‐*Fe2O3 was then tested under the same inlet gas conditions, but raising the temperature only to 500°C, which is the lattice stability temperature limit of *α‐*Fe2O3 [54]. In the presence of *α‐*Fe2O3, the conversion of CO started at ∼220°C and complete conversion occurred at 398°C, while these values were shifted to ∼300 and 930°C on NCM, respectively (see **Figure 2**). It should be mentioned that the reusability of the obtained α-Fe2O3 films and reproducibility of their catalytic performance were satisfactory within experimental uncertainty. The obtained results were compared to results reported by Walker et al. [53] who used unsupported Fe2O3 and several supported catalysts for CO oxidation. The temperature at 50% of CO conversion (defined as T50) was found to be 398°C with Fe2O3 [53], while PSE-CVD made *α‐*Fe2O3 thin films [55] enabled 50% conversion of CO to CO2 at 320°C, 78°C lower than Fe2O3 from Walker et al.'s work, revealing the better catalytic performance of the PSE-CVD deposited thin films. More details regarding the comparison can be found in **Table 2** (see Section 2.1.6).

some noble metal catalysts. The high activity of Co3O4 is likely to be related to the relatively low enthalpy of vaporization (ΔHvap) of O2 [47]. Therefore, the Co–O bond strength of Co3O4 can affect desorption of lattice oxygen [49]. For example, CO frequently seems to react with pre-adsorbed or lattice oxygen to give CO2, which may further react to form surface carbonate

Kouotou et al. have studied the total oxidation of C3H6 and CO over PSE-CVD made spinel Co3O4 deposited on stainless steel grid mesh [23] and the as-deposited Co3O4 catalysts exhibited good activity compared with the reaction over non-coated mesh (NCM) as blank experiment (see **Figure 1**). The total conversion of the investigated compounds to CO2 was obtained respectively at around 380°C for C3H6 and 350°C for CO, which gives an obvious temperature shift relative to the NCM. This result shows that Co3O4 prepared by PSE-CVD was active for total oxidation of CO and C3H6. The catalytic performance of Co3O4 toward the oxidation of CO and C3H6 was attributed to the abundance of active Co3+/Co2+ cations and oxygen vacancies generally present at the surface of such materials. Co3+/Co2+ cations and oxygen vacancies were suggested to act as active sites for the oxidation process and are key parameters governing catalytic process during the total conversion rate of CO and C3H6.

**Figure 1.** Light-off curves of C3H6 and CO catalytic conversion over grid-mesh of stainless steel coated with Co3O4 and non-coated mesh as reference. Reproduced from [23] with permission. Copyright 2013, the Royal Society of Chemistry.

Hematite, known as iron oxide (α-Fe2O3), has been extensively studied because of its excellent chemical stability, natural abundance, low cost, relatively nontoxic and environmentally benign [50]. As an important precursor, α-Fe2O3 is the most stable iron oxide phase, featuring some unique properties such as *n*-type semiconductor and magnetic as well as corrosion-

species.

222 Magnetic Spinels- Synthesis, Properties and Applications

*3.1.2. Hematite*

**Figure 2.** CO conversion profiles over stainless steel grid meshes coated with α-Fe2O3 films at a WHSV of 45,000 gcat-1 h-1. Two samples (S1 and S2) prepared at the same conditions were tested. First run and 2nd run are the catalytic tests performed for the first and second time, respectively. The performance for NCM is included as a reference. Reproduced from [55] with permission. Copyright 2013, Elsevier.

The reaction mechanism involved in CO oxidation over various Fe2O3 surfaces has been widely discussed in the literature [56–58]. As an example, a suprafacial mechanism describing the oxidation of CO over hematite has been proposed by Kandalam et al. [57] and was suggested to be applicable to the catalytic reaction with PSE-CVD made hematite [55]. According to Kandalam et al. [57], CO molecule first adsorbs onto hematite, weakening a Fe-O bond near the crystal surface and then, a second CO molecule adsorbs and forms CO2 by breaking the weakened Fe–O bond [57]. As hematite films are composed of a bulk and surface region on which active adsorption sites exist, part of these sites could be occupied by trapped oxygen atoms, which can originate either via dissociation of adsorbed oxygen during the oxidation reaction or via diffusion from the lattice to the surface. In addition, the number of surface iron atoms with which the adsorbed O2 can interact was earlier reported to play a significant role [59].

Besides CO oxidation, deep oxidations of C3H6 have been achieved at low temperature over a series of *α‐*Fe2O3 thin films coated at different temperature on stainless steel. The effect of the deposition temperature on the film morphology and redox properties has been systematically investigated and their influences on the catalytic oxidation of C3H6 have been clearly demonstrated [54]. Since the samples presented different redox behavior that is the temperatureprogrammed reduction (TPR), oxygen states distribution (lattice and adsorbed) obtained by the X-ray photoelectron spectroscopy (XPS) and morphology by helium ion microscopy (HIM, see **Figure 3**) [54], the catalytic tests were performed for all samples. The objective was to investigate these effects on the catalytic properties. To analyze the catalytic behavior after reoxidation as a prerequisite for application in consecutive cycles, the catalytic conversion of C3H6 was carried out both before and after pretreatment of the catalyst under oxygen flow. This was done because of the large amount of adventitious carbon, C–O and O–C=O type moieties detected at the surface of the samples prepared at 400 and 450°C (see **Table 1** and **Figure 3**) would be expected to limit their performance by forming a barrier layer between the active sites of the catalysts and the reactant molecules. The catalytic tests on fresh samples were performed over catalysts prepared at 350 and 450°C which present the lowest and the highest concentration of carbonaceous species, respectively. In addition, all pretreated samples were used at least twice in the catalytic tests to assess the reproducibility, as shown in **Figure 4**. **Figure 4** displays the light-off curves of C3H6 conversion obtained with α-Fe2O3 fresh samples (**Figure 4a**) and with α-Fe2O3 pretreated samples (**Figure 4b**). With fresh samples, catalyst (CAT) prepared at 350°C (CAT350) with lower adsorbed species presents better performance than CAT450°C. For each sample, the conversion begins at around 250°C. Temperatures for 10, 50 and 90% C3H6 conversion were presented for fresh and pretreated samples. The film prepared at 350°C was the most active one regarding the conversion profile as a function of temperature, followed by the films obtained at 400 and 450°C, respectively (**Table 1**). The experiments were repeated and the results were reproducible (**Figure 4b**). α-Fe2O3 thin films present competitive activity to that reported for supported noble metals. As an example, with only 20 mg of α-Fe2O3 deposited at 350°C, the oxidation of 50% of propene was reached at 331°C, whereas that of 200 mg Au/Al2O3 and La1.7Sr0.3CuO4S0.2 was obtained at 365 and 419°C, respectively. More details regarding the comparison can be found in **Table 2** (see Section 2.1.6). CVD‐Made Spinels: Synthesis, Characterization and Applications for Clean Energy http://dx.doi.org/10.5772/66285 225

The reaction mechanism involved in CO oxidation over various Fe2O3 surfaces has been widely discussed in the literature [56–58]. As an example, a suprafacial mechanism describing the oxidation of CO over hematite has been proposed by Kandalam et al. [57] and was suggested to be applicable to the catalytic reaction with PSE-CVD made hematite [55]. According to Kandalam et al. [57], CO molecule first adsorbs onto hematite, weakening a Fe-O bond near the crystal surface and then, a second CO molecule adsorbs and forms CO2 by breaking the weakened Fe–O bond [57]. As hematite films are composed of a bulk and surface region on which active adsorption sites exist, part of these sites could be occupied by trapped oxygen atoms, which can originate either via dissociation of adsorbed oxygen during the oxidation reaction or via diffusion from the lattice to the surface. In addition, the number of surface iron atoms with which the adsorbed O2 can interact was earlier reported

Besides CO oxidation, deep oxidations of C3H6 have been achieved at low temperature over a series of *α‐*Fe2O3 thin films coated at different temperature on stainless steel. The effect of the deposition temperature on the film morphology and redox properties has been systematically investigated and their influences on the catalytic oxidation of C3H6 have been clearly demonstrated [54]. Since the samples presented different redox behavior that is the temperatureprogrammed reduction (TPR), oxygen states distribution (lattice and adsorbed) obtained by the X-ray photoelectron spectroscopy (XPS) and morphology by helium ion microscopy (HIM, see **Figure 3**) [54], the catalytic tests were performed for all samples. The objective was to investigate these effects on the catalytic properties. To analyze the catalytic behavior after reoxidation as a prerequisite for application in consecutive cycles, the catalytic conversion of C3H6 was carried out both before and after pretreatment of the catalyst under oxygen flow. This was done because of the large amount of adventitious carbon, C–O and O–C=O type moieties detected at the surface of the samples prepared at 400 and 450°C (see **Table 1** and **Figure 3**) would be expected to limit their performance by forming a barrier layer between the active sites of the catalysts and the reactant molecules. The catalytic tests on fresh samples were performed over catalysts prepared at 350 and 450°C which present the lowest and the highest concentration of carbonaceous species, respectively. In addition, all pretreated samples were used at least twice in the catalytic tests to assess the reproducibility, as shown in **Figure 4**. **Figure 4** displays the light-off curves of C3H6 conversion obtained with α-Fe2O3 fresh samples (**Figure 4a**) and with α-Fe2O3 pretreated samples (**Figure 4b**). With fresh samples, catalyst (CAT) prepared at 350°C (CAT350) with lower adsorbed species presents better performance than CAT450°C. For each sample, the conversion begins at around 250°C. Temperatures for 10, 50 and 90% C3H6 conversion were presented for fresh and pretreated samples. The film prepared at 350°C was the most active one regarding the conversion profile as a function of temperature, followed by the films obtained at 400 and 450°C, respectively (**Table 1**). The experiments were repeated and the results were reproducible (**Figure 4b**). α-Fe2O3 thin films present competitive activity to that reported for supported noble metals. As an example, with only 20 mg of α-Fe2O3 deposited at 350°C, the oxidation of 50% of propene was reached at 331°C, whereas that of 200 mg Au/Al2O3 and La1.7Sr0.3CuO4S0.2 was obtained at 365 and 419°C, respectively. More details regarding the comparison can be found in **Table 2** (see Section 2.1.6).

to play a significant role [59].

224 Magnetic Spinels- Synthesis, Properties and Applications

**Figure 3.** TPR profiles obtained for α-Fe2O3 thin films, showing variation of the reduction properties in function of deposition temperature (A); XPS O 1s core-shell for α-Fe2O3 thin films, thin and thick lines are fitted and experimental results respectively (B) and HIM images of α-Fe2O3 thin films coated on stainless steel at different temperatures (C). Reproduced from [54] with permission. Copyright 2013, the Royal Society of Chemistry.

**Figure 4.** Light-off-curves of C3H6 conversion over a series of α-Fe2O3 coated on the mesh of stainless steel at different temperature and NCM. CAT350, CAT400 and CAT450 represent α-Fe2O3 deposited at 350, 400 and 450°C, respectively. The total flow rate was kept at 15 ml min-1 with 1% of C3H6 and 10% of O2 diluted in Ar at the WHSV of 45,000 gcat-1 h-1. Reproduced with permission from [54]. Copyright 2013, the Royal Society of Chemistry.


**Table 1.** XPS peak deconvolution result in percentage Oabs/OL ratio, T10, T50, T90 of α-Fe2O3 at different deposition temperature.

To identify the phenomenon governing the difference in activity between the three catalysts, a correlation between the catalytic behavior and the α-Fe2O3 characterization results was made [54]. XPS results revealed the presence of both adsorbed and lattice oxygen (**Figure 3B**). The Oadsorbed/Olattice ratio of CAT350 was the lowest among the three samples [54]. The catalytic tests over fresh and pretreated samples reveal the negligible role of the adsorbed oxygen, suggesting that the lattice oxygen plays a key role in the reaction sequence. This was corroborated by the TPR experiments, as shown in **Figure 3A**. In addition, it has been observed by Xie et al. [60] and Wang et al. [61] that the morphology or the crystal plane of metal oxides nano-crystal can remarkably alter their catalytic performance. As revealed by the HIM analysis, the film morphology of PSE-CVD made hematite was found to be significantly dependent on the deposition temperature (see **Figure 3C**). The increase of the preparation temperature leads to a variation of film morphology, which is responsible of the increase of the grain size. The grain size is one of the most crucial factors that determine the catalytic performances of catalysts. It is known that the smaller the grain size, the larger the specific surface area. Therefore, it was assumed that CAT350 with the smallest grain size and fine crystal structure possesses the largest specific surface area. This argument was consistent with the experimental observation that CAT350 presents the best performance.

Based on the TPR, XPS and HIM results, the possible oxidation mechanism of C3H6 over the as-prepared catalysts was suggested to follow a Mars-van Krevelen (MvK) mechanism, an intrafacial mechanism which involves migration of bulk oxygen to the surface, where it participates in the reaction with the reactant and replacement of bulk oxygen by oxygen from the gas phase [62]. Previous investigations have pointed out that this mechanism is valid in the combustion of hydrocarbons over transition metal oxide catalysts [63]. Also, catalytic combustion of C3H6 over Co3O4 has been reported by Liotta et al. [10] to proceed with the MvK mechanism.

#### *3.1.3. Manganese oxides*

Considering the low toxicity and availability, manganese oxides have attracted great attention among the various different transition metal oxides. Besides its unique physic-chemical properties, high activity and durability, they were extensively studied as catalysts [64– 66]. Manganese oxides possess a wide range of crystal phases (*β‐*MnO2, *γ‐*MnO2, *α‐*Mn2O3, *γ‐*Mn2O3, *α‐*Mn3O4 and Mn5O8) as well as variable oxidations states (+II, +III, +IV), which confer strong ability to switch from one oxidation state to another one and enable the formation of defects in the lattice, beneficial to the high oxygen mobility and oxygen storage [67]. Catalytic oxidation of VOCs (benzene and toluene) was investigated over a series of manganese oxide catalysts (Mn3O4, Mn2O3 and MnO2). The sequence of catalytic activity was found as follows: Mn3O4 > Mn2O3 > MnO2, which was closely correlated with the oxygen mobility on the catalyst [68]. Following the same logic, an investigation of the role of lattice oxygen in catalytic activity of manganese oxides toward the oxidation of ethanol, ethyl acetate and toluene has been performed by Santos et al. [69]. The results indicate that Mn3O4 improves catalytic performance due to the increased reactivity and mobility of lattice oxygen. In a recent paper, Tian et al. [43] have reported that Mn3O4 is a highly stable and active catalyst in many respects comparable to conventional catalysts based on noble metals. The catalytic performance was investigated with respect to the total oxidation of CO and C3H6 at atmospheric pressure. The full comparison of the reactants and products is given in **Figure 5**. The oxidation of CO over Mn3O4 (**Figure 5a1**, **a3**) becomes observable at 190°C and complete conversion occurs at 343°C with a temperatures shift to 280 and 820°C for an experiment without Mn3O4 (**Figure 5b1**, **b3**). Compared to the CO oxidation over manganese oxides (MnOx and Mn2O3) prepared by precipitation [70], *T25* was observed at 298°C, which is higher than the value (250°C) obtained with PSE-CVD Mn3O4 [43]. Moreover, the temperature at 50% (*T50*) of CO conversion was observed at 271°C, demonstrating that Mn3O4 prepared by PSE-CVD is highly active in the deep oxidation of CO.

**OL (%) OH‐**

226 Magnetic Spinels- Synthesis, Properties and Applications

temperature.

mechanism.

*3.1.3. Manganese oxides*

 **(%) O–C=O, C=O (%) H2O (%) Oads/OL T10 (°C) T50 (°C) T90 (°C)**

CAT350 46.52 25.87 22.76 – 1.05 264 260 326 315 405 350 CAT400 31.19 23.69 41.20 3.92 2.21 – 295 – 355 – 400 CAT450 27.31 25.47 43.82 3.41 2.66 326 302 405 380 470 435

**Table 1.** XPS peak deconvolution result in percentage Oabs/OL ratio, T10, T50, T90 of α-Fe2O3 at different deposition

To identify the phenomenon governing the difference in activity between the three catalysts, a correlation between the catalytic behavior and the α-Fe2O3 characterization results was made [54]. XPS results revealed the presence of both adsorbed and lattice oxygen (**Figure 3B**). The Oadsorbed/Olattice ratio of CAT350 was the lowest among the three samples [54]. The catalytic tests over fresh and pretreated samples reveal the negligible role of the adsorbed oxygen, suggesting that the lattice oxygen plays a key role in the reaction sequence. This was corroborated by the TPR experiments, as shown in **Figure 3A**. In addition, it has been observed by Xie et al. [60] and Wang et al. [61] that the morphology or the crystal plane of metal oxides nano-crystal can remarkably alter their catalytic performance. As revealed by the HIM analysis, the film morphology of PSE-CVD made hematite was found to be significantly dependent on the deposition temperature (see **Figure 3C**). The increase of the preparation temperature leads to a variation of film morphology, which is responsible of the increase of the grain size. The grain size is one of the most crucial factors that determine the catalytic performances of catalysts. It is known that the smaller the grain size, the larger the specific surface area. Therefore, it was assumed that CAT350 with the smallest grain size and fine crystal structure possesses the largest specific surface area. This argument was consistent with the

Based on the TPR, XPS and HIM results, the possible oxidation mechanism of C3H6 over the as-prepared catalysts was suggested to follow a Mars-van Krevelen (MvK) mechanism, an intrafacial mechanism which involves migration of bulk oxygen to the surface, where it participates in the reaction with the reactant and replacement of bulk oxygen by oxygen from the gas phase [62]. Previous investigations have pointed out that this mechanism is valid in the combustion of hydrocarbons over transition metal oxide catalysts [63]. Also, catalytic combustion of C3H6 over Co3O4 has been reported by Liotta et al. [10] to proceed with the MvK

Considering the low toxicity and availability, manganese oxides have attracted great attention among the various different transition metal oxides. Besides its unique physic-chemical properties, high activity and durability, they were extensively studied as catalysts [64– 66]. Manganese oxides possess a wide range of crystal phases (*β‐*MnO2, *γ‐*MnO2, *α‐*Mn2O3,

Reproduced with permission from [54]. Copyright 2013, the Royal Society of Chemistry.

experimental observation that CAT350 presents the best performance.

**FS PS FS PS FS PS**

**Figure 5.** Production and light-off curves of CO and C3H6 oxidation over Mn3O4-coated mesh (a1, a2 and a3) and NCM (b1, b2 and b3), respectively. Light-off curve of CO (a3) and C3H6 (b3) oxidation over Mn3O4-coated mesh and NCM. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of C3H6 and 10% of O2 diluted in Ar corresponding the WHSV of 45,000 gcat-1 h-1. Reproduced with permission from [43]. Copyright 2013, American Chemical Society.

It was also observed that the conversion of C3H6 on Mn3O4 (**Figure 5, a3**) becomes detectable at 220°C and conversion of C3H6 approaches 100% at 433°C. However, the oxidation of C3H6 on NCM starts at ∼290°C and a temperature as high as 821°C (**Figure 5b2, b3**) was required for complete conversion of C3H6. More importantly, at low C3H6 conversions, the main product is CO2, but a trace of CO peaking at 769°C (**Figure 5b2**) was formed on NCM, which results from partial oxidation. However, except CO2 and water, no other by-products were detected in the presence of Mn3O4, indicating that Mn3O4 is effective to reduce CO emission.

The catalytic oxidation of CO and C3H6 over Mn3O4 was suggested to likely follow the MvK mechanism which involves reversible reduction/reoxidation steps of the lattice oxygen atoms. As reported in the above-mentioned text, trapped oxygen was apparently released with the increase of the temperature and Mn3O4 tends to be reduced through the subsurface oxidation of CO or C3H6 with the lattice or surface oxygen. Subsequently, the reduced metal oxide site is reoxidized by the atmospheric O2. According to the redox results [43], Mn3O4 was easier to be reoxidized than to be reduced, demonstrating that the reduction step could play a crucial role in the kinetics of the catalytic oxidation process. The abundance of adsorbed oxygen, revealed in the XPS results in addition to the excellent redox behavior [43], can play a key role in the oxidation process.

#### *3.1.4. Copper oxide*

Copper oxide is a promising semiconductor material which is widely exploited for a broad field of applications. As CuO is nontoxic and its constituents are abundantly available, the synthesis of CuO is considered as an important research topic for catalytic processes. Copper oxide is unique as it has square planar coordination of copper by oxygen in the monoclinic structure. TMOs, such as Co3O4 and Mn3O4, are effective catalysts for the oxidative destruction of VOCs [24]. However, not much is known concerning the abatement of CO and VOCs with CuO [71–73]. The selective oxidation of propene with O2 to propylene oxide and acrolein has been study by Hua et al. [74] and the results indicate that Cu2O nanocrystals control the catalytic selectivity as well as the activity in propylene oxidation with O2. In addition, the authors also reveal the underlying structure-activity relationships of this complex heterogeneous catalytic reaction at the molecular level and identify the catalytically active sites. Most recently, Tian et al. [44] and Pan et al. [45] have studied the low-temperature complete oxidation of VOCs of olefins type such as C2H2 and C3H6 over CuOx. They have demonstrated that CuO and Cu2O are active and promising catalyst for the abatement of VOCs.

The catalytic performance of the CuO was evaluated for the complete oxidation of C3H6 by Tian et al. [44]. The background effect of the mesh on the combustion process was examined by carrying out the oxidation of C3H6 on NCM under the same gas inlet conditions. **Figure 6** compares the conversion temperature of C3H6 production over CuO films and NCM. The conversion plots show clearly that CuO favors the total oxidation C3H6 (**Figure 6a**) at lower temperatures relative to NCM. In the presence of CuO, the consumption of C3H6 was observable at about 190°C and complete conversion was reached within 310°C, while these two values shift toward higher temperatures for the reaction on NCM. No trace of CO was detected in the oxidation process with CuO (**Figure 6b**). However, a significant amount of CO was formed in the reaction without catalyst, which was assigned to come from the partial oxidation reaction. In the two cases, CO2 was observed to be the final product. The temperatures *T10*, *T50,* and *T90*, corresponding to the 10, 50 and 90% C3H6 conversion during the temperature-programmed reaction, were selected to compare the catalytic performance of the deposited CuO as well as some representative catalysts available in the literature (see **Table 2** in Section 2.1.6) toward C3H6 oxidation.

CO2, but a trace of CO peaking at 769°C (**Figure 5b2**) was formed on NCM, which results from partial oxidation. However, except CO2 and water, no other by-products were detected in the

The catalytic oxidation of CO and C3H6 over Mn3O4 was suggested to likely follow the MvK mechanism which involves reversible reduction/reoxidation steps of the lattice oxygen atoms. As reported in the above-mentioned text, trapped oxygen was apparently released with the increase of the temperature and Mn3O4 tends to be reduced through the subsurface oxidation of CO or C3H6 with the lattice or surface oxygen. Subsequently, the reduced metal oxide site is reoxidized by the atmospheric O2. According to the redox results [43], Mn3O4 was easier to be reoxidized than to be reduced, demonstrating that the reduction step could play a crucial role in the kinetics of the catalytic oxidation process. The abundance of adsorbed oxygen, revealed in the XPS results in addition to the excellent redox behavior [43], can play a key role in the

Copper oxide is a promising semiconductor material which is widely exploited for a broad field of applications. As CuO is nontoxic and its constituents are abundantly available, the synthesis of CuO is considered as an important research topic for catalytic processes. Copper oxide is unique as it has square planar coordination of copper by oxygen in the monoclinic structure. TMOs, such as Co3O4 and Mn3O4, are effective catalysts for the oxidative destruction of VOCs [24]. However, not much is known concerning the abatement of CO and VOCs with CuO [71–73]. The selective oxidation of propene with O2 to propylene oxide and acrolein has been study by Hua et al. [74] and the results indicate that Cu2O nanocrystals control the catalytic selectivity as well as the activity in propylene oxidation with O2. In addition, the authors also reveal the underlying structure-activity relationships of this complex heterogeneous catalytic reaction at the molecular level and identify the catalytically active sites. Most recently, Tian et al. [44] and Pan et al. [45] have studied the low-temperature complete oxidation of VOCs of olefins type such as C2H2 and C3H6 over CuOx. They have demonstrated that CuO

The catalytic performance of the CuO was evaluated for the complete oxidation of C3H6 by Tian et al. [44]. The background effect of the mesh on the combustion process was examined by carrying out the oxidation of C3H6 on NCM under the same gas inlet conditions. **Figure 6** compares the conversion temperature of C3H6 production over CuO films and NCM. The conversion plots show clearly that CuO favors the total oxidation C3H6 (**Figure 6a**) at lower temperatures relative to NCM. In the presence of CuO, the consumption of C3H6 was observable at about 190°C and complete conversion was reached within 310°C, while these two values shift toward higher temperatures for the reaction on NCM. No trace of CO was detected in the oxidation process with CuO (**Figure 6b**). However, a significant amount of CO was formed in the reaction without catalyst, which was assigned to come from the partial oxidation reaction. In the two cases, CO2 was observed to be the final product. The temperatures *T10*, *T50,* and *T90*, corresponding to the 10, 50 and 90% C3H6 conversion during the temperature-programmed reaction, were selected to compare the catalytic performance of the deposited CuO as well

and Cu2O are active and promising catalyst for the abatement of VOCs.

presence of Mn3O4, indicating that Mn3O4 is effective to reduce CO emission.

oxidation process.

228 Magnetic Spinels- Synthesis, Properties and Applications

*3.1.4. Copper oxide*

**Figure 6.** Light‐off curves for C3H6 oxidation: conversion of C3H6 (a); associated CO production (b) and CO2 production (c) with CuO‐coated and NCM. The reaction was carried at the total flow rate was kept at 15 ml min‐1 with 1% of C3H6 and 10% of O2 diluted in Ar corresponding the WHSV of 45,000 gcat‐1  h‐1 . Reproduced with permission from [44]. Copy‐ right 2013, Elsevier.

Cu2O was also tested by Pan et al. [45] as catalyst for the deep oxidation C2H2 and C3H6. The catalytic tests were carried out three times for the same sample and the results were quite close, demonstrating that the prepared Cu2O has good reusability with reproduced results. **Figure 7A** compares the temperature‐dependent conversion ratio of C2H2 and C3H6 with Cu2O‐coated mesh and NCM. Compared to the NCM condition, the complete oxidation of C2H2 decreased from 450 to 300°C and for C3H6 decreased from 675 to 425°C over Cu2O. It should be mentioned that during the oxidation of C2H2 and C3H6 over Cu2O‐coated mesh, CO was not detected. However, CO was detected abundantly over NCM. Compared to oth‐ er TMOs such as Mn3O4 [43] and Co3O4 [23] for the oxidation of C2H2 and C3H6, Cu2O exhibits much better catalytic performances. With an Arrhenius expression, apparent activation energies (*Ea*) of <15% of C2H2 and C3H6 conversion were deduced. The *Ea* of C2H2 and C3H6 oxidation over NCM was 93.5 and 92.0 kJ mol-1, while these values shift to 51.7 and 57.0 kJ  mol-1 with Cu2O-coated samples (average values for three times), respectively. Compared to *Ea* obtained with other TMOs, such as Co3O4 (128.9 kJ mol-1 for C2H2, 127.1 kJ mol-1 for C3H6) [48] and Mn3O4 (84.7 kJ mol-1 for C3H6) [43], the reaction with Cu2O also shows lower *Ea*. Thus, the relatively low *Ea* was suggested to contribute in the acceleration of the oxidation processes and enhances the catalytic performance for the oxidation of C2H2 and C3H6. As redox mechanism is generally accepted to be the dominant mechanism in the oxidation of low-rank hydrocarbons over TMOs, both authors also suggested CuO and Cu2O follow a redox process. According to Pan et al. [45], Cu2O is suggested to react first with oxygen, giving rise to CuO. Second, the reaction of C2H2 and C3H6 with the trapped or lattice oxygen occurs, leading to CuO reduction and release of oxygen to form Cu2O. From the XPS results (**Figure 7B**), O 1s core shell shows mainly O2 2- and O species. Both O2 2- and O are known as strongly electrophilic reactants capable to attack an organic molecule in the region of its highest electron density and result in the oxidation of the carbon skeleton. As the electrophilic oxygen species such as lattice and adsorbed oxygen generally participates in the total oxidation of hydrocarbons to CO2, these electrophilic oxygen species (O2 2- or O- ) presented at the surface of Cu2O were suggested to be also benefit for the complete conversion of C2H2 and C3H6.

**Figure 7.** Outlet profiles of C2H2 A (a) and C3H6 A (b) oxidation over NCM and mesh grid of stainless steel coated with Cu2O. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of fuel and 10% O2 diluted in Ar, corresponding the WHSV of 45,000 gcat-1 h-1. Cu LM2 B (a) and O 1s signals B (b) of representative Cu2O thin film. Reproduced with permission from [45]. Copyright 2015, the Royal Society of Chemistry.

#### *3.1.5. Chromium oxide*

er TMOs such as Mn3O4 [43] and Co3O4 [23] for the oxidation of C2H2 and C3H6, Cu2O exhibits much better catalytic performances. With an Arrhenius expression, apparent activation energies (*Ea*) of <15% of C2H2 and C3H6 conversion were deduced. The *Ea* of C2H2 and C3H6 oxidation over NCM was 93.5 and 92.0 kJ mol-1, while these values shift to 51.7 and 57.0 kJ  mol-1 with Cu2O-coated samples (average values for three times), respectively. Compared to *Ea* obtained with other TMOs, such as Co3O4 (128.9 kJ mol-1 for C2H2, 127.1 kJ mol-1 for C3H6) [48] and Mn3O4 (84.7 kJ mol-1 for C3H6) [43], the reaction with Cu2O also shows lower *Ea*. Thus, the relatively low *Ea* was suggested to contribute in the acceleration of the oxidation processes and enhances the catalytic performance for the oxidation of C2H2 and C3H6. As redox mechanism is generally accepted to be the dominant mechanism in the oxidation of low-rank hydrocarbons over TMOs, both authors also suggested CuO and Cu2O follow a redox process. According to Pan et al. [45], Cu2O is suggested to react first with oxygen, giving rise to CuO. Second, the reaction of C2H2 and C3H6 with the trapped or lattice oxygen occurs, leading to CuO reduction and release of oxygen to form Cu2O. From the XPS results

2- and O-

strongly electrophilic reactants capable to attack an organic molecule in the region of its highest electron density and result in the oxidation of the carbon skeleton. As the electrophilic oxygen species such as lattice and adsorbed oxygen generally participates in the total

at the surface of Cu2O were suggested to be also benefit for the complete conversion of C2H2

**Figure 7.** Outlet profiles of C2H2 A (a) and C3H6 A (b) oxidation over NCM and mesh grid of stainless steel coated with Cu2O. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of fuel and 10% O2 diluted in Ar, corresponding the WHSV of 45,000 gcat-1 h-1. Cu LM2 B (a) and O 1s signals B (b) of representative Cu2O thin film. Re-

produced with permission from [45]. Copyright 2015, the Royal Society of Chemistry.

oxidation of hydrocarbons to CO2, these electrophilic oxygen species (O2

species. Both O2

2- and O-

2- or O-

are known as

) presented

(**Figure 7B**), O 1s core shell shows mainly O2

230 Magnetic Spinels- Synthesis, Properties and Applications

and C3H6.

Chromium oxide (Cr2O3) has been broadly and successfully used for numerous applications such as wear resistance, corrosion protection, optics and electronics due to its high melting point, heat resistance, mechanical strength, chemical inertness, optical characteristics, high hardness and low friction coefficient [75–79]. It has been proved to be active when used as catalyst for the destruction of halogenated compounds. In addition to its good thermal stability, Cr2O3 exhibited attractive performance in the low-temperature abatement of VOCs. Recently, Liang et al. [46] reported the synthesis of Cr2O3 by PSE-CVD for the total oxidation of propene. **Figure 8** compares the temperature-dependent conversion of C3H6 to CO2 as the final product and the associated CO production over Cr2O3 films and NCM. The conversion plots show clearly that Cr2O3 enables the oxidation of C3H6 at lower temperatures relative to the NCM. In the presence of Cr2O3, the consumption of C3H6 becomes observable at about 400°C and complete conversion occurs around 550°C, while these temperatures were observed to shift, respectively, to 600 and 775°C for the reaction on NCM. Compared to the *T50* of C3H6 conversion, *T50* obtained with Cr2O3 was quite close to the values reported for Au/Al2O3 and La1.7Sr0.3CuO4S0.2, which indicates that the PSE-CVD made Cr2O3 exhibits similar catalytic performance to the noble metal and perovskite. During the oxidation of C3H6 over Cr2O3-coated mesh, only a small amount of CO was detected. However, considerable quantities of CO (**Figure 8b**) were formed in the reaction without catalyst, which could originate from the partial oxidation reaction.

**Figure 8.** Outlet profiles of C3H6 oxidation over Cr2O3-coated and NCM. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of fuel and 10% O2 diluted in Ar, corresponding the WHSV of 62,500 gcat-1 h-1. Reproduced with permission from [46]. Copyright 2015, Wiley-VCH.

Diffuse reflectance infrared Fourier transformed spectroscopy (DRIFTS) was used to study the possible catalytic reaction mechanism of Cr2O3. A feed of argon gas containing 1% of C3H6 and 10% of O2 was introduced into the reactor which was equipped with a small piece of stainless steel grid mesh as the catalyst support. The IR spectra at different temperatures are presented in **Figure 9**. Surface adsorption species with IR bands at 729 and 913 cm-1 are clearly observed at temperature lower than 100°C, which were assigned to adsorption of C3H6 on the deposited film. These bands were still detectable below 300°C. When the temperature increases to 500°C, the coordinated C3H6 vanishes gradually and CO2 was detected. The temperature for the complete conversion of C3H6 to CO2 was around 500°C, in good agreement with the catalytic tests.

**Figure 9.** In-situ DRIFTS spectra of Cr2O3 during C3H6 adsorption at different temperatures. Reproduced with permission from [46]. Copyright 2015, Wiley-VCH.

#### *3.1.6. Comparison of the catalytic performances of single TMOs with literature*

**Table 2** summarizes and compares the catalytic performance of the PSE-CVD made single TMOs with that of the selected catalysts in the literature. Even though the experimental condition was not exactly the same, regarding the weight of the catalyst and the WHSV, our catalysts present comparable activity to the systems presented in the table. Special attention is paid to the comparison of the catalytic performance with noble metals and transition metal oxides. As it can be seen from **Table 2**, noble metals and TMOs (single) present attractive results in terms of CO and propene total oxidation.


CVD‐Made Spinels: Synthesis, Characterization and Applications for Clean Energy http://dx.doi.org/10.5772/66285 233


a : *Weight hourly space velocity*

Diffuse reflectance infrared Fourier transformed spectroscopy (DRIFTS) was used to study the possible catalytic reaction mechanism of Cr2O3. A feed of argon gas containing 1% of C3H6 and 10% of O2 was introduced into the reactor which was equipped with a small piece of stainless steel grid mesh as the catalyst support. The IR spectra at different temperatures are presented in **Figure 9**. Surface adsorption species with IR bands at 729 and 913 cm-1 are clearly observed at temperature lower than 100°C, which were assigned to adsorption of C3H6 on the deposited film. These bands were still detectable below 300°C. When the temperature increases to 500°C, the coordinated C3H6 vanishes gradually and CO2 was detected. The temperature for the complete conversion of C3H6 to CO2 was around 500°C, in good agreement with the catalytic

**Figure 9.** In-situ DRIFTS spectra of Cr2O3 during C3H6 adsorption at different temperatures. Reproduced with permis-

**Table 2** summarizes and compares the catalytic performance of the PSE-CVD made single TMOs with that of the selected catalysts in the literature. Even though the experimental condition was not exactly the same, regarding the weight of the catalyst and the WHSV, our catalysts present comparable activity to the systems presented in the table. Special attention is paid to the comparison of the catalytic performance with noble metals and transition metal oxides. As it can be seen from **Table 2**, noble metals and TMOs (single) present attractive results

Single TMOs α-Fe2O3 20 1% C3H6/10% O2 in Ar 45,000 260 313 350 [54]

α-Fe2O3 400°C 295 355 400 α-Fe2O3 450°C 302 380 435 Co3O4 12 1% C3H6/10% O2 in Ar 75,000 306 347 396 [23]

**(ml g‐1 h‐1)** 

**T10** *<sup>b</sup>* **(°C) T50** *<sup>b</sup>* **(°C)** **T90** *<sup>b</sup>* **(°C)**  **Refs.**

*3.1.6. Comparison of the catalytic performances of single TMOs with literature*

**Type of catalyst Material Weight (mg) Gas composition WHSV** *<sup>a</sup>*

tests.

sion from [46]. Copyright 2015, Wiley-VCH.

232 Magnetic Spinels- Synthesis, Properties and Applications

in terms of CO and propene total oxidation.

b : *Temperature at X% conversion (X=10, 50 and 90 %)*

**Table 2.** Overview of the catalytic performances of PSE-CVD single made TMOs compared with the literature.

#### **3.2. CO and VOCs oxidation over PSE‐CVD made binary TMOs catalysts**

Among the different mixed-oxide structures, numerous references can be found, dealing with the reactivity of perovskites such as compounds with A2BO4 structure or spinel-type mixed oxide (AB2O3) for the CO and VOCs oxidation reaction. Since the 1970s, some spinel-type mixed oxides are known to exhibit activity for CO and VOCs oxidation. In order to further improve the catalytic performance of TMOs, some combinations of oxides have been formulated and the obtained binary or mixed TMOs have exhibited better activities than the single or mixed components or even comparable with that of noble metals. Such catalysts include Co-Mn [24], Cu-Co [28] and Co-Fe [25–27]. It is important to note that, because of the weak performances of Cr2O3 catalyst compared to other single oxide presented in this review, in addition to highly toxicity, we have restricted the application of Cr-based catalysts to low operation temperatures.

#### *3.2.1. Catalytic oxidation of VOCs over mixed Co‐Mn oxides made by PSE‐CVD*

Tian et al. have studied the catalytic oxidation of VOCs over spinels Co3-xMnxO4 (0 ≤ x ≤ 0.34) binary oxides obtained by PSE-CVD. The grown Co-Mn binary oxides were tested toward the total oxidation of C2H2 and C3H6 as illustrative examples [24]. The TPR/TPO results and the light-off curves of the samples are presented in **Figures 10** and **11** and summarized in **Table 3**. **Figure 10A** compares the temperature-dependent conversion of C2H2 and C3H6 over Co3 xMnxO4 films and NCM. The detailed outlet profiles of the fuels and products are given in **Figure 11**. The temperatures *T50*, *T90*, *T50,* and *T90* of CO2, corresponding to the temperatures of CO2 production, were selected as parameters to indicate the catalytic activity of the deposited samples toward the deep oxidation of hydrocarbons, as shown in **Table 3**. The catalytic performances were improved from Co3O4 to Co2.66Mn0.34O4, both being superior to the NCM.

**Figure 10.** (A) Light-off curves of C2H2 and C3H6 over Co3-xMnxO4 (x = 0 and 0.34) spinel structures grown on mesh of stainless steel substrates and NCM. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 2% of fuel and 20% O2 diluted in Ar at a total flow rate of 0.015 L/min; (B) TPO/TPR patterns of Co-Mn oxide. Reproduced with permission from [45]. Copyright 2015, Elsevier.

CVD‐Made Spinels: Synthesis, Characterization and Applications for Clean Energy http://dx.doi.org/10.5772/66285 235

**Figure 11.** Production profiles of C2H2/C3H6, CO, CO2 and CH4 in the oxidation of C2H2 and C3H6 over Co3-xMnxO4 oxides (x = 0 and 0.34) grown on mesh of stainless steel substrates and NCM. Reproduced with permission from [24]. Copyright 2015, Elsevier.


Reproduced with permission from [24]. Copyright 2012, Elsevier.

**3.2. CO and VOCs oxidation over PSE‐CVD made binary TMOs catalysts**

*3.2.1. Catalytic oxidation of VOCs over mixed Co‐Mn oxides made by PSE‐CVD*

operation temperatures.

234 Magnetic Spinels- Synthesis, Properties and Applications

with permission from [45]. Copyright 2015, Elsevier.

Among the different mixed-oxide structures, numerous references can be found, dealing with the reactivity of perovskites such as compounds with A2BO4 structure or spinel-type mixed oxide (AB2O3) for the CO and VOCs oxidation reaction. Since the 1970s, some spinel-type mixed oxides are known to exhibit activity for CO and VOCs oxidation. In order to further improve the catalytic performance of TMOs, some combinations of oxides have been formulated and the obtained binary or mixed TMOs have exhibited better activities than the single or mixed components or even comparable with that of noble metals. Such catalysts include Co-Mn [24], Cu-Co [28] and Co-Fe [25–27]. It is important to note that, because of the weak performances of Cr2O3 catalyst compared to other single oxide presented in this review, in addition to highly toxicity, we have restricted the application of Cr-based catalysts to low

Tian et al. have studied the catalytic oxidation of VOCs over spinels Co3-xMnxO4 (0 ≤ x ≤ 0.34) binary oxides obtained by PSE-CVD. The grown Co-Mn binary oxides were tested toward the total oxidation of C2H2 and C3H6 as illustrative examples [24]. The TPR/TPO results and the light-off curves of the samples are presented in **Figures 10** and **11** and summarized in **Table 3**. **Figure 10A** compares the temperature-dependent conversion of C2H2 and C3H6 over Co3 xMnxO4 films and NCM. The detailed outlet profiles of the fuels and products are given in **Figure 11**. The temperatures *T50*, *T90*, *T50,* and *T90* of CO2, corresponding to the temperatures of CO2 production, were selected as parameters to indicate the catalytic activity of the deposited samples toward the deep oxidation of hydrocarbons, as shown in **Table 3**. The catalytic performances were improved from Co3O4 to Co2.66Mn0.34O4, both being superior to the NCM.

**Figure 10.** (A) Light-off curves of C2H2 and C3H6 over Co3-xMnxO4 (x = 0 and 0.34) spinel structures grown on mesh of stainless steel substrates and NCM. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 2% of fuel and 20% O2 diluted in Ar at a total flow rate of 0.015 L/min; (B) TPO/TPR patterns of Co-Mn oxide. Reproduced a T50 and T90 represent the temperatures at which the conversion of C2H2/C3H6 reaches 50 and 90%, respectively.

b T50 CO2 and T90 CO2 represent the temperatures at which 50 and 90% conversion to CO2 is reached.

c Ea is the apparent activation energy for the hydrocarbon activation at atmospheric pressure.

**Table 3.** Catalytic oxidation of C2H2/C3H6 over Co-Mn oxides catalysts deposited on grid mesh of stainless steel substrates and NCM.

The selectivity-conversion plots show clearly that the investigated coatings favor the total oxidation of both unsaturated hydrocarbons at lower temperatures relative to the NCM. For C2H2, *T50* and *T90* were observed at 355 and 386°C over NCM, whereas these values decreased by respective 56 and 51°C over cobalt oxide and further decreased by 17 and 41°C over Co2.66Mn0.34O4, which reveals that the catalyst with small amount of manganese content shows the highest activity (see **Figure 10** and **Table 3**). Tian et al. [48] reported that *T50* of C2H2 oxidation with Co3O4 deposited on monolithic cordierite support was 297°C, which was in good agreement with the current Co3O4 sample prepared on mesh. Note that the *T50* of C2H2 (282°C, **Table 3**) with Co2.66Mn0.34O4 [24] was quite close to the value (280°C) reported by Ivanova et al. [80] who used Au/Al2O3, revealing that Co2.66Mn0.34O4 prepared by CVD features a competitive activity. The low-temperature shift of the light-off curve was more pronounced for C3H6 since the *T90* difference between the NCM and Co2.66Mn0.34O4 is 336°C, demonstrating that Co2.66Mn0.34O4 was very active for the total oxidation of propene and supporting the conclusion of Liang et al. [81] who reported that the manganese insertion into the cobalt oxide spinel could enhance the catalytic activity of the oxidation of light olefins. Considering the higher activity, Mn-doped spinel of cobalt oxide could be a potential system for catalytic oxidation of hydrocarbons. As presented in **Table 3**, the values of *T50* and *T90* of hydrocarbons conversion and those of CO2 selectivity were found to be almost identical for Co3O4. However, differences were observed relative to NCM especially in the oxidation of C3H6, illustrating an occurrence of partial oxidation by generating products other than CO2 (see **Figure 11f**). With Co3O4 and Co2.66Mn0.34O4, CO2 was observed to be the only carbonaceous product and no secondary products were formed all over the entire oxidation process of C3H6, while CO and CH4, coming from incomplete oxidation, were detected in addition to the production of CO2 on NCM. Even though some by-products were formed at certain temperatures with NCM, all indicate that CO2 was the final product. Taking the temperature at which hydrocarbons were completely converted to CO2 on Co2.66Mn0.34O4 sample as a reference, Co3O4 and NCM show a respective increase of 51 and 96°C for C2H2 and 36 and 358°C for C3H6, which indicate that the use of cobalt together with manganese significantly affects the catalytic activity. By applying the Arrhenius equation in the conversion range within 15% [48], the *Ea* was calculated (see **Table 3**). In general, the manganese-doped cobalt oxide exhibits lower *Ea*. Compared to those activation energies obtained in the oxidation of C2H2 and C3H6 over Co3O4-coated monolith with a large flow rate (500 sccm) [48], the values obtained on Co3O4-coated mesh were about 10 and 30 kJ mol-1 larger, the difference of which may come from the comprehensive effect of the support and inlet condition. Zhang et al. [82] and Aguilera et al. [83] also reported that the modification of Co3O4 by MnOx could promote the preferential oxidation of CO, toluene and alcohols at lower temperatures. The catalytic oxidation of the two hydrocarbons employing Co-Mn oxides systems consists of two irreversible steps, namely the reaction of the hydrocarbon with the lattice or trapped oxygen leading to its reduction and release of oxygen from the surface of the metal oxide and the reoxidation of the partly reduced metal oxide site by means of oxygen in a subsequent step. The TPO analysis indicates that the manganese introduction does not really influence the bulk reoxidation behavior of Co-Mn oxides [24]. Considering the TPR (**Figure 10B**) and catalytic results which reveal that Co3O4 has higher reducibility but lower activity than Co2.66Mn0.34O4, the formation of oxygen vacant site plays a key role in the redox mechanism, the importance of which was pointed out previously by Noller and Vinek [84]. According to the enhancement of the thermal stability, the amount of oxygen vacancy was increased upon manganese incorporation, which accelerates the oxidation process. Furthermore, the catalytic activity could also benefit from the substitution of cobalt with manganese that was more active and from the cooperative effect among metallic species by increasing the oxide reduction sites.

#### *3.2.2. PSE‐CVD made cobalt ferrite for low‐temperature oxidation of CO and VOCs*

C2H2, *T50* and *T90* were observed at 355 and 386°C over NCM, whereas these values decreased by respective 56 and 51°C over cobalt oxide and further decreased by 17 and 41°C over Co2.66Mn0.34O4, which reveals that the catalyst with small amount of manganese content shows the highest activity (see **Figure 10** and **Table 3**). Tian et al. [48] reported that *T50* of C2H2 oxidation with Co3O4 deposited on monolithic cordierite support was 297°C, which was in good agreement with the current Co3O4 sample prepared on mesh. Note that the *T50* of C2H2 (282°C, **Table 3**) with Co2.66Mn0.34O4 [24] was quite close to the value (280°C) reported by Ivanova et al. [80] who used Au/Al2O3, revealing that Co2.66Mn0.34O4 prepared by CVD features a competitive activity. The low-temperature shift of the light-off curve was more pronounced for C3H6 since the *T90* difference between the NCM and Co2.66Mn0.34O4 is 336°C, demonstrating that Co2.66Mn0.34O4 was very active for the total oxidation of propene and supporting the conclusion of Liang et al. [81] who reported that the manganese insertion into the cobalt oxide spinel could enhance the catalytic activity of the oxidation of light olefins. Considering the higher activity, Mn-doped spinel of cobalt oxide could be a potential system for catalytic oxidation of hydrocarbons. As presented in **Table 3**, the values of *T50* and *T90* of hydrocarbons conversion and those of CO2 selectivity were found to be almost identical for Co3O4. However, differences were observed relative to NCM especially in the oxidation of C3H6, illustrating an occurrence of partial oxidation by generating products other than CO2 (see **Figure 11f**). With Co3O4 and Co2.66Mn0.34O4, CO2 was observed to be the only carbonaceous product and no secondary products were formed all over the entire oxidation process of C3H6, while CO and CH4, coming from incomplete oxidation, were detected in addition to the production of CO2 on NCM. Even though some by-products were formed at certain temperatures with NCM, all indicate that CO2 was the final product. Taking the temperature at which hydrocarbons were completely converted to CO2 on Co2.66Mn0.34O4 sample as a reference, Co3O4 and NCM show a respective increase of 51 and 96°C for C2H2 and 36 and 358°C for C3H6, which indicate that the use of cobalt together with manganese significantly affects the catalytic activity. By applying the Arrhenius equation in the conversion range within 15% [48], the *Ea* was calculated (see **Table 3**). In general, the manganese-doped cobalt oxide exhibits lower *Ea*. Compared to those activation energies obtained in the oxidation of C2H2 and C3H6 over Co3O4-coated monolith with a large flow rate (500 sccm) [48], the values obtained on Co3O4-coated mesh were about 10 and 30 kJ mol-1 larger, the difference of which may come from the comprehensive effect of the support and inlet condition. Zhang et al. [82] and Aguilera et al. [83] also reported that the modification of Co3O4 by MnOx could promote the preferential oxidation of CO, toluene and alcohols at lower temperatures. The catalytic oxidation of the two hydrocarbons employing Co-Mn oxides systems consists of two irreversible steps, namely the reaction of the hydrocarbon with the lattice or trapped oxygen leading to its reduction and release of oxygen from the surface of the metal oxide and the reoxidation of the partly reduced metal oxide site by means of oxygen in a subsequent step. The TPO analysis indicates that the manganese introduction does not really influence the bulk reoxidation behavior of Co-Mn oxides [24]. Considering the TPR (**Figure 10B**) and catalytic results which reveal that Co3O4 has higher reducibility but lower activity than Co2.66Mn0.34O4, the formation of oxygen vacant site plays a key role in the redox mechanism, the importance of which was pointed out previously by Noller and Vinek [84]. According to the enhancement of the thermal stability,

236 Magnetic Spinels- Synthesis, Properties and Applications

Thin films of cobalt ferrite binary oxides (with the general formula Co3-xFexO4) at different composition (Co0.9Fe2.1O4, Co1.8Fe1.2O4 and Co2.1Fe0.9O4) were prepared by PSE-CVD. The systematic characterization of their properties and their potential application as catalysts for low-temperature CO, C3H6, *n*-C4H8 and C2H6O oxidation has been reported by Kouotou et al. and Tian et al. [27] The effect of iron substitution by cobalt in the structure on the optical and redox properties was investigated. The catalytic performance of the Co-Fe oxides was discussed with respect to the participation of surface and lattice oxygen in the oxidation process. According to XPS and temperature-programmed reduction/oxidation (TPR/ TPO) results, a suprafacial mechanism was the dominant mechanism for CO oxidation to CO2, while C3H6, *n‐*C4H8 and C2H6O were oxidized through an intrafacial process (MvK mechanism).

**Figure 12.** (A) Redox behaviors of Co-Fe oxides: (a) TPR, (b) TPO, (c) progressive loss of the spinel structure and (d) recover of Co-Fe-O IR vibration; (B) O1s XPS spectra of the Co-Fe-O samples; (C) HIM image displaying films morphology; (D) light-off curves of CO conversion with the Co-Fe-O samples, αFe2O3 and NCM. The results obtained over αFe2O3 [55]. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of CO and 10% of O2 diluted in Ar, corresponding the WHSV of 45,000 gcat-1 h-1. Reproduced with permission from [25]. Copyright 2014, the Royal Society of Chemistry.

CO oxidation was performed at atmospheric pressure over all samples. The results were compared with those obtained with pure α-Fe2O3 [55] and a blank sample of NCM, as shown in **Figure 12D**. Single α-Fe2O3 becomes active in CO oxidation at around 230°C and achieves complete CO conversion to CO2 above 400°C. Co-Fe-O composites exhibit lower temperatures for the initiation of CO oxidation than the single α-Fe2O3. The complete CO oxidation over Co0.9Fe2.1O4, Co1.8Fe1.2O4 and Co2.1Fe0.9O4 occurs at 255, 275 and 325°C, respectively. The results indicate that cobalt ferrites were more catalytically active than single α-Fe2O3. The obtained performance order was the following: α-Fe2O3 < Co2.1Fe0.9O4 < Co1.8Fe1.2O4 < Co0.9Fe2.1O4. It was observed that the catalytic performance was decreased with the increase of the Co content in the matrix of Co-Fe-O [25]. The sample with low Co content (Co0.9Fe2.1O4) exhibits the highest catalytic performance. An attempt to explain such behavior was made with respect to the difference in the film morphology, redox property and the chemical composition as well as the ionic state at the surface of the material as clearly displayed in **Figure 12** [25].

The TPR and TPO result shows that Co-Fe-O samples were reduced at higher temperature, increasing the Co content (**Figure 12A**). Generally, Co3O4, Co2+-Co3+ ion pairs are known to be very active in low-temperature CO oxidation [22]. Therefore, the presence in Co-Fe-O composites of both Co2+ and Co3+ together with Fe3+ in the octahedral and tetrahedral sites should enable a decrease of the reduction temperature and an improvement of the catalytic performance of CO conversion to CO2 over samples with higher Co content (Co2.1Fe0.9O4 and Co1.8Fe1.2O4), which own the [Co2+Fe3+Co3+]O-siteO4 cationic distribution in the O-site. Surprisingly, the opposite behavior was observed with Co0.9Fe2.1O4 (the most active sample) in which only Fe3+ and Co2+ were present in the octahedral sites ([Fe3+Co2+]O-siteO4). It was thus suggested that the CO oxidation over Co-Fe-O catalyst does not proceed with the redox mechanism, even if Co0.9Fe2.1O4 presents the lowest reduction temperature (**Figure 12A**). The redox mechanism hypothesis was strongly supported by the fact that CO oxidation of Co2.1Fe0.9O4 was initiated at ∼200°C while the reduction started at ∼280°C (**Figure 12A, D**). Since the formation of carbonates on the cobalt surface has been suggested by Thormählen et al. to play an important role in the low-temperature oxidation of CO [85], the earlier initiation of the reaction at low temperature was therefore assigned to the surface-adsorbed oxygen revealed by XPS analysis (**Figure 12B**). It was thus proposed that CO oxidation over Co-Fe-O follows a suprafacial mechanism where CO molecules react with adsorbed oxygen, mainly as Co3 2- and OH- , giving rise to form CO2.

In order to study further the catalytic activity of cobalt ferrite thin films, Tian et al. have investigated the catalytic performance of cobalt ferrite (Co2.1Fe0.9O4) with respect to the total oxidation of propene, *n*-butene and DME at atmospheric pressure referring to NCM [27]. The catalytic effect of the mesh has been excluded by the observation that there is no significant difference between the oxidation over NCM and in a blank system (**Figure 13**). The results show that the cobalt ferrite films favor the complete conversion of the reactants at much lower temperatures relative to NCM. Besides the reactant gas, CO2 was detected as the unique product in the oxidation processes over cobalt ferrite, while additional CO was observed in the reaction on NCM. *T50* and *T90*, corresponding to respective 50 and 90% conversion of the reactant gas, were used as parameters to compare the performance of the deposited samples. With cobalt ferrite, *T50* and *T90* of propene oxidation were 348 and 382°C. These values shifted to 578 and 691°C for the experiment carried out with NCM (see **Figure 13a**, **b**), respectively. Compared to the reaction over Co3O4 with *Ea* of 158.32 kJ/mol [24], the introduction of iron tends to initiate the oxidation of propene with lower *Ea*, which makes the cobalt ferrite more suitable for the catalytic applications. **Figure 13c** and **d** compares the results of *n*-C4H8 oxidation with and without cobalt ferrite. *T50* and *T90* of *n*-C4H8 oxidation over cobalt ferrite were observed at 358 and 402°C, whereas these values shifted by respective 100 and 135°C toward higher temperatures over NCM. It has been reported that a small quantity of 1,3butadiene was selectively formed in the oxidation of *n‐*C4H8 at temperature higher than 350°C over MnMoO4-based catalysts [86]. It should be noted that 1,3-butadiene was not detected in the reaction with the PSE-CVD made Co2.1Fe0.9O4 indicating that the material presents higher catalytic activity.

observed that the catalytic performance was decreased with the increase of the Co content in the matrix of Co-Fe-O [25]. The sample with low Co content (Co0.9Fe2.1O4) exhibits the highest catalytic performance. An attempt to explain such behavior was made with respect to the difference in the film morphology, redox property and the chemical composition as well as the

The TPR and TPO result shows that Co-Fe-O samples were reduced at higher temperature, increasing the Co content (**Figure 12A**). Generally, Co3O4, Co2+-Co3+ ion pairs are known to be very active in low-temperature CO oxidation [22]. Therefore, the presence in Co-Fe-O composites of both Co2+ and Co3+ together with Fe3+ in the octahedral and tetrahedral sites should enable a decrease of the reduction temperature and an improvement of the catalytic performance of CO conversion to CO2 over samples with higher Co content (Co2.1Fe0.9O4 and Co1.8Fe1.2O4), which own the [Co2+Fe3+Co3+]O-siteO4 cationic distribution in the O-site. Surprisingly, the opposite behavior was observed with Co0.9Fe2.1O4 (the most active sample) in which only Fe3+ and Co2+ were present in the octahedral sites ([Fe3+Co2+]O-siteO4). It was thus suggested that the CO oxidation over Co-Fe-O catalyst does not proceed with the redox mechanism, even if Co0.9Fe2.1O4 presents the lowest reduction temperature (**Figure 12A**). The redox mechanism hypothesis was strongly supported by the fact that CO oxidation of Co2.1Fe0.9O4 was initiated at ∼200°C while the reduction started at ∼280°C (**Figure 12A, D**). Since the formation of carbonates on the cobalt surface has been suggested by Thormählen et al. to play an important role in the low-temperature oxidation of CO [85], the earlier initiation of the reaction at low temperature was therefore assigned to the surface-adsorbed oxygen revealed by XPS analysis (**Figure 12B**). It was thus proposed that CO oxidation over Co-Fe-O follows a suprafacial

ionic state at the surface of the material as clearly displayed in **Figure 12** [25].

238 Magnetic Spinels- Synthesis, Properties and Applications

mechanism where CO molecules react with adsorbed oxygen, mainly as Co3

In order to study further the catalytic activity of cobalt ferrite thin films, Tian et al. have investigated the catalytic performance of cobalt ferrite (Co2.1Fe0.9O4) with respect to the total oxidation of propene, *n*-butene and DME at atmospheric pressure referring to NCM [27]. The catalytic effect of the mesh has been excluded by the observation that there is no significant difference between the oxidation over NCM and in a blank system (**Figure 13**). The results show that the cobalt ferrite films favor the complete conversion of the reactants at much lower temperatures relative to NCM. Besides the reactant gas, CO2 was detected as the unique product in the oxidation processes over cobalt ferrite, while additional CO was observed in the reaction on NCM. *T50* and *T90*, corresponding to respective 50 and 90% conversion of the reactant gas, were used as parameters to compare the performance of the deposited samples. With cobalt ferrite, *T50* and *T90* of propene oxidation were 348 and 382°C. These values shifted to 578 and 691°C for the experiment carried out with NCM (see **Figure 13a**, **b**), respectively. Compared to the reaction over Co3O4 with *Ea* of 158.32 kJ/mol [24], the introduction of iron tends to initiate the oxidation of propene with lower *Ea*, which makes the cobalt ferrite more suitable for the catalytic applications. **Figure 13c** and **d** compares the results of *n*-C4H8 oxidation with and without cobalt ferrite. *T50* and *T90* of *n*-C4H8 oxidation over cobalt ferrite were observed at 358 and 402°C, whereas these values shifted by respective 100 and 135°C toward higher temperatures over NCM. It has been reported that a small quantity of 1,3-

rise to form CO2.

2- and OH-

, giving

**Figure 13.** Outlet profiles of C3H6, n-C4H8 and DME oxidation over cobalt ferrite-coated and non-coated meshes. The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of Fuel and 10% of O2 diluted in Ar, corresponding the WHSV of 45,000 gcat-1 h-1. Reproduced with permission from [27]. Copyright 2015, Elsevier.

For DME, cobalt ferrite also exhibits much better performance than NCM, as revealed in **Figure 13e** and **f**. *T50* and *T90* of the reaction over cobalt ferrite are 356 and 409°C. These values were observed to be 613 and 682°C for the reaction on NCM, respectively. According to Liu et al. [87], the reaction network for DME conversion contains four pathways giving rise to CH3OH, HCHO, HCOOCH3 and COx. Low selectivity of CH3OH and HCOOCH3 was reported [87]. HCHO was not detected, which could be resulted either from the low concentration or fast conversion to COx.

The good catalytic performance of cobalt ferrite could be correlated with the Oadsorbed on the surface and attractive redox properties (**Figure 12A**, **B**). It is widely accepted that the catalytic performance of ferrite-type catalysts depends on its oxygen mobility since the reaction follows MvK mechanism [88]. The abundance of OAdsorbed could participate in the catalytic oxidation process, as proposed by Veleva and Trifirò [86]. The good reducibility and reoxidability tend to enable the catalytic reactions at relatively low temperatures by involving various cations distributed in the octahedral and tetrahedral sites. Moreover, the slightly low band gap energy of cobalt ferrite could also indicate that the migration of Olattice or O2- from the bulk to the surface becomes easier and leads to good reducibility [25].

Because cobalt ferrite binary oxide with high Co concentration (Co2.1Fe0.9O4) showed very interesting catalytic performance against the deep oxidation of C3H6, *n*-C4H8 and DME, Kouotou et al. [26] explored the catalytic activity of the same material, but with low Co concentration (CoFe2O4). Authors aimed to compare the performances of CoFe2O4 (low cobalt content) versus Co2.1Fe0.9O4 (high cobalt content) from recent investigation toward catalytic oxidation of CO, DME, C3H6 and *n*-C4H8. The role of Co or Fe atom in the matrix of cobalt ferrite mixed oxide and adsorbed/lattice oxygen at the surface of the deposited material in the catalytic reaction was identified. The light-off curves of CO, C3H6, *n*-C4H8 and DME oxidation over CoFe2O4 are shown in **Figure 14**. The results indicate that CoFe2O4 is a very active catalyst for the total oxidation of CO, C3H6, *n*-C4H8 and DME at atmospheric pressure. Total destruction of such pollutants at low temperature enables the production of CO2 as the only detectable Ccontaining product. In the presence of CoFe2O4, the oxidation becomes detectable at around 150, 200, 230 and 270°C for CO, C3H6, *n*-C4H8 and DME, respectively and complete conversion occurs at 253°C for CO, 336°C for C3H6 and n-C4H8 and 502°C for DME.

**Figure 14.** Light-off curves of CO, C3H6, n-C4H8 and DME oxidation over the PSE-CVD made CoFe2O4 sample. Reproduced with permission from [26]. Copyright 2015, Wiley-VCH.

For the CO oxidation, the comparison of the temperature at 90% conversion (*T90*) over CoFe2O4 (260°C) as catalyst with that obtained with Co0.9Fe2.1O4 (230°C), Co1.8Fe1.2O4 (250°C) and Co2.1Fe0.9O4 (298°C) from our previous work has been made. The results indicate the following performance order: Co0.9Fe2.1O4 > Co1.8Fe1.2O4 > CoFe2O4 > Co2.1Fe0.9O4. CoFe2O4 owns similar composition as Co0.9Fe2.1O4, but it exhibits less activity for CO oxidation. By the way, the comparison of the T50 of CO over CoFe2O4 (243°C) with results obtained with Pt/Al2O3 (304°C) and Pt/CoOx/Al2O3 (340°C) catalysts reported by Torncrona et al. [89] indicates that CoFe2O4 is more active at low temperature.

In the investigation of the CO oxidation over cobalt ferrite oxides with different composition (Co0.9Fe2.1O4, Co1.8Fe1.2O4 and Co2.1Fe0.9O4) [25], it was established that the performance of cobalt ferrite oxide was not dependent on the variation of the Co content in the mixed oxides. Therefore, these results suggest that CO oxidation over CoFe2O4 follows a suprafacial mechanism where CO molecules react with adsorbed oxygen at the surface of the catalyst to form CO2. Considering that the difference in the catalytic performance of CO over CoFe2O4 with seemingly the same composition as Co2.1Fe0.9O4 [25], it was therefore evident that the process is independent on the Co population in the cobalt ferrite oxide, but on the lattice and adsorbed oxygen, mainly adsorbed oxygen at the surface of the material.

content) versus Co2.1Fe0.9O4 (high cobalt content) from recent investigation toward catalytic oxidation of CO, DME, C3H6 and *n*-C4H8. The role of Co or Fe atom in the matrix of cobalt ferrite mixed oxide and adsorbed/lattice oxygen at the surface of the deposited material in the catalytic reaction was identified. The light-off curves of CO, C3H6, *n*-C4H8 and DME oxidation over CoFe2O4 are shown in **Figure 14**. The results indicate that CoFe2O4 is a very active catalyst for the total oxidation of CO, C3H6, *n*-C4H8 and DME at atmospheric pressure. Total destruction of such pollutants at low temperature enables the production of CO2 as the only detectable Ccontaining product. In the presence of CoFe2O4, the oxidation becomes detectable at around 150, 200, 230 and 270°C for CO, C3H6, *n*-C4H8 and DME, respectively and complete conversion

**Figure 14.** Light-off curves of CO, C3H6, n-C4H8 and DME oxidation over the PSE-CVD made CoFe2O4 sample. Repro-

For the CO oxidation, the comparison of the temperature at 90% conversion (*T90*) over CoFe2O4 (260°C) as catalyst with that obtained with Co0.9Fe2.1O4 (230°C), Co1.8Fe1.2O4 (250°C) and Co2.1Fe0.9O4 (298°C) from our previous work has been made. The results indicate the following performance order: Co0.9Fe2.1O4 > Co1.8Fe1.2O4 > CoFe2O4 > Co2.1Fe0.9O4. CoFe2O4 owns similar composition as Co0.9Fe2.1O4, but it exhibits less activity for CO oxidation. By the way, the comparison of the T50 of CO over CoFe2O4 (243°C) with results obtained with Pt/Al2O3 (304°C) and Pt/CoOx/Al2O3 (340°C) catalysts reported by Torncrona et al. [89] indicates that

In the investigation of the CO oxidation over cobalt ferrite oxides with different composition (Co0.9Fe2.1O4, Co1.8Fe1.2O4 and Co2.1Fe0.9O4) [25], it was established that the performance of cobalt ferrite oxide was not dependent on the variation of the Co content in the mixed oxides. Therefore, these results suggest that CO oxidation over CoFe2O4 follows a suprafacial mech-

duced with permission from [26]. Copyright 2015, Wiley-VCH.

CoFe2O4 is more active at low temperature.

occurs at 253°C for CO, 336°C for C3H6 and n-C4H8 and 502°C for DME.

240 Magnetic Spinels- Synthesis, Properties and Applications

The oxidation of C3H6 and *n*-C4H8 on CoFe2O4 starts at temperatures lower than that of DME and the trend for the rate of oxidation is C3H6 ≈ *n‐*C4H8 > DME [26]. In the presence of CoFe2O4 nanoparticles, the oxidation becomes detectable at around 250°C for C3H6 and *n‐* C4H8 and 285°C for DME and the T90 is 334, 371 and 470°C for C3H6, *n‐*C4H8 and DME, respectively. The temperature at the complete oxidation was 475°C for C3H6, 485°C for *n‐*C4H8 and 506°C for DME. These values appear to be higher than those for Co2.1Fe0.9O4 catalysts, where complete oxidation was observed at 400, 425 and 446°C for C3H6, *n‐*C4H6 and DME, respectively [27]. However, they were lower than those observed for C3H6 on Au/Al2O3 [80] and La1.7Sr0.3CuO4S0.2 [90]. The activity in *n‐*C4H8 oxidation was comparable to the measurement on Co3O4 [48]. In general, for the oxidation reactions over CoFe2O4 at low temperatures, no toxic or partial oxidation products are formed when complete conversion of VOCs is achieved. For DME oxidation, it has been established that the reaction network consists of four pathways giving rise to CH3OH, HCHO, HCOOCH3 and COx [87]. Kouotou et al. [26] detected only CO2, indicating the higher activity of CoFe2O4 catalyst for total oxidation of DME at low temperature.

## *3.2.3. PSE‐CVD made CuCo2O4 for low‐temperature catalytic combustion of CO and VOCs*

To investigate the potential application of the grown Co-Cu oxides, Tian et al. [28] investigated the catalytic performance of the prepared samples with respect to the total oxidation of CO and C3H6. The authors carried out the same tests on bare meshes to exclude the background effect of the mesh structures on the oxidation processes. **Figure 15** compares the light-off curves of CO and C3H6 by using Co3O4 and Co-Cu oxide referring to the bare mesh. With Co3O4, the conversion of CO occurs at 270°C and complete conversion was achieved at 362°C. For Co-Cu oxide, these values shift to 125 and 221°C, respectively. CO oxidation over bare mesh was observed to begin at around 395°C and finish at 833°C. Under similar conditions, CO oxidation on other TMOs such as Mn3O4 was reported to happen at 190°C and complete at 343°C [43]. Compared to Mn3O4, Co-Cu oxide exhibits better performance. This could be explained by the abundance of adsorbed oxygen, Co3+ and Cu2+ revealed in the XPS analysis. By applying the Arrhenius equation at low conversion profiles (within 5%), the *Ea* values were calculated to be 68.2, 77.3 and 82.9 kJ mol-1 for the CO oxidation with Co3O4, Co-Cu oxide and bare mesh, respectively. As shown in **Figure 15b**, both Co-Cu oxide and Co3O4 improve the catalytic oxidation of C3H6 compared to the bare mesh. In the presence of Co-Cu oxide, the propene conversion becomes detectable at 220°C. This value was 30 and 120°C lower than that with Co3O4 and bare mesh, respectively. With Co-Cu oxide and Co3O4, complete conversion of C3H6 was reached within 412 and 454°C, while this value was observed to be 772°C for the reaction on NCM. It should be mentioned that CO2 is the unique measurable product and no trace of CO was detected in the oxidation process with the prepared CuCo2O4 and Co3O4, as

depicted in **Figure 16**. This result agrees well that the conclusion drawn in our recent work that the active transition metal oxides prevent the formation of CO in the oxidation of lowrank hydrocarbons. In the reaction on bare mesh, a large amount of CO peaking at 758°C was produced, which comes from partial oxidation.

**Figure 15.** Catalytic performance of Co-Cu oxides: oxidation of CO (a) and propene (b). The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of fuel and 10% of O2 diluted in Ar, corresponding the WHSV of 75,000 gcat-1 h-1. Reproduced with permission from [28]. Copyright 2014, the Royal Society of Chemistry.

**Figure 16.** Outlet profiles of CO in the propene oxidation over CuCo2O4, Co3O4 and NCM. Reproduced with permission from [28]. Copyright 2014, the Royal Society of Chemistry.

Compared to the reported *T*50 values in the temperature-programmed reaction, 330°C obtained with CuCo2O4 was lower than that measured with Au/Al2O3 (349°C) [80], Co3O4 (354°C) [48], Au/MgO/Al2O3 (359°C) [91], La1.7Sr0.3CuO4S0.2 (419°C) [90], (327°C) [91] and Co2.66Mn0.34O4 (321°C) [24]. Although CuO shows better performance with the *T*<sup>50</sup> at 272°C [44], CuCo2O4 has more application prospects than CuO by considering the higher thermal stability. It is worth mentioning that the magnitude order of *T*<sup>50</sup> for CuO, CuCo2O4 and Co3O4 follows the order of the band gaps [28]. It can be explained that the migration of O2- from the bulk to the surface gets easier for samples with lower band gaps which have better electron mobility. Moreover, the *Ea* for propene oxidation with CuCo2O4 is 82.6 kJ mol-1, which is lower than that obtained with Co3O4 (130 kJ mol-1) [24], CuO (109.5 kJ mol-1) [44], Co2.66Mn0.34O4 (115.5 kJ mol-1) [24] as well as the non-catalyzed reaction (138.3 kJ mol-1). The MvK mechanism has been established for CO and propene oxidation over CuO [44], Co–Mn oxides and Co–Fe oxides [25–27]. The same redox behavior could exist in the reaction with Co-Cu oxides.

#### *3.2.4. Comparison of the catalytic performances of binary TMOs with literature*

depicted in **Figure 16**. This result agrees well that the conclusion drawn in our recent work that the active transition metal oxides prevent the formation of CO in the oxidation of lowrank hydrocarbons. In the reaction on bare mesh, a large amount of CO peaking at 758°C was

**Figure 15.** Catalytic performance of Co-Cu oxides: oxidation of CO (a) and propene (b). The reaction was carried at the total flow rate was kept at 15 ml min-1 with 1% of fuel and 10% of O2 diluted in Ar, corresponding the WHSV of 75,000

**Figure 16.** Outlet profiles of CO in the propene oxidation over CuCo2O4, Co3O4 and NCM. Reproduced with permis-

Compared to the reported *T*50 values in the temperature-programmed reaction, 330°C obtained with CuCo2O4 was lower than that measured with Au/Al2O3 (349°C) [80], Co3O4 (354°C) [48],

sion from [28]. Copyright 2014, the Royal Society of Chemistry.

gcat-1 h-1. Reproduced with permission from [28]. Copyright 2014, the Royal Society of Chemistry.

produced, which comes from partial oxidation.

242 Magnetic Spinels- Synthesis, Properties and Applications

**Table 4** compares the catalytic performance of the PSE-CVD made mixed TMOs with that of the selected catalysts in the literature. It can be noted that our catalysts present comparable activity to the other systems. Special attention is paid to the comparison of the catalytic performance with noble metals and transition metal oxides. As can be seen from **Table 4**, noble metals and TMOs (mixed) present attractive results in terms of CO, propene, *n‐*butene and DME total oxidation.



b : *Temperature at X% conversion (X = 10, 50 and 90 %)*

**Table 4.** Overview of the catalytic performances of PSE-CVD made binary TMOs compared with the literature.

## **4. Remarks and perspectives**

This review highlights the newly developed CVD approach called PSE-CVD for single and binary TMOs thin films coating and their performances in the catalytic oxidation of CO and VOCs. The advances in the synthesis and properties of single and binary oxides thin films have shown potential applications in several catalytic processes. In this review, we first presented different single oxides fabricated by using a more elaborated CVD method. In particular, the complete oxidation of CO and VOCs over Co3O4, CuO/Cu2O, Fe2O3, Cr2O3 as well as Mn3O4 is comprehensively summarized. In addition, the catalytic activity in CO and VOCs oxidation is systematically compared and it can be well-documented that the CO catalytic oxidation activity is related to either the abundance of active metallic cation sites (Co3+, Cu2+, Fe3+, Cr3+ and Mn3+) present in Co3O4, CuO/Cu2O, Fe2O3, Cr2O3 as well as Mn3O4, respectively, or the weakly bonded molecular oxygen species at their surface. On the basis of the present results, we could conclude that the VOCs catalytic oxidation activity of single TMOs is related with either the reactive surface oxygen species or bulk oxygen mobility, reoxidation of metallic cations species and active oxygen vacancies of single oxides.

Regarding binary systems, Co-based binary oxides generally prepared by co-precipitation, impregnation, surfactant-template methods and by combined impregnation and combustion synthesis were prepared by PSE-CVD and summarized in the present review. The idea was to formulate active and stable mixed oxide based on the performance obtained with single oxide. In fact, Co3O4 is known as the most active single oxides for CO and VOCs deep oxidation and the activity strongly depends on Co3+/Co2+ redox couple. Compared with pure Co3O4 catalysts, the binary oxide catalysts are generally composed of one or more catalytically active components and a functional support, in which the interaction between the catalytic components and the supports can provide improved redox properties and enhanced catalytic behavior [92–94]. As for single oxide catalysts, the active component alone can catalyze the various reactions at a certain kinetic rate, which is usually relatively low. The synergistic effect between binary oxide catalysts leads to the improved catalytic activity. Therefore, the combination of Co and individual metal of single oxide such as Fe2O3, Mn3O4 and CuO enables the obtaining of binary oxide like Co-Mn-O, Co-Fe-O and Co-Cu-O with new chemical and cationic distribution benefits to the obtained catalytic performances summarized in this review. From the obtained results, it seems like combining two transition metals influence the interaction of active cobalt oxide with the second oxide, affecting the textural and structural properties of the binary oxides. In particular, the investigations of the interaction mechanisms as a function of the surface redox properties of the binary oxides are comprehensively summarized. The catalytic activity of the binary oxides in CO and VOC oxidation is compared and different reaction mechanisms occurring in CO and VOC oxidation are presented. Probably, the surface adsorbed oxygen species over the surface oxygen vacancies of the catalysts are mainly attributed to the CO oxidation activity at low temperatures. The total oxidations of VOCs over transition metal oxides are suggested to follow the MvK redox process.

**Type of catalyst Material Weight (mg) Gas composition WHSV** *<sup>a</sup>*

a

b

: *Weight hourly space velocity*

: *Temperature at X% conversion (X = 10, 50 and 90 %)*

244 Magnetic Spinels- Synthesis, Properties and Applications

**4. Remarks and perspectives**

cations species and active oxygen vacancies of single oxides.

**(ml g‐1 h‐1)** 

CuCo2O4 12 1% CO/10% O2 89%Ar 75,000 – 190 [28] CoFe2O4 20 45,000 – – 260 [23] Co0.9Fe2.1O4 – – 230 [25]

Pt/Al2O3 2/200 1% CO/1.38%O2 in N2 90,000 – 333 430 [89] Au/SiO2 100 1% CO/99% dry air 12,000 – 337 423 [99]

Co2.1Fe0.9O4 1% *n*C4H8/10%O2 in Ar – 358 402 [27]

Co2.1Fe0.9O4 – 356 409 [27]

Co2.1Fe0.9O4 – – 298 Co1.8Fe1.2O4 – – 250

Precious metals Pt/H2SO4/ZrO2 50 3.5% CO/4%O2 in N2 120,000 – – 290 [98]

Mixed TMOs CoFe2O4 20 1% *n‐*C4H8/10% O2 in Ar 45,000 – 371 [23]

Mixed TMOs CoFe2O4 20 1% C2H6O/10%O2 in Ar 45,000 – – 470 [26]

**Table 4.** Overview of the catalytic performances of PSE-CVD made binary TMOs compared with the literature.

This review highlights the newly developed CVD approach called PSE-CVD for single and binary TMOs thin films coating and their performances in the catalytic oxidation of CO and VOCs. The advances in the synthesis and properties of single and binary oxides thin films have shown potential applications in several catalytic processes. In this review, we first presented different single oxides fabricated by using a more elaborated CVD method. In particular, the complete oxidation of CO and VOCs over Co3O4, CuO/Cu2O, Fe2O3, Cr2O3 as well as Mn3O4 is comprehensively summarized. In addition, the catalytic activity in CO and VOCs oxidation is systematically compared and it can be well-documented that the CO catalytic oxidation activity is related to either the abundance of active metallic cation sites (Co3+, Cu2+, Fe3+, Cr3+ and Mn3+) present in Co3O4, CuO/Cu2O, Fe2O3, Cr2O3 as well as Mn3O4, respectively, or the weakly bonded molecular oxygen species at their surface. On the basis of the present results, we could conclude that the VOCs catalytic oxidation activity of single TMOs is related with either the reactive surface oxygen species or bulk oxygen mobility, reoxidation of metallic

Regarding binary systems, Co-based binary oxides generally prepared by co-precipitation, impregnation, surfactant-template methods and by combined impregnation and combustion

**T10** *<sup>b</sup>*  **(°C)**

**T50** *<sup>b</sup>*  **(°C)**

**T90** *<sup>b</sup>*  **(°C)**

**Refs.**

Nevertheless, in the case of VOCs oxidation, the catalytic activity depends on the nature of catalysts and VOCs. While CO oxidation over single oxide seems to be mainly governed by the amount of available surface oxygen species, VOCs oxidation activity of PSE-CVD made single oxide in contrary is scarcely determined by the surface oxygen species. In the case of CO oxidation, PSE-CVD made single and binary oxides show relevant concentration of surface oxygen species which might be involved in the oxidation at low temperatures. Moreover, the higher activity of Co3O4 and Co0.9Fe2.1O4 binary oxides with respect to the Fe2O3 or CuO with respect to CuCO2O4 and other binary oxides is mainly attributed to their higher mobility of lattice oxygen species.

The investigation of the catalytic properties of single and binary TMOs made by PSE-CVD has provide valuable results in the literature, about the employ of a well-established nonconventional CVD approach for the selective synthesis of oxide made of Co, Cu, Fe, Cr and Mn and their performances toward low temptation conversion of CO, C2H2, C3H6, n-C4H8 and C2H6O in harmless CO2 and H2O. As perspectives, we think that the following aspects deserving further investigations are: (1) further research study of the interactions between different kinds of binary oxide catalytic systems and related corrections between their redox properties and catalytic activities; (2) investigation of the interactions among three or more component catalytic systems, such as noble metal doped or noble metal alloys doped Co-Mn, Co-Fe and Co-Cu binary oxides. It is expected that the novel structures matrix can offer new opportunities to expand our understanding of this kind of interaction as well as relationship between structure and property; (3) because of his lattice oxygen reservoir character when used as supporter, single Co3O4, CuO, Fe2O3 crystals structure needs to be deposited on CeO2 nanoparticles in order to control this kind of interaction between these two oxides and to enhance application potentials of these multifunctional materials. It is expected that such an integration of the current investigations on different systems will bring a more comprehensive understanding of the interactions during the catalytic process and therefore will be of great significance in searching for and find novel multi-oxide catalysts of further enhanced catalytic activity, selectivity as well as durability.

## **5. Summary**

The catalysis science of non-precious metals has significantly improved recently, due to development of more elaborated synthesis approaches for their controlled synthesis and advanced characterization techniques that allow more fundamental insights into the reaction mechanisms. As summary of the present review on the low-temperature oxidation of CO and VOCs over selected non-precious metal such as TMOs catalytic materials either in their single and binary or mixed phases prepared using PSE-CVD, the following main results can be summarized:


## **Conflict of interest**

structure and property; (3) because of his lattice oxygen reservoir character when used as supporter, single Co3O4, CuO, Fe2O3 crystals structure needs to be deposited on CeO2 nanoparticles in order to control this kind of interaction between these two oxides and to enhance application potentials of these multifunctional materials. It is expected that such an integration of the current investigations on different systems will bring a more comprehensive understanding of the interactions during the catalytic process and therefore will be of great significance in searching for and find novel multi-oxide catalysts of further enhanced catalytic

The catalysis science of non-precious metals has significantly improved recently, due to development of more elaborated synthesis approaches for their controlled synthesis and advanced characterization techniques that allow more fundamental insights into the reaction mechanisms. As summary of the present review on the low-temperature oxidation of CO and VOCs over selected non-precious metal such as TMOs catalytic materials either in their single and binary or mixed phases prepared using PSE-CVD, the following main results can be

**1.** The catalytic combustion has become the most popular method for the environmental emission control. Thus, catalytic oxidation of CO and VOCs is highly desirable to proceed at low temperature for the consideration of energy savings, low cost, operation safety and

**2.** To reduce the temperature of VOCs catalytic oxidation, great efforts have been made to develop efficient and low cost catalysts via an elaborated synthesis method, namely pulsed spray evaporation chemical vapor deposition. The grid mesh of stainless steel as inert substrates was used instead of catalytically active support to enable the real evaluation of the catalytic performance of the as-deposited thin films layer toward CO and

**3.** TMOs exhibited comparative activity versus noble metals toward the catalytic oxidation of VOCs at low temperature. The catalytic performance of TMOs was found to be generally affected by many factors such as the composition, the valence of metallic particles, doping,

**4.** Single oxides systems composed of transition metals (Co, Cu, Cr, Fe and Mn) are efficient for the abatement of CO and VOCs. In addition, the effect of doping should be taken into consideration since they greatly improve the thermal properties and the catalytic performances of the as-prepared materials. For instance, the combination of two transition metals or the doping that is the introduction of another transition metal ion in the matrix of single oxides to form a mixed oxide catalysts such as Co-Mn, Co-Fe and Co-Cu oxides also contributes to stabilize the as-deposited catalysts, enable the total oxidation of VOCs

VOCs oxidation, offering the major advantage of low-pressure drop.

the film morphology and the particle size of the metals oxide.

and exhaust stream at low temperature.

activity, selectivity as well as durability.

246 Magnetic Spinels- Synthesis, Properties and Applications

environmental friendliness.

**5. Summary**

summarized:

The authors confirm that this review article has no conflict of interest.

## **Acknowledgements**

Prof. Dr. Tian and Dr. Mountapmbeme Kouotou are grateful for the support from the Recruitment Program of Global Youth Experts and the Chinese Academy of Sciences visiting Professorship for Senior International scientist (grant no. 2015PT016). The authors are grateful to Prof. Dr Katharina Kohse-Höinghaus for her support and discussions and for allowing us to perform part of this research in her laboratory in Bielefeld.

## **Author details**

Patrick Mountapmbeme Kouotou1,3\*, Guan-Fu Pan1,2 and Zhen-Yu Tian1,2\*


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#### **Spinel‐Structured Nanoparticles for Magnetic and Mechanical Applications Spinel‐Structured Nanoparticles for Magnetic and Mechanical Applications**

Malik Anjelh Baqiya, Ahmad Taufiq, Sunaryono, Khuroti Ayun, Mochamad Zainuri, Suminar Pratapa, Triwikantoro and Darminto Malik Anjelh Baqiya, Ahmad Taufiq, Sunaryono, Khuroti Ayun, Mochamad Zainuri, Suminar Pratapa, Triwikantoro and Darminto

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

http://dx.doi.org/10.5772/66293

#### **Abstract**

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144–150.

Nanoparticles of Fe3O4 have been successfully synthesized using a simple coprecipita‐ tion technique from natural iron sands, employing HNO3 and NH4OH as dispersing and precipitating agents, respectively. The substitution of Fe with Mn to result in Fe3‐ *<sup>x</sup>*Mn*x*O4 (0 ≤ *x* ≤ 3) was conducted to control the magnetic strength of this nano‐sized spinel powder. It is shown that magnetic properties depend not only on the particle size and Mn doping but also on the particles clustering. The applications for magnetic fluids, gels, and coating are extensively described. Meanwhile, the spinel MgAl2O4 nanoparti‐ cles have also been prepared by the same simple method from commercial starting materials. This powder was used as a nano‐reinforcer of Al‐matrix composites. In addition, MgAl2O4 micro‐sized powder forming a thick layer was successfully grown by electroless plating on the interface of matrix‐filler in Al/SiC composites. The strengthening of mechanical properties with respect to the varying uses of these MgAl2O4 powders is discussed.

**Keywords:** Fe3O4, MgAl2O4, Powders, Coprecipitation, Magnetic, mechanical proper‐ ties

## **1. Introduction**

Spinel structures have the general formula of AB2X4, where X can be oxygen (oxides) or a chalcogen element, such as sulfur (thio‐spinels) and selenium (seleno‐spinels). A and B in the spinel structures can be divalent, trivalent, or tetravalent cations, such as iron, magnesium,

© 2017 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. © 2017 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.

aluminum (Al), zinc, manganese, chromium, and titanium. The various compounds of spinel family including their chemical and physical properties depend not only on the arrangement of cations and anions in the structures but also on how the spinels are produced. The spinels can be magnetic or non‐magnetic compounds depending on the nature of the cations among the A and B sub‐lattices. Fe3O4 (FeO.Fe2O3) is one of the so‐called spinel ferrite having magnetic properties, whereas MgAl2O4 is also spinel (should not be confused with the spinel structure described earlier) which is non‐magnetic one. This chapter covers some of the physical properties of the spinel compounds in more detail.

## **2. Magnetite (Fe3O4) nanoparticles**

Fe3O4 (magnetite) or the black iron oxide in the form of nanoparticles has huge potential for applications in many fields. However, preparing Fe3O4 particles in small size with high quality is enormously challenging for several reasons. Therefore, the preparation of Fe3O4 nanoparti‐ cles in several forms and various sizes with high purity and homogeneity becomes an essential task before applications. The following section provides more details on the preparation and applications of Fe3O4 nanoparticles.

Several synthesis methods to prepare high‐purity Fe3O4 nanoparticles in various forms and sizes have been reported in the literature. In general, Fe3O4 nanoparticles have been success‐ fully prepared through sonochemical [1], thermal decomposition [2], ionothermal [3], hydro‐ thermal [4], micro‐emulsion [5], sol‐gel [6], modified hydrolysis [7], solvothermal [8], electrospinning [9], coprecipitation [10] methods, and so forth. However, high quality of the particles, such as good particle size distribution with high crystalline, structural and compo‐ sitional homogeneity, still has become one of the main challenges in preparing Fe3O4 nano‐ particles [11].

The improvement of synthesis methods has also been carried out for producing Fe3O4 nanoparticles in order to reduce cost and time of preparation. A simple method to produce Fe3O4 nanoparticles via coprecipitation method has been introduced using a local natural iron sand from Indonesia as a starting material. This coprecipitation method is considered to be the easiest method for preparing Fe3O4 nanoparticles [12] because of its simplicity, efficiency, and inexpensiveness [13]. The use of starting material from commercial product can be substituted by natural iron sand.

As for the preparation of Fe3O4 nanoparticles using coprecipitation method, the ratio of ferrous and ferric ions is very crucial affecting the synthesis result. The effects include the particle diameter and magnetic properties. Iida et al. [14] have reported that the valence of the metal salt is an important factor for determining the particle size of the magnetic particles. In their experiments, they have succeeded in obtaining Fe3O4 nanoparticles in the size range of 9–40 nm with various molar contents of ferrous ions in the total amount of the iron ions. Further‐ more, Gnanaprakash et al. [15] have also reported that initial pH and temperature of the iron salt solution during coprecipitation process have strong influences for the formation of magnetite nanoparticles. They reported that the average of particle size having initial pH value of 0.7, 1.5, 3, and 4.7 are 6.7, 7.6, 9.9, and 9.9 nm, respectively. Determination of particle size is very important according to how they will be used for certain applications, namely for magnetic fluids and gels.

In the magnetic nanoparticles' syntheses, the so‐called templating method has been employed to form certain particle shapes and sizes, such as one dimensional (1‐D) particles or nanorods. Lian et al. [16] have successfully synthesized nanorods of Fe3O4 through the addition of polyethylene glycol (PEG) 1000 as a template with the average particle diameter and length of 80 nm and 2 µm, respectively. Therefore, templating technique has important role on the particle growth of magnetic nanomaterials. The greater value of the PEG's molecular weight will lead to the higher possibility to obtain nanorods resulted from the synthesis [17]. Fur‐ thermore, based on the research of He et al. [18], Fe3O4 nanowire has successfully been prepared depending on the volume ratio of PEG and H2O. They have found that the best formation is achieved with the ratio of 1:3; a higher ratio gives higher viscosity and inhibits the particle formation.

## **3. Fe3O4 spinels for magnetic fluids and gels**

aluminum (Al), zinc, manganese, chromium, and titanium. The various compounds of spinel family including their chemical and physical properties depend not only on the arrangement of cations and anions in the structures but also on how the spinels are produced. The spinels can be magnetic or non‐magnetic compounds depending on the nature of the cations among the A and B sub‐lattices. Fe3O4 (FeO.Fe2O3) is one of the so‐called spinel ferrite having magnetic properties, whereas MgAl2O4 is also spinel (should not be confused with the spinel structure described earlier) which is non‐magnetic one. This chapter covers some of the physical

Fe3O4 (magnetite) or the black iron oxide in the form of nanoparticles has huge potential for applications in many fields. However, preparing Fe3O4 particles in small size with high quality is enormously challenging for several reasons. Therefore, the preparation of Fe3O4 nanoparti‐ cles in several forms and various sizes with high purity and homogeneity becomes an essential task before applications. The following section provides more details on the preparation and

Several synthesis methods to prepare high‐purity Fe3O4 nanoparticles in various forms and sizes have been reported in the literature. In general, Fe3O4 nanoparticles have been success‐ fully prepared through sonochemical [1], thermal decomposition [2], ionothermal [3], hydro‐ thermal [4], micro‐emulsion [5], sol‐gel [6], modified hydrolysis [7], solvothermal [8], electrospinning [9], coprecipitation [10] methods, and so forth. However, high quality of the particles, such as good particle size distribution with high crystalline, structural and compo‐ sitional homogeneity, still has become one of the main challenges in preparing Fe3O4 nano‐

The improvement of synthesis methods has also been carried out for producing Fe3O4 nanoparticles in order to reduce cost and time of preparation. A simple method to produce Fe3O4 nanoparticles via coprecipitation method has been introduced using a local natural iron sand from Indonesia as a starting material. This coprecipitation method is considered to be the easiest method for preparing Fe3O4 nanoparticles [12] because of its simplicity, efficiency, and inexpensiveness [13]. The use of starting material from commercial product can be substituted

As for the preparation of Fe3O4 nanoparticles using coprecipitation method, the ratio of ferrous and ferric ions is very crucial affecting the synthesis result. The effects include the particle diameter and magnetic properties. Iida et al. [14] have reported that the valence of the metal salt is an important factor for determining the particle size of the magnetic particles. In their experiments, they have succeeded in obtaining Fe3O4 nanoparticles in the size range of 9–40 nm with various molar contents of ferrous ions in the total amount of the iron ions. Further‐ more, Gnanaprakash et al. [15] have also reported that initial pH and temperature of the iron salt solution during coprecipitation process have strong influences for the formation of magnetite nanoparticles. They reported that the average of particle size having initial pH value

properties of the spinel compounds in more detail.

**2. Magnetite (Fe3O4) nanoparticles**

254 Magnetic Spinels- Synthesis, Properties and Applications

applications of Fe3O4 nanoparticles.

particles [11].

by natural iron sand.

Generally, the study of Fe3O4 nanoparticles has attracted attention for several disciplines. Physicists investigate the physical properties and propose the theories of the Fe3O4 nanopar‐ ticles. Chemists focus on the synthesis methods of the Fe3O4 nanoparticles. Biologists concen‐ trate on developing Fe3O4 nanoparticles in bioapplications. Meanwhile, engineers study Fe3O4 nanoparticles with a focus on technological applications. Therefore, there is a great deal of interest in Fe3O4 nanoparticles by scientists and engineers, starting from synthesis and characterizations and followed by their many application potentials.

Fe3O4 ferrofluids or magnetic fluids are colloidal suspension containing magnetite nanoparti‐ cles in single domain, which are dispersed in a liquid carrier. Ferrofluids have unique charac‐ teristics because they have both the liquid and magnetic properties. In the ferrofluids, the liquid carriers can be polar or nonpolar medium. In order to prevent agglomerations, the Fe3O4 nanoparticles should be layered or covered by an appropriate material as a shell [19]. In the last decades, it has been reported that the Fe3O4 ferrofluids can be applied in various fields, that is, optical grating, optical switching [20], thermoelectric conversion [21], quenching media [22], sensor [23], controlling lubricant migration [24], biomedical applications [25], and so forth.

Another application for the magnetite spinel is in ferrogels, namely a hydrogel containing magnetic ferrous particles. The Fe3O4‐ferrogel can be defined as a polymer (nano)composites with polyvinyl alcohol (PVA) and Fe3O4 as the matrix and filler, respectively, cross‐linked together with a certain amount of liquid (water). Ferrogels can be physically or chemically cross‐linked without any chemical reaction between them. The physical properties of ferrogels depend on the preparation conditions, such as polymer and solvent ratio, particle size, distribution and concentration of the magnetic material. By combining the elastic medium and the magnetic properties from the particles, the elastic behavior of ferrogel can be manipulated by external magnetic field. This leads to a great opportunity for various advanced applications, namely soft actuators and artificial muscles [26].

In the development of magneto‐elastic materials, Li et al. [27] have prepared a polymeric gel in variety of shapes depending on the temperature. Furthermore, Zrínyi et al. [28] have developed a magneto‐sensitive gel with Fe3O4 as the magnetic filler. Ramanujan and Lao [26] have successfully produced magnetic gel consisting of PVA and micro‐sized Fe3O4 particles for artificial tissues. They have demonstrated an elastic behavior of the gel controlled by an external magnetic field.

## **4. Spinel MgAl2O4 grown as a thick coating and fillerreinforcer in composites**

Compared to the conventional aluminum alloys, aluminum metal matrix composites (Al‐ MMC) have been used in many kinds of applications (e.g. automobiles) due to their high specific strength and Young's modulus, improved yield and creep strengths, light weight, and excellent properties of wear resistance. The important parameter for improving their perform‐ ances lies in the structure and bonding properties between the matrix and fillers. Some materials in any form of particles have been widely used as the particle‐matrix interfaces in order to improve the wettability [29], namely MgO, SiC, and Al2O3 in the micro‐ or nanostruc‐ tures. It has been noted that the interface, particles size, and its distribution as well as the concentration (volume fraction) play a basic role in the modification of Al‐MMC.

A typical way to improve the mechanical properties of Al‐MMC is by additional reinforcement or interfacial modification. Recently, the spinel MgAl2O4 has been applied for this purpose. In the last decade, MgAl2O4 has been produced by several methods such as solid‐state reaction [30], high‐energy ball milling [31], polymerization method [32], pyrolysis [33], sol‐gel proc‐ esses [34, 35], and coprecipitation techniques [36–38]. The shape and size of the spinel particles vary depending on the preparation techniques [39, 40]. This spinel can be prepared in the form of ceramic foams [41, 42], whiskers [43–45], thin films [46, 47], nanoparticles [37, 48, 49], and also layer of particle bonding in composite materials [50]. For the Al/SiC composites, in which Al acts as the matrix and SiC acts as the filler reinforcement, Al4C3 phase can be formed at the interface between matrix and fillers. This phase weakens the mechanical strength of the composite. Up to now, several experimental techniques have been proposed to prevent this problem including adding excessive Si into the matrix and introducing a thin coating layer on the SiC particle surface. The formation of interlayer phase is believed to be a stable interfacial bonding that improves the mechanical properties of the composite. However, the detail of this mechanism remains unclear.

Particulate‐reinforced composites with micron‐sized filler of various materials are the most commonly used composites in daily life. Particles are typically introduced to enhance the matrix elastic modulus and yield strength. It has been shown that the novel properties of composites are improved by adding nano‐scaled particles as the fillers [51]. The key role of mechanical strength in composites is the interface‐bonding quality between matrix and filler materials. The nano‐sized particles have much greater total area compared to particles with larger size at the same volume. The interface interaction can therefore be extended by reducing particle size of fillers down to nanometer scale.

Meanwhile, synthesis of particles in nanometer size of various materials is still an active area of research. Because of their very small size, nanoparticles tend to be aggregate to each other to form larger size particles [52]. As nanometer‐scaled filler, the particles have to be kept separate in the matrix. In this context, an effort to maintain nanoparticles as fillers from agglomeration was achieved by the introduction of a surfactant. Tetra‐methyl‐ammonium‐ hydroxide (TMAH) could be used as surfactant in the process of filler‐matrix‐mixing media. Aluminum was chosen as matrix, while MgO and MgAl2O4 were used as nano‐fillers. In the case of Al/SiC composites, a thin‐layer MgAl2O4 is coated onto the SiC filler to enhance matrix‐ filler interfacial bonding leading to increased mechanical strength. The SiC fillers were coated by spinel MgAl2O4 by electroless plating, soaking SiC particulate powder in solutions con‐ taining Mg and Al ions before heating process to grow spinel phase on the surface.

## **5. Preparation of Fe3O4 nanoparticles from iron sands**

by external magnetic field. This leads to a great opportunity for various advanced applications,

In the development of magneto‐elastic materials, Li et al. [27] have prepared a polymeric gel in variety of shapes depending on the temperature. Furthermore, Zrínyi et al. [28] have developed a magneto‐sensitive gel with Fe3O4 as the magnetic filler. Ramanujan and Lao [26] have successfully produced magnetic gel consisting of PVA and micro‐sized Fe3O4 particles for artificial tissues. They have demonstrated an elastic behavior of the gel controlled by an

Compared to the conventional aluminum alloys, aluminum metal matrix composites (Al‐ MMC) have been used in many kinds of applications (e.g. automobiles) due to their high specific strength and Young's modulus, improved yield and creep strengths, light weight, and excellent properties of wear resistance. The important parameter for improving their perform‐ ances lies in the structure and bonding properties between the matrix and fillers. Some materials in any form of particles have been widely used as the particle‐matrix interfaces in order to improve the wettability [29], namely MgO, SiC, and Al2O3 in the micro‐ or nanostruc‐ tures. It has been noted that the interface, particles size, and its distribution as well as the

A typical way to improve the mechanical properties of Al‐MMC is by additional reinforcement or interfacial modification. Recently, the spinel MgAl2O4 has been applied for this purpose. In the last decade, MgAl2O4 has been produced by several methods such as solid‐state reaction [30], high‐energy ball milling [31], polymerization method [32], pyrolysis [33], sol‐gel proc‐ esses [34, 35], and coprecipitation techniques [36–38]. The shape and size of the spinel particles vary depending on the preparation techniques [39, 40]. This spinel can be prepared in the form of ceramic foams [41, 42], whiskers [43–45], thin films [46, 47], nanoparticles [37, 48, 49], and also layer of particle bonding in composite materials [50]. For the Al/SiC composites, in which Al acts as the matrix and SiC acts as the filler reinforcement, Al4C3 phase can be formed at the interface between matrix and fillers. This phase weakens the mechanical strength of the composite. Up to now, several experimental techniques have been proposed to prevent this problem including adding excessive Si into the matrix and introducing a thin coating layer on the SiC particle surface. The formation of interlayer phase is believed to be a stable interfacial bonding that improves the mechanical properties of the composite. However, the detail of this

Particulate‐reinforced composites with micron‐sized filler of various materials are the most commonly used composites in daily life. Particles are typically introduced to enhance the matrix elastic modulus and yield strength. It has been shown that the novel properties of composites are improved by adding nano‐scaled particles as the fillers [51]. The key role of mechanical strength in composites is the interface‐bonding quality between matrix and filler

**4. Spinel MgAl2O4 grown as a thick coating and fillerreinforcer in**

concentration (volume fraction) play a basic role in the modification of Al‐MMC.

namely soft actuators and artificial muscles [26].

256 Magnetic Spinels- Synthesis, Properties and Applications

external magnetic field.

mechanism remains unclear.

**composites**

The detail of synthesis method for the preparation of Fe3O4 nanoparticles was described in our previous works [53–57]. The iron sand was extracted by a magnetic separator to obtain dominantly Fe3O4 powders. Using a magnetic stirrer, the hydrochloric acid (HCl) was used to dissolve the Fe3O4 powder to produce FeCl3 and FeCl2 solutions at room temperature as written in Eq. (1).

$$\text{Fe}\_3\text{O}\_4 + 8\text{HCl} \rightarrow 2\text{FeCl}\_3 + \text{FeCl}\_2 + 4\text{H}\_2\text{O} \tag{1}$$

The solution of FeCl3 and FeCl2 was then reacted by dropping slowly ammonium hydroxide (NH4OH) to produce black precipitate at room temperature. The reaction mechanism for producing Fe3O4 particles is described as Eq. (2) [54]

$$\text{2FeCl}\_3 + \text{FeCl}\_2 + 8\text{NH}\_4\text{OH} \rightarrow \text{Fe}\_3\text{O}\_4 + 8\text{NH}\_4\text{Cl} + 4\text{H}\_2\text{O} \tag{2}$$

In order to synthesize Mn*x*Fe3‐*x*O4 nanoparticles discussed later and investigate its magnetic properties, the variation of Mn concentration from MnCl2 solution was also prepared to form Mn*x*Fe3‐*x*O4 (0 <*x* ≤ 1). The dissolved MnCl2 solution was mixed using a magnetic stirrer followed by a drop‐wise addition of NH4OH to obtain the precipitate of Fe3O4 and Mn*x*Fe3‐*x*O4. The magnetic moment of Mn2+ ion is 25% higher than that of Fe2+ ion. Therefore, introducing Mn2+ into Fe3O4 to form Mn*x*Fe3‐*x*O4 structure will theoretically enhance the magnetization of Fe3O4 [58]. A washing process of the precipitate was then carried out for several times using distilled water until achieving normal pH condition. Finally, the Fe3O4 powders were prepared by drying process at 100°C for 1 h. For the preparation of Fe3O4 magnetic fluid, tetra‐methyl‐ ammonium‐hydroxide (TMAH), as the stabilizing layers of the particles, was added in the precipitate. In the present investigation, Fe3O4 ferrogels were also synthesized using freezing‐ thawing method as described in Refs. [57, 59].

To study the crystal structure, the particle size and its distribution as well as form and structural factor, and the magnetic properties of the sample, the Fe3O4 powders were then characterized by means of X‐ray diffractometry (XRD), small‐angle neutron scattering (SANS) spectrometer and superconducting quantum interference device (SQUID) magnetometer, respectively. The XRD measurements were carried out in the 2*θ* range from 20° to 70° using Cu‐Kα radiation. A 36‐m SANS spectrometer was used to investigate the primary and secondary particles as well as the fractal structure of Fe3O4 nanoparticles. The detailed SANS measurements have been given elsewhere [54, 60]. The two‐lognormal analysis was also performed to investigate the clustering effect of the magnetic fluid and hydrogel [55, 57].

## **6. Preparation of spinel MgAl2O4 as filler reinforcers**

Magnesium oxide (MgO) and magnesium‐aluminum oxide (MgAl2O4) as fillers were synthe‐ sized by employing coprecipitation process. Mg (Aldrich, 99.9%) and Al (Aldrich, 99.9%) were dissolved in HCl (12.63 molar) with stoichiometric molar fraction. The NH4OH (6.5 molar) was used as precipitating agent to produce Mg(OH)2 and MgAl2O4 after filtering, washing process with distilled water for several times, and finally drying at 100°C. The resulting powders were then heated at 500°C for 1 h to convert them into MgO and MgAl2O4 phases. The powders were checked with XRD and transmission electron microscope (TEM) to explore the phase purity, crystallite size, particle shape, and size.

The MgO and MgAl2O4 powders were mixed with TMAH as surfactant to form a colloid system, consisting of individual MgO or MgAl2O4 particles which were homogeneously dispersed. To fabricate composites, Al powders as matrix were added into the colloid system and followed by thorough stirring and grinding to achieve homogeneity. The mixed powders were dried at 100°C for 3 h and then pressed into pellets having a diameter of 1 cm. During pelletization, the powders were compacted using a force of 1.5 kN applied for 15 min in an isostatic die with zinc stearate on the inner disc. Before sintering process at 500°C for 1 h, pellets were pre‐sintered at, respectively, 200 and 400°C for 20 min each. All heat treatments were conducted in a furnace with a controlled atmosphere of low vacuum (∼10‐3 atm).

The samples produced consist of Al/MgO and Al/MgAl2O4 composites with various volume fractions (% vol) of fillers covering 10, 20, and 30%. We have also prepared the corresponding samples fabricated without TMAH for comparison. To study the effects of surfactant addition during the process on the fabricated composites, measurements of densities, porosity, elastic modulus by compressive test, and microhardness (Vicker's hardness number, VHN) of the samples were carried out.

Furthermore, to prepare the Al/SiC composites, the aluminum (PA) powders and particles of SiC ceramics (220 mesh) were employed as starting materials. The SiC reinforcement particles were cleaned by ultrasonic cleaner in alcohol (90%), and then dried in oven at 100°C. The SiC particles were then soaked in an electrolyte media of 40 ml HNO3 containing Al and Mg ions as the part of the electroless‐plating mechanism. The process of metal oxide coating was done using a magnetic stirrer at 125°C. The oxidation of SiC particles to grow the MgAl2O4 layer was performed in a furnace at 200°C for 1 h, and continued at 400°C for another 1 h. The electrolyte media for electroless plating can be controllably adjusted to grow various metal oxides on the SiC surface, such as MgO, CuO, Al2O3, and so forth, besides MgAl2O4. The metal oxide‐coated SiC particles were mixed with Al powders in n‐butanol (wet mixing) using magnetic stirrer at 100°C, having SiC volume fraction (%vol) of 10, 20, and 30%. The heating process of the green pellets was carried out in a vacuum (10‐3 Torr, rotary pump) by applying a pre‐sintering at 200°C for 20 min and followed by sintering at 600°C for 1 h.

## **7. Results and discussion**

ammonium‐hydroxide (TMAH), as the stabilizing layers of the particles, was added in the precipitate. In the present investigation, Fe3O4 ferrogels were also synthesized using freezing‐

To study the crystal structure, the particle size and its distribution as well as form and structural factor, and the magnetic properties of the sample, the Fe3O4 powders were then characterized by means of X‐ray diffractometry (XRD), small‐angle neutron scattering (SANS) spectrometer and superconducting quantum interference device (SQUID) magnetometer, respectively. The XRD measurements were carried out in the 2*θ* range from 20° to 70° using Cu‐Kα radiation. A 36‐m SANS spectrometer was used to investigate the primary and secondary particles as well as the fractal structure of Fe3O4 nanoparticles. The detailed SANS measurements have been given elsewhere [54, 60]. The two‐lognormal analysis was also performed to investigate

Magnesium oxide (MgO) and magnesium‐aluminum oxide (MgAl2O4) as fillers were synthe‐ sized by employing coprecipitation process. Mg (Aldrich, 99.9%) and Al (Aldrich, 99.9%) were dissolved in HCl (12.63 molar) with stoichiometric molar fraction. The NH4OH (6.5 molar) was used as precipitating agent to produce Mg(OH)2 and MgAl2O4 after filtering, washing process with distilled water for several times, and finally drying at 100°C. The resulting powders were then heated at 500°C for 1 h to convert them into MgO and MgAl2O4 phases. The powders were checked with XRD and transmission electron microscope (TEM) to explore the phase purity,

The MgO and MgAl2O4 powders were mixed with TMAH as surfactant to form a colloid system, consisting of individual MgO or MgAl2O4 particles which were homogeneously dispersed. To fabricate composites, Al powders as matrix were added into the colloid system and followed by thorough stirring and grinding to achieve homogeneity. The mixed powders were dried at 100°C for 3 h and then pressed into pellets having a diameter of 1 cm. During pelletization, the powders were compacted using a force of 1.5 kN applied for 15 min in an isostatic die with zinc stearate on the inner disc. Before sintering process at 500°C for 1 h, pellets were pre‐sintered at, respectively, 200 and 400°C for 20 min each. All heat treatments were

The samples produced consist of Al/MgO and Al/MgAl2O4 composites with various volume fractions (% vol) of fillers covering 10, 20, and 30%. We have also prepared the corresponding samples fabricated without TMAH for comparison. To study the effects of surfactant addition during the process on the fabricated composites, measurements of densities, porosity, elastic modulus by compressive test, and microhardness (Vicker's hardness number, VHN) of the

Furthermore, to prepare the Al/SiC composites, the aluminum (PA) powders and particles of SiC ceramics (220 mesh) were employed as starting materials. The SiC reinforcement particles

conducted in a furnace with a controlled atmosphere of low vacuum (∼10‐3 atm).

thawing method as described in Refs. [57, 59].

258 Magnetic Spinels- Synthesis, Properties and Applications

crystallite size, particle shape, and size.

samples were carried out.

the clustering effect of the magnetic fluid and hydrogel [55, 57].

**6. Preparation of spinel MgAl2O4 as filler reinforcers**

## **7.1. Fe3O4 nanoparticles: magnetic properties and applications**

The XRD pattern of Fe3O4 nano‐powders is shown in **Figure 1**. All peaks in the pattern show a single phase of spinel structure corresponding to the crystal structure of Fe3O4 with PDF No. 19‐0629 without any impurity. Based on the Rietveld analysis, the sample has lattice parameter *a* = *b* = *c* of approximately 8.377 Å.

**Figure 1.** The XRD pattern of Fe3O4 nano‐powders (using Cu‐Kα radiation).

The SANS pattern of the Fe3O4 nano‐powders is presented in **Figure 2**. The SANS data were analyzed using a lognormal spherical model as a form factor *P*(*R*) and mass fractal model as a structure factor *S*(*q*), following Eqs. (3) and (4), respectively,

(3)

**Figure 2.** SANS data of Fe3O4 nano‐powders.

with *R*0 and σ representing the radius of the distribution and standard deviation, respectively [61],

$$S(q) = 1 + \frac{D\Gamma\left(D - 1\right)}{\left(qR\right)^D \left[1 + 1/\left(q^2 \xi\right)^2\right]^{(D-1)/2}} \sin\left[\left(D - 1\right)\tan^{-1}(q\xi)\right] \tag{4}$$

where *q*, Γ, *D*, and *ξ* represent the scattering vector, gamma function, fractal dimension, and cut‐off distance, respectively [62]. The SANS curve was fitted globally using two lognormal spherical model as form factor combining with mass fractal model as structure factor regarding Eq. (5) [54]

$$I(q) \approx \bigcap\_{0}^{\alpha} N\_1(R\_1) F\_N^2(q, R\_1) dR\_1 + \int\_0^{\alpha} N\_2\left(R\_2\right) F\_N^2(q, R\_2) dR\_2 S\left(q, \xi, D, R\_2\right) \tag{5}$$

Here, *I* is the density, *N* is the number density of particles, *R*<sup>1</sup> is the primary particles, *R*2 is the secondary particles or clusters, and *F* is the scattering amplitude.

Based on the analysis using lognormal and mass fractal models, the Fe3O4 nano‐powders have hierarchical nanostructure with the primary particles of 3.8 nm as a building block constructing secondary particles as clusters of 9.3 nm. The clusters of Fe3O4 nanoparticles have a fractal structure in three‐dimension with fractal dimension of 2.9. The SANS data analysis of magnetic nanoparticles coincides with the image produced by high‐resolution transmission electron microscopy (HRTEM) as shown in the previous work [54]. The details of the analysis of SANS data of the Fe3O4 nanoparticles were presented in the previous work [54].

( )

260 Magnetic Spinels- Synthesis, Properties and Applications

**Figure 2.** SANS data of Fe3O4 nano‐powders.

( ) ( ) ( )

0 0

secondary particles or clusters, and *F* is the scattering amplitude.

¥ ¥

<sup>=</sup> *<sup>D</sup> <sup>D</sup> D D*

*qR q*

[61],

Eq. (5) [54]

s

2

è ø

s

<sup>1</sup> ln ( / ) exp <sup>2</sup> <sup>2</sup> æ ö <sup>=</sup> ç ÷

with *R*0 and σ representing the radius of the distribution and standard deviation, respectively

( 1)/2 2 2

1 sin 1 tan ( )


where *q*, Γ, *D*, and *ξ* represent the scattering vector, gamma function, fractal dimension, and cut‐off distance, respectively [62]. The SANS curve was fitted globally using two lognormal spherical model as form factor combining with mass fractal model as structure factor regarding

> ( ) ( ) ( ) ( ) ( ) ( ) 2 2 11 1 1 2 2 2 2 2

Here, *I* is the density, *N* is the number density of particles, *R*<sup>1</sup> is the primary particles, *R*2 is the

» ò ò *N N I q N R F q R dR N R F q R dR S q D R*

ë û é ù <sup>+</sup> ë û

, + , ,, ,

+ - é ù

1

x

*S q D q*

1 1/( )

G -

( ) <sup>1</sup>


(4)

x

(5)

x

*R R P R*

 p*R*

0 2

(3)

The magnetic properties of the Fe3O4 nanoparticles as well as M versus H variations and zero‐ field‐cooling (ZFC) curves are presented in **Figures 3** and **4**. The M versus H curve of the Fe3O4 nanoparticles was collected at room temperature by sweeping the magnetic field from ‐5 to 5 T. The Fe3O4 nanoparticles have saturated magnetization of 37.1 emu/g. Regarding **Figure 3**, it is clear that the M‐H curve of sample has S‐shape with nearly zero coercivity field, indicating the superparamagnetic behavior at room temperature. This result is consistent with the recent paper by Abboud et al. [63] where the magnetization was almost zero in the absence of external magnetic field. The superparamagnetic phenomenon in the magnetic nanoparticles has also been observed in the Fe3O4@SiO2 core‐shell composites [64], the dispersed Fe3O4 in the polymer matrix [65], and Fe3O4‐LiMo3Se3 [66]. Moreover, the zero‐field‐cooling (ZFC) meas‐ urement was also carried out to investigate the superparamagnetic phenomenon of the sample as shown in **Figure 4**. Based on the results in **Figure 4**, it is evident that the sample has a maximum peak of magnetic blocking temperature, *T*B, at 243 K. Theoretically, *T*B depends on the particle size and shape so the increase of nanoparticle's volume will increase the *T*B value [67].

**Figure 3.** M‐H curve of Fe3O4 nanoparticles at room temperature.

**Figure 4.** Zero‐field‐cooling (ZFC) curve of Fe3O4 nanoparticles (*T*B = 243 K).

Dutta et al. [68] have reported that Fe3O4 nanoparticles with the particles size ranging from 4 to 12 nm have blocking temperature lower than 100 K, which is not similar to the result of this work. The difference of these results can be explained in a sense of the presence of clusters or aggregations phenomena in the present work. In Ref. [68], the samples were constructed with primary particles in relatively homogeneous samples without any clusters or aggregations. On the other hand, the sample in this work consists of hierarchical nanostructures of primary particles forming secondary particles or clusters with fractal dimension in three‐dimensions. Despite the particles size and anisotropy constant, the clusters of magnetic nanoparticles give an effect on the blocking temperature. Theoretically, at magnetic‐blocking temperature, the thermal energy of particles is comparable with the anisotropy energy barrier. In bulk, the Fe3O4 is ferrimagnetic generated by net magnetic moments at tetrahedral and octahedral sites. The effect of clusters on magnetic properties was also documented by other researchers [69, 70].

In the magnetic nanoparticles, there is the so‐called blocking temperature or the energy barrier that can be obtained from ZFC and field‐cooling (FC) magnetization curves. With increasing temperature, a curve peak should appear in the ZFC measurement. This peak temperature can be considered as the average *T*B in the magnetic material. On the other hand, the increase of magnetic field should decrease the barrier energy, and results in a shift of the blocking temperature to lower temperatures. A typical of nanocrystalline Fe3O4 has shown superpara‐ magnetic behavior (**Figure 3**). The superparamagnetism is related to the random fluctuation of the magnetization of a single domain when the thermal energy overcomes the anisotropy energy barrier. Below the *T*B, the magnetization of each domain of nanoparticles is oriented parallel to a certain crystallographic direction or the easy axis with minimum energy. Conse‐ quently, the magnetization is blocked in the nanoparticles. Above *T*B, the thermal energy can overcome the anisotropy energy (barrier energy) and the magnetization starts fluctuating and the magnetic susceptibility follows modified Curie‐law behavior at higher temperatures [71].

The effect of Mn2+ substitution on the magnetic properties of Mn*x*Fe3‐*x*O4 nanoparticles has been intensively studied recently [54, 72]. It has been found that the lattice parameters and crystal volume of the Mn*x*Fe3‐*x*O4 increase with the increase of Mn content. This is due to the fact that the ionic radius of Fe2+ is smaller than that of Mn2+ substituted in the spinel structure. Fe3O4 without any substitution has a cubic spinel structure, where the Fe2+ and Fe3+ ions occupy the tetrahedral (A) and octahedral (B) sites represented by (Fe3+)A(Fe3+Fe2+)BO2‐ 4. The Fe3O4 is one example of the cubic inverse spinels, in which there is a mixed valence of Fe ions on the octahedral sublattice. In the Mn*x*Fe3‐*x*O4, the cationic distribution can be written as follows [54]:

$$\left(\mathrm{Mn}\_{\mathrm{x}}^{2+}\mathrm{Fe}\_{\mathrm{l}-\mathrm{x}}^{3+}\right)\_{\mathrm{A}}\left[\mathrm{Fe}\_{\mathrm{l}-\mathrm{x}}^{2+}\mathrm{Fe}\_{\mathrm{l}+\mathrm{x}}^{3+}\right]\_{\mathrm{B}}\tag{6}$$

The Mn2+ ions occupy the A sites of the spinel structure. A further analysis of the SANS data of the Mn*x*Fe3‐*x*O4 nanoparticles from *x* = 0 to 1 using the lognormal function has found that the particle size of the Mn*x*Fe3‐*x*O4 depends on the Mn content. In this study, the increase of Mn content results in the smaller size of the nanoparticles. Moreover, the nanoparticles of Mn*x*Fe3‐ *<sup>x</sup>*O4 tend to become larger aggregates or clusters. The Mn content, particle size, and particle clustering are important factors influencing the magnetic properties of nanomaterials.

**Figure 4.** Zero‐field‐cooling (ZFC) curve of Fe3O4 nanoparticles (*T*B = 243 K).

262 Magnetic Spinels- Synthesis, Properties and Applications

Dutta et al. [68] have reported that Fe3O4 nanoparticles with the particles size ranging from 4 to 12 nm have blocking temperature lower than 100 K, which is not similar to the result of this work. The difference of these results can be explained in a sense of the presence of clusters or aggregations phenomena in the present work. In Ref. [68], the samples were constructed with primary particles in relatively homogeneous samples without any clusters or aggregations. On the other hand, the sample in this work consists of hierarchical nanostructures of primary particles forming secondary particles or clusters with fractal dimension in three‐dimensions. Despite the particles size and anisotropy constant, the clusters of magnetic nanoparticles give an effect on the blocking temperature. Theoretically, at magnetic‐blocking temperature, the thermal energy of particles is comparable with the anisotropy energy barrier. In bulk, the Fe3O4 is ferrimagnetic generated by net magnetic moments at tetrahedral and octahedral sites. The effect of clusters on magnetic properties was also documented by other researchers [69, 70].

In the magnetic nanoparticles, there is the so‐called blocking temperature or the energy barrier that can be obtained from ZFC and field‐cooling (FC) magnetization curves. With increasing temperature, a curve peak should appear in the ZFC measurement. This peak temperature can be considered as the average *T*B in the magnetic material. On the other hand, the increase of magnetic field should decrease the barrier energy, and results in a shift of the blocking temperature to lower temperatures. A typical of nanocrystalline Fe3O4 has shown superpara‐ magnetic behavior (**Figure 3**). The superparamagnetism is related to the random fluctuation of the magnetization of a single domain when the thermal energy overcomes the anisotropy energy barrier. Below the *T*B, the magnetization of each domain of nanoparticles is oriented parallel to a certain crystallographic direction or the easy axis with minimum energy. Conse‐ quently, the magnetization is blocked in the nanoparticles. Above *T*B, the thermal energy can

Taufiq et al. [55] have successfully synthesized Fe3O4 magnetic fluid with a chain‐like spinel structure using coprecipitation method. These results show that the Fe3O4 ferrofluid has a primary particle size of about 7.6 nm and the size of fractal aggregates of about 45 nm con‐ structing the chain‐like structure. Even though there were chain aggregates, a homogene‐ ous particle distribution with low polydispersity value of about 0.4 was found, which is lower than that in the former paper [55]. It is reported that both nanoparticles of Fe3O4 and Mn*x*Fe3‐*x*O4 (*x* up to 1) exhibit superparamagnetic behavior and their saturation magnetiza‐ tion decreases with increasing Mn content [54, 72].

In the research of magnetic hydrogels, Fe3O4 ferrogels have been investigated by Sunaryono et al. [57]. It has been reported that the magneto‐elastic properties of the ferrogels were strongly affected by the preparation technique, the ratio of the magnetic particle and the polymer, and magnetic particle content as well as the particle clustering and distribution. It has also been shown that the magneto‐elasticity of the ferrogel tends to decrease with the increasing Fe3O4 content from 2.5 to 15% [56, 57]. Further analysis of the SAXS data using the two‐lognormal distribution function showed that Fe3O4 in the hydrogel has primary particles of about 3 nm with an average particle distance of about 18 nm. Higher Fe3O4 concentration in the hydrogel leads to the increase of Fe3O4 cluster size. Due to the surface effect, it influences the magneti‐ zation value of the ferrogel. Consequently, the saturation magnetization drops with decreasing particle size and the cluster of Fe3O4 nanoparticles in the hydrogel. It has been confirmed that Fe3O4 ferrogel exhibits a superparamagnetic behavior at room temperature which is a crucial parameter for biomedical applications.

## **7.2. Improved mechanical properties of aluminum matrix composites with surfactant‐ coated MgAl2O4 nanofillers and MgAl2O4‐coated SiC particulates**

**Figure 5** presents the XRD spectra (using Cu‐Kα radiation, *λ* = 1.54 angstrom) of MgO and MgAl2O4 particles as results of coprecipitation process. The powders seem to be single phase as can be seen from the spectra, where no other additional diffraction peak is observed, except those associated with Miller indexes belonging to the MgO (periclase phase) and MgAl2O4 (spinel phase) as specified in **Figure 5**. Besides, the diffraction peaks are much broadened, featuring that the crystal is in nanometer size. If we take the full‐width at half maximum (FWHM) of diffraction peaks, by using Scherrer formula [73] and taking the apparatus‐ broadening correction into account, we arrive at crystallite sizes of 17.5 ± 2.3 and 6.3 ± 0.7 nm for MgO and MgAl2O4 powders, respectively.

**Figure 5.** XRD spectra of MgO and MgAl2O4.

**Figure 6.** TEM images of (a) MgO and (b) MgAl2O4.

It is interesting to correlate the above crystallite size to the particle size of powders from TEM images depicted in **Figure 6**. The particle size for both powders, MgO in **Figure 6a** and MgAl2O4 in **Figure 6b**, is seen to be between 50 and less than 100 nm, confirming that the powders can be classified as nanoparticles. The shape of particles is observed to be varied from spherical, square, square with rounded edge, to semi‐rod. Comparing the crystal and the particle sizes, one can consider that both MgO and MgAl2O4 powders contain secondary particles, where each particle is constituted by several crystals or grains.

**7.2. Improved mechanical properties of aluminum matrix composites with surfactant‐**

**Figure 5** presents the XRD spectra (using Cu‐Kα radiation, *λ* = 1.54 angstrom) of MgO and MgAl2O4 particles as results of coprecipitation process. The powders seem to be single phase as can be seen from the spectra, where no other additional diffraction peak is observed, except those associated with Miller indexes belonging to the MgO (periclase phase) and MgAl2O4 (spinel phase) as specified in **Figure 5**. Besides, the diffraction peaks are much broadened, featuring that the crystal is in nanometer size. If we take the full‐width at half maximum (FWHM) of diffraction peaks, by using Scherrer formula [73] and taking the apparatus‐ broadening correction into account, we arrive at crystallite sizes of 17.5 ± 2.3 and 6.3 ± 0.7 nm

**coated MgAl2O4 nanofillers and MgAl2O4‐coated SiC particulates**

for MgO and MgAl2O4 powders, respectively.

264 Magnetic Spinels- Synthesis, Properties and Applications

**Figure 5.** XRD spectra of MgO and MgAl2O4.

**Figure 6.** TEM images of (a) MgO and (b) MgAl2O4.

**Figure 7.** Porosity of sintered Al/MgO and Al/MgAl2O4 composites with fillers specified.

**Figure 8.** Elastic moduli of sintered Al/MgO and Al/MgAl2O4 composites with fillers specified.

To discuss further, the physical properties of composites, the density of pellets in the green state (after compaction) and after sintering significantly increases with increasing volume fraction of fillers and the use of surfactant (coated fillers by surfactant). The enhanced density has led to the lowering porosity in the pellets. By applying the formula to define porosity [52]: *ρ* = 1 –*ρ*s/*ρ*<sup>t</sup> , where *ρ*s and *ρ*<sup>t</sup> stand for, respectively, sintered and theoretical densities, the porosity of samples is exhibited in **Figure 7**. The increase of pellet densification has directly affected the mechanical properties, as reflected by the results of compressive tests (yielding elastic modulus) and hardness (VHN) measurements (see **Figures 8** and **9**).

**Figure 9.** Micro‐hardness of sintered Al/MgO and Al/MgAl2O4 composites with fillers specified.

It is worth noting that the role of surfactant is to enhance the physical and therefore mechanical properties of these Al/MgO and Al/MgAl2O4 nanocomposites. The TMAH as a mixing media of matrix and filler is an ionic molecule which will attach on the surface of nanoparticles. A metal‐oxide nanoparticle is usually deficient of oxygen ion [74], hence its surface is positively charged. The negatively charged end of TMAH molecule is therefore attracted by metal‐oxide surface, while the opposite end is dangled. It may further create the situation where one may consider that MgO or MgAl2O4 nanoparticle is "coated" by TMAH. This in turn avoids nanoparticles from aggregation to each other in the matrix. So, the filler is still in nanometer scale to maintain its larger surface area, especially compared to the agglomerated fillers having the same volume fraction in the composites fabricated without TMAH. The larger filler surface will lead to stronger interface bonding between matrix and filler.

The next attractive point to address is that the surfactant‐coated MgO filler has led to consid‐ erably enhanced elastic modulus (∼60%) and microhardness (∼40%) of the Al‐matrix nano‐ composites (**Figures 8** and **9**). The periclase phase of MgO has simpler crystal structure than the spinel MgAl2O4 one [75, 76]. The former phase is transformed from Mg(OH)2 at 500°C, whereas the latter one is already stable from 100°C as a result of coprecipitation process. Besides particles consolidation during the sintering at 500°C, the MgO filler will be possibly able to react further in a limited amount with Al matrix to form spinel phase on the matrix‐ filler interface, creating a strong bonding. Such kind of process, except consolidation, may not occur in the Al/MgAl2O4 composite, because of stable MgAl2O4.

**Figure 10.** Elastic modulus of Al/SiC composites depending on filler content, using SiC coated by metal oxides as specified.

Going to further study on Al/SiC composites, it can be seen in **Figure 10** that a very clear distinction exists between SiC particles without coating and with coated metal oxides on the surface based on the modulus elasticity. The entire filler volume fraction (%vol) of SiC in Al‐ SiC composites using coatings has higher value than that without coating. The experimental results show that the modulus of elasticity of the composites using metal‐oxide coating approaches the theoretical prediction [77], estimating a limit of lower and upper bounds. Based on the stress analysis, the composite having filler without metal oxide coating, for all low‐ volume fraction of reinforcement, the modulus value lies in the outside of the upper and lower bounds, which indicates weak interaction between matrix and filler. This is much different from composites with filler using metal oxides, where for the entire volume fraction of reinforcement, the elastic modulus lies inside the region of the upper and lower bounds, implying that the compactibility between matrix and filler, and reinforcements are therefore going well. One can see in **Figure 10** that coating with MgAl2O4 onto the SiC surface has produced the superior Al/SiC composites among the prepared samples, followed by Al2O3 and MgO. The three metal oxides coating materials mentioned seem to be compatible in facilitating and therefore enhancing interfacial bonding between Al matrix and SiC filler in the composites.

### **8. Conclusions**

has led to the lowering porosity in the pellets. By applying the formula to define porosity [52]:

porosity of samples is exhibited in **Figure 7**. The increase of pellet densification has directly affected the mechanical properties, as reflected by the results of compressive tests (yielding

elastic modulus) and hardness (VHN) measurements (see **Figures 8** and **9**).

**Figure 9.** Micro‐hardness of sintered Al/MgO and Al/MgAl2O4 composites with fillers specified.

will lead to stronger interface bonding between matrix and filler.

occur in the Al/MgAl2O4 composite, because of stable MgAl2O4.

It is worth noting that the role of surfactant is to enhance the physical and therefore mechanical properties of these Al/MgO and Al/MgAl2O4 nanocomposites. The TMAH as a mixing media of matrix and filler is an ionic molecule which will attach on the surface of nanoparticles. A metal‐oxide nanoparticle is usually deficient of oxygen ion [74], hence its surface is positively charged. The negatively charged end of TMAH molecule is therefore attracted by metal‐oxide surface, while the opposite end is dangled. It may further create the situation where one may consider that MgO or MgAl2O4 nanoparticle is "coated" by TMAH. This in turn avoids nanoparticles from aggregation to each other in the matrix. So, the filler is still in nanometer scale to maintain its larger surface area, especially compared to the agglomerated fillers having the same volume fraction in the composites fabricated without TMAH. The larger filler surface

The next attractive point to address is that the surfactant‐coated MgO filler has led to consid‐ erably enhanced elastic modulus (∼60%) and microhardness (∼40%) of the Al‐matrix nano‐ composites (**Figures 8** and **9**). The periclase phase of MgO has simpler crystal structure than the spinel MgAl2O4 one [75, 76]. The former phase is transformed from Mg(OH)2 at 500°C, whereas the latter one is already stable from 100°C as a result of coprecipitation process. Besides particles consolidation during the sintering at 500°C, the MgO filler will be possibly able to react further in a limited amount with Al matrix to form spinel phase on the matrix‐ filler interface, creating a strong bonding. Such kind of process, except consolidation, may not

stand for, respectively, sintered and theoretical densities, the

*ρ* = 1 –*ρ*s/*ρ*<sup>t</sup>

, where *ρ*s and *ρ*<sup>t</sup>

266 Magnetic Spinels- Synthesis, Properties and Applications


## **Acknowledgements**

This chapter is based on research partially supported by LPPM ITS, Ministry of Research and Technology, and Ministry of Education and Culture, Republic of Indonesia, 2006–2015.

## **Author details**

Malik Anjelh Baqiya1 , Ahmad Taufiq2 , Sunaryono2 , Khuroti Ayun1 , Mochamad Zainuri1 , Suminar Pratapa1 , Triwikantoro1 and Darminto1\*

\*Address all correspondence to: darminto@physics.its.ac.id

1 Advanced Materials Research Group, Department of Physics, Sepuluh Nopember Institute of Technology (ITS), Surabaya, Indonesia

2 Department of Physics, State University of Malang (UM), Malang, Indonesia

## **References**


**3.** The magnetic properties have also been investigated by observing the hysteresis (M‐H) behavior. It is found that Fe3O4 and Mn*x*Fe3‐*x*O4 (0 < *x* < 1) nanoparticles exhibit superpar‐ amagnetic behavior at room temperature which depends on the particle size, doping

**4.** The physical and mechanical properties of Al/MgO and Al/MgAl2O4 nanocomposites have been enhanced by the introduction of TMAH as surfactant during the mixing process of matrix and filler. The surfactant‐coated MgO seems to be a more efficient filler compared to the surfactant‐coated MgAl2O4 for the lower content (<30%). The experiment results show that the elastic modulus of the samples with volume fraction of fillers up to 30% was enhanced by more than 60% and 40%, respectively, for Al/MgO and Al/MgAl2O4. The use of other kind of surfactant needs to be explored to further enhance the mechanical

**5.** A layer of the spinel MgAl2O4 micro‐sized grown on the surface of SiC filler has induced

This chapter is based on research partially supported by LPPM ITS, Ministry of Research and Technology, and Ministry of Education and Culture, Republic of Indonesia, 2006–2015.

, Sunaryono2

1 Advanced Materials Research Group, Department of Physics, Sepuluh Nopember Institute

[1] D. Ghanbari, M. Salavati‐Niasari, M. Ghasemi‐Kooch, J. Ind. Eng. Chem. 20 (2014) 3970–

[2] L.F. Cótica, V.F. Freitas, G.S. Dias, I.A. Santos, S.C. Vendrame, N.M. Khalil, R.M. Mainardes, M. Staruch, and M. Jain, J. Magn. Magn. Mater. 324 (2012) 559–563.

and Darminto1\*

2 Department of Physics, State University of Malang (UM), Malang, Indonesia

, Khuroti Ayun1

, Mochamad Zainuri1

,

the enhancement of mechanical properties of the Al/SiC composites.

, Ahmad Taufiq2

\*Address all correspondence to: darminto@physics.its.ac.id

, Triwikantoro1

of Technology (ITS), Surabaya, Indonesia

content, and particle clustering.

268 Magnetic Spinels- Synthesis, Properties and Applications

properties.

**Acknowledgements**

**Author details**

Suminar Pratapa1

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Malik Anjelh Baqiya1


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#### **Photothermal Conversion Applications of the Transition Metal (Cu, Mn, Co, Cr, and Fe) Oxides with Spinel Structure Photothermal Conversion Applications of the Transition Metal (Cu, Mn, Co, Cr, and Fe) Oxides with Spinel Structure**

Pengjun Ma, Qingfen Geng and Gang Liu Pengjun Ma, Qingfen Geng and Gang Liu

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

http://dx.doi.org/10.5772/67210

#### **Abstract**

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272 Magnetic Spinels- Synthesis, Properties and Applications

67–74.

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133.

The transition metal (Cu, Mn, Co, Cr, and Fe) oxides with spinel structure can be used as solar absorber materials due to their unique properties. Copper-based spinel ceramic pigments have been successfully prepared by sol-gel combustion method at low temperatures. Subsequently, spinel ceramic pigments have been employed to fabricate selective absorber paint coating by spraying-coating. The paint coating showed good spectral selectivity and thermal stability at low-to-mid temperature region. Spinel ceramic films have also been deposited on metal substrates by one dipping/annealing cycle. Spinel phase for single-layer ceramic film could be achieved at low annealing temperatures, and the single-layer ceramic film showed good spectral selectivity and benign thermal stability. Results presented here show that spinel compounds based on transition metal (Cu, Mn, Co, Cr, and Fe) oxides are promising materials for photothermal conversion applications.

**Keywords:** solar absorber coating, sol-gel combustion, spinel ceramic pigment, sol-gel dip-coating, spinel ceramic film, spectral selectivity

## **1. Introduction**

Spectrally selective absorber coatings have attracted more attention because of the ability of its absorber to convert solar radiation into heat. The high photothermal conversion efficiency is usually dependent on selectivity absorption of the coating, which is required to have a high solar absorbance (*α*<sup>s</sup> ) in the solar wavelength range (0.3–2.5 μm) and low

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

thermal emittance (*ε*T) in infrared wavelength ranges (2.5–20 μm) [1]. During the last few decades, the transition metal (Cu, Mn, Co, Cr, and Fe) oxides with spinel structure have attracted significant attention due to their unique properties such as chemical inertness, high corrosion resistance, high mechanical strength, good thermal-shock resistance, and excellent optical and catalytic properties. This makes them very suitable for potential applications ranging from optics, electronics, magnetism, and catalysis to energy conversion and storage [2–4]. Furthermore, the existence of numerous spin-allowed electron transitions between partially filled d-orbitals contributes to their high absorption of radiation across the whole solar radiation spectrum [5]. These features endow these materials with promising application as solar selective absorber in solar-thermal energy conversion systems.

In recent years, several different approaches to the preparation of the spinel ceramic powders have been developed. For example, to prepare the spinel ceramic pigment, techniques of high temperature solid-state reaction, coprecipitation, sol-gel, hydrothermal synthesis, and thermal decomposition have been employed [6–8]. However, it is the major drawback of the above methods that the presence of heterogeneous products and requirement of high temperature and long durations result in tremendous wastage of energy. The sol-gel selfcombustion method contributes to synthesize the better chemical homogeneity, small grain size, and high purity powder, which requires relatively low crystallization temperature [9]. In addition, crystallinity, size, and morphology of the particle, surface area, as well as other particular properties of the particles could be directed by tuning reaction parameters such as kinds of fuel, the mole ratio of fuel and oxidizer ratio, kinds and contents of the solvent, and the annealing temperature [10]. The as-prepared spinel ceramic powders are then utilized as solar-absorbing pigments to fabricate thickness-sensitive spectrally selective (TSSS) paint coatings by a convenient and economical spray-coating technology. Spray-coating technique is quick, easily adaptable to different coating solutions, and complex shapes can be coated. This makes it adoptable for an in-line process with minimum of material waste. These advantages with the spray-coating method suggest that this is the technique to prefer when scaling up the process.

Sol-gel techniques are promising synthesis methods for these spectral selectivity absorber coatings. The optical properties and durability of the spectral selectivity absorber coating can be easily controlled by fine-tuning relevant design parameters such as heating temperature or precursor concentrations in the synthesis process. In light of this, there are many knowledge gaps that need to be filled in the context of technicalities regarding the sol-gel processes and the optical and morphological characteristics of these coatings. The sol-gel processes are a soft chemistry technique where the precursors are generally in the form of a colloidal-based solution that eventually "transforms" into a widespread network of either discrete or continuously linked molecules [11]. Sol-gel techniques facilitate control of coating parameters such as absorber particle size, particle size distribution, homogeneity, chemical composition, and film thickness. The techniques also show good potential for scaling up to an industrial scale [12]. The most significant advantage of sol-gel over other established coating methods is its ability to tailor the microstructure of the deposited film at relatively low temperatures [13].

## **2. Experimental procedures**

thermal emittance (*ε*T) in infrared wavelength ranges (2.5–20 μm) [1]. During the last few decades, the transition metal (Cu, Mn, Co, Cr, and Fe) oxides with spinel structure have attracted significant attention due to their unique properties such as chemical inertness, high corrosion resistance, high mechanical strength, good thermal-shock resistance, and excellent optical and catalytic properties. This makes them very suitable for potential applications ranging from optics, electronics, magnetism, and catalysis to energy conversion and storage [2–4]. Furthermore, the existence of numerous spin-allowed electron transitions between partially filled d-orbitals contributes to their high absorption of radiation across the whole solar radiation spectrum [5]. These features endow these materials with promising applica-

In recent years, several different approaches to the preparation of the spinel ceramic powders have been developed. For example, to prepare the spinel ceramic pigment, techniques of high temperature solid-state reaction, coprecipitation, sol-gel, hydrothermal synthesis, and thermal decomposition have been employed [6–8]. However, it is the major drawback of the above methods that the presence of heterogeneous products and requirement of high temperature and long durations result in tremendous wastage of energy. The sol-gel selfcombustion method contributes to synthesize the better chemical homogeneity, small grain size, and high purity powder, which requires relatively low crystallization temperature [9]. In addition, crystallinity, size, and morphology of the particle, surface area, as well as other particular properties of the particles could be directed by tuning reaction parameters such as kinds of fuel, the mole ratio of fuel and oxidizer ratio, kinds and contents of the solvent, and the annealing temperature [10]. The as-prepared spinel ceramic powders are then utilized as solar-absorbing pigments to fabricate thickness-sensitive spectrally selective (TSSS) paint coatings by a convenient and economical spray-coating technology. Spray-coating technique is quick, easily adaptable to different coating solutions, and complex shapes can be coated. This makes it adoptable for an in-line process with minimum of material waste. These advantages with the spray-coating method suggest that this is the technique to prefer when scaling

Sol-gel techniques are promising synthesis methods for these spectral selectivity absorber coatings. The optical properties and durability of the spectral selectivity absorber coating can be easily controlled by fine-tuning relevant design parameters such as heating temperature or precursor concentrations in the synthesis process. In light of this, there are many knowledge gaps that need to be filled in the context of technicalities regarding the sol-gel processes and the optical and morphological characteristics of these coatings. The sol-gel processes are a soft chemistry technique where the precursors are generally in the form of a colloidal-based solution that eventually "transforms" into a widespread network of either discrete or continuously linked molecules [11]. Sol-gel techniques facilitate control of coating parameters such as absorber particle size, particle size distribution, homogeneity, chemical composition, and film thickness. The techniques also show good potential for scaling up to an industrial scale [12]. The most significant advantage of sol-gel over other established coating methods is its ability to tailor the microstructure of the deposited film at relatively

tion as solar selective absorber in solar-thermal energy conversion systems.

274 Magnetic Spinels- Synthesis, Properties and Applications

up the process.

low temperatures [13].

#### **2.1. Synthesis of spinel ceramic pigments and spectrally selective paint coatings**

Spinel ceramic pigments were synthesized by sol-gel self-combustion technique. Metal nitrate was first dissolved in an adequate amount of ultrapure water with the appropriate molar ratio. An appropriate amount of citric acid was then added into the prepared aqueous solution to chelate metal ion. After stirring for some time, the polyethylene glycol was added to the solution as an esterifying agent, which took part in chelation reaction. The mixture solution was adjusted to pH = 6.0–4.5 by slowly dropping ammonia and successively stirred to obtain a homogeneous solution. The prepared solution was subsequently heated for the adequate period of time to form the xerogel. Then, the xerogel was ignited in the atmosphere and burned in a self-combustion manner with rapid evolution of a large quantity of fume, yielding voluminous powders. Finally, the as-burned powders were annealed at different temperatures to obtain spinel ceramic pigments. Pigment dispersion was first prepared by mixing the pigments with the commercially corresponding binders and solvent in specific proportions and ground in a ball mill to form paint. Ultimately, the paint was sprayed on metal substrate to obtain paint coatings. A diagram for the sample preparation procedures is shown in **Figure 1**.

 **Figure 1.** Spectrally selective paint coatings based on spinel ceramic pigments are fabricated [14].

#### **2.2. Synthesis of spinel ceramic film for spectrally selective coatings**

Metallic precursor sol was prepared by dissolving metal nitrates in absolute ethanol with suitable mole ratio of 1:1. This was followed by the addition of the citric acid as chelating agents. After stirring for a period of time, appropriate amount of polyethylene glycol was added under magnetic stirring. The resulting solution was then used for coating deposited on metallic substrates by soakage method and subsequently heated to obtain xerogel films. The films were then air annealed in an oven at different temperatures.

#### **3. Results and discussion**

#### **3.1. Spectral selectivity paint coating based on Cu1.5Mn1.5O4 spinel ceramic pigments**

Cu1.5Mn1.5O4 spinel ceramic pigments have been prepared by a facile and cost-effective solgel self-combustion method and annealed at the temperature ranging from 500 to 900°C for 1 h [14]. Ceramic pigments are utilized to fabricate the TSSS paint coatings by means of the convenient and practical spray-coating technique, and TSSS paint coatings based on pigments annealed at 700°C show solar absorbance of *α*<sup>s</sup> = 0.914–0.923 and thermal emittance of *ε*<sup>T</sup> = 0.244–0.357, which are calculated from the reflectance spectral shown in **Figure 2**. Furthermore, it is seen from reflectance spectral that the absorption edge is shifted toward longer wavelength, which means that the coating becomes thick and both of solar absorptance and thermal emittance are increased. As can be seen from **Figure 3**, the reflectance spectra show the marginal changes after the accelerated thermal test is carried out. Therefore, TSSS paint coating shows the thermal stability at the temperature of 227°C. Meanwhile, TSSS paint coatings exhibit no observable visual changes, and the performance criterion (PC) values reach the qualified requirement.

 **Figure 2.** Reflectance spectra of paint coatings with different thickness based on spinel ceramic pigments [14].

Photothermal Conversion Applications of the Transition Metal (Cu, Mn, Co, Cr, and Fe) Oxides with Spinel Structure http://dx.doi.org/10.5772/67210 277

 **Figure 3.** Reflectance spectra of paint coating subjected to thermal stability test [14].

under magnetic stirring. The resulting solution was then used for coating deposited on metallic substrates by soakage method and subsequently heated to obtain xerogel films. The films

spinel ceramic pigments have been prepared by a facile and cost-effective sol-

gel self-combustion method and annealed at the temperature ranging from 500 to 900°C for 1 h [14]. Ceramic pigments are utilized to fabricate the TSSS paint coatings by means of the convenient and practical spray-coating technique, and TSSS paint coatings based on pigments annealed at 700°C show solar absorbance of *α*<sup>s</sup> = 0.914–0.923 and thermal emittance of *ε*<sup>T</sup> = 0.244–0.357, which are calculated from the reflectance spectral shown in **Figure 2**. Furthermore, it is seen from reflectance spectral that the absorption edge is shifted toward longer wavelength, which means that the coating becomes thick and both of solar absorptance and thermal emittance are increased. As can be seen from **Figure 3**, the reflectance spectra show the marginal changes after the accelerated thermal test is carried out. Therefore, TSSS paint coating shows the thermal stability at the temperature of 227°C. Meanwhile, TSSS paint coatings exhibit no observable visual changes, and the performance criterion (PC) val-

 **Figure 2.** Reflectance spectra of paint coatings with different thickness based on spinel ceramic pigments [14].

 **spinel ceramic pigments**

were then air annealed in an oven at different temperatures.

**3.1. Spectral selectivity paint coating based on Cu1.5Mn1.5O4**

**3. Results and discussion**

276 Magnetic Spinels- Synthesis, Properties and Applications

ues reach the qualified requirement.

Cu1.5Mn1.5O4

#### **3.2. Spectral selectivity paint coating based on CuCr2 O4 spinel ceramic pigments**

Single-phase CuCr2 O4 spinel crystals are obtained after heat treatment of the as-burned powder at a low temperature (600°C) for 1 h, and the average crystallite size, morphology, and crystallinity of the CuCr2 O4 are greatly influenced by the annealing temperature. It can be seen from SEM images (**Figure 4**) that the as-burned powder has numerous voids and pores embedding into lamella. Increasing the annealing temperature, there are obvious flat face and clear edge appearing on the particles. Particles take on regular octahedron-shaped morphology and perfection of crystals at the annealing temperature of 800°C [15]. Comparison of TSSS paint coating based on metal oxide powder, the as-burned powder, and CuCr2 O4 spinel ceramic pigments as solar absorber pigments shows that TSSS paint coatings based on the spinel ceramic pigment exhibit the relative high solar selective absorption. The SEM representative morphologies of the similar thickness TSSS paint coatings ((a) sample A3 based on CuO and Cr2 O3 , (b) sample B4 based on as-burned pigment, (c) sample C4 based on pigment annealed at 600°C, and (d) sample D2 based on pigments annealed at 800°C) are shown in **Figure 5**. It can be observed that the surface morphologies of all samples exhibit microscale bumps and protrusions, which are brought about particles agglomeration. For the TSSS paint coating, pigment particle agglomeration in the resin can cause uneven distribution and formation of clusters. The corresponding 3D surface profile images of samples are coincident with the surface morphologies shown by the SEM images. Water contact angles exhibited on sample surface can testify the different surface roughness (**Figure 6**). Furthermore, the TSSS paint coatings based on the spinel ceramic pigment show low surface roughness value and good hydrophobicity.

**Figure 4.** FE-SEM images and corresponding (inset) photographs of the powders: (a) the as-burned powder and powders annealed at (b) 600°C, (c) 700°C, (d) 800°C, and (e) 900°C for 1 h [15].

 **Figure 5.** SEM images of paint coatings for (a) the sample A3, (b) the sample B4, (c) the sample C4, and (d) the sample D2 [15].

Photothermal Conversion Applications of the Transition Metal (Cu, Mn, Co, Cr, and Fe) Oxides with Spinel Structure http://dx.doi.org/10.5772/67210 279

 **Figure 6.** (a) 3D surface roughness profiles of samples and (b) images of water droplet on samples [15].

**Figure 4.** FE-SEM images and corresponding (inset) photographs of the powders: (a) the as-burned powder and powders

 **Figure 5.** SEM images of paint coatings for (a) the sample A3, (b) the sample B4, (c) the sample C4, and (d) the sample D2 [15].

annealed at (b) 600°C, (c) 700°C, (d) 800°C, and (e) 900°C for 1 h [15].

278 Magnetic Spinels- Synthesis, Properties and Applications

#### **3.3. Spectral selectivity absorber coating based on CuMnO***<sup>x</sup>*  **spinel ceramic film**

Spray-coating technique is quick, low-cost, easily adaptable to different coating solutions, and suitable for the establishment of a large-scale process, and there is a minimum of material waste. But as paint coatings are comparatively thicker and the organic binders also absorb in the thermal IR range, these coatings usually suffer from the higher thermal emittance (*ε*<sup>100</sup> > 0.2) [16]. Hence, the preparation and investigation of the spinel thin films by sol-gel route have attracted considerable attention. There is a great demand for low-cost and environmentally friendly techniques for synthesizing high-quality spectral selectivity absorber coatings. Such coatings are capable of absorbing most of the incoming solar radiation (high solar absorbance) without losing much of the thermal energy through reradiation from heated surface (low thermal emittance) [17].

The term spinel refers to a group of minerals, which crystallize in a cubic (isometric) crystal structure. Kaluza et al. [18] have succeeded in synthesizing CuFeMnO4 black film spinel solar absorber coating using sol-gel dip-coating and heat treatment at 500°C. Mn-acetate tetrahydrate, Cu-chloride, and Fe-chloride hexahydrate precursors are used in a molar ratio of 3:3:1, respectively. To protect the spinel from corrosion, a 3-aminopropyltriethoxysilane (3-APTES) silica precursor is added to the Cu, Mn, and Fe sol precursors with a molar ratio of (Mn-Cu-Fe): silica = 1:1. Analytical results show that the films consist of two layers: the lower is amorphous SiO2 , and the upper is a spinel having the composition of Cu1.4Mn1.6O4 . The films exhibit solar absorbance values of around *α*<sup>s</sup> = 0.6 and thermal emittance values of *ε* = 0.29–0.39. Copper-cobalt oxide thin coatings have been deposited on highly infrared-reflecting aluminum substrate via a four-dipping/annealing-cycle sol-gel dip-coating route [19]. Nevertheless, the high annealing temperature and long annealing time would severely wreck the mechanical strength of aluminum substrate. Furthermore, the low solar absorptance (*α*<sup>s</sup> = 0.834) was merely obtained. He and Chen [5] added a complexing agent and an esterifying agent to fabricate the precursor sol, and thus CuCoMnO*<sup>x</sup>* coatings were deposited on aluminum substrate with a *α*<sup>s</sup> value up to 0.93. Mahallawy et al. [17] also successfully synthesized CuCoMnO*<sup>x</sup>* coatings on aluminum and copper substrates by sol-gel dip-coating method. CuMnO*<sup>x</sup>* monolayer coating, CuMnO*<sup>x</sup>* /SiO2 two-layer coating, and CuMnSiO*<sup>x</sup>* /CuMnSiO*<sup>x</sup>* /SiO2 three-layer coating have been fabricated [20], which provide good design strategies for ceramic spectral selectivity (CSS) coatings. Cu1.5Mn1.5O4 -based CSS coating is deposited on aluminum substrate using sol-gel dip-coating method from a stable metal nitrate precursor sol [21]. The Chelating agent citric acid, acting as a reducing agent simultaneously in the exothermic redox reaction, lowered the annealing temperature required by the formation of crystalline Cu1.5Mn1.5O4 . By optimizing the withdrawal rates and annealing temperatures, coating with optical parameter values as good as *α*<sup>s</sup> = 0.876 and *ε*<sup>100</sup> = 0.057 is achieved after only one dipping/annealing cycle. Furthermore, the recycling experiment should be implemented to certify the reproducibility and stability of the metallic precursor sol. After reserved for 20 days, the precursor sol was deposited on aluminum substrate to obtain the CSS coating. **Figure 7** shows the typical x-ray diffraction (XRD) patterns and field emission scanning electron microscopy (FE-SEM) images of the CSS coating, which is deposited at 120 mm/min and annealed at 500°C for 1 h after recycling experiment. As can been seen from the XRD diffraction spectra, the diffraction peaks of the sample at 2*θ* values of 30.46°, 35.85°, 37.52°, and 57.82° correspond to (2 2 0), (3 1 1), (2 2 2), and (3 3 3) crystal planes of cubic spinel structure Cu1.5Mn1.5O4 . The morphology of Cu1.5Mn1.5O4 coating indicates the presence of jagged and uneven pores, which can be attributed to the liberation of abundant H2 O, CO2 , O2 , NO2 , and other NO*<sup>x</sup>* during the heat treating process. One interesting thing worthy to discuss is that the pores are conducive to enhance solar absorber of the CSS coatings. The pores look like light traps where the reflected light can be refracted consecutively and enters into the absorber layer again.

Photothermal Conversion Applications of the Transition Metal (Cu, Mn, Co, Cr, and Fe) Oxides with Spinel Structure http://dx.doi.org/10.5772/67210 281

 **Figure 7.** The XRD patterns and FE-SEM images of the solar selective absorber coating deposited on aluminum substrate after recycling experiment [21].

## **4. Conclusions**

IR range, these coatings usually suffer from the higher thermal emittance (*ε*<sup>100</sup> > 0.2) [16]. Hence, the preparation and investigation of the spinel thin films by sol-gel route have attracted considerable attention. There is a great demand for low-cost and environmentally friendly techniques for synthesizing high-quality spectral selectivity absorber coatings. Such coatings are capable of absorbing most of the incoming solar radiation (high solar absorbance) without losing much of the thermal energy through reradiation from heated surface (low thermal emittance) [17].

The term spinel refers to a group of minerals, which crystallize in a cubic (isometric) crys-

nel solar absorber coating using sol-gel dip-coating and heat treatment at 500°C. Mn-acetate tetrahydrate, Cu-chloride, and Fe-chloride hexahydrate precursors are used in a molar ratio of 3:3:1, respectively. To protect the spinel from corrosion, a 3-aminopropyltriethoxysilane (3-APTES) silica precursor is added to the Cu, Mn, and Fe sol precursors with a molar ratio of (Mn-Cu-Fe): silica = 1:1. Analytical results show that the films consist of two layers: the lower

exhibit solar absorbance values of around *α*<sup>s</sup> = 0.6 and thermal emittance values of *ε* = 0.29–0.39. Copper-cobalt oxide thin coatings have been deposited on highly infrared-reflecting aluminum substrate via a four-dipping/annealing-cycle sol-gel dip-coating route [19]. Nevertheless, the high annealing temperature and long annealing time would severely wreck the mechanical strength of aluminum substrate. Furthermore, the low solar absorptance (*α*<sup>s</sup> = 0.834) was merely obtained. He and Chen [5] added a complexing agent and an esterifying agent to fab-

coatings on aluminum and copper substrates by sol-gel dip-coating method. CuMnO*<sup>x</sup>*

two-layer coating, and CuMnSiO*<sup>x</sup>*

coating have been fabricated [20], which provide good design strategies for ceramic spectral

using sol-gel dip-coating method from a stable metal nitrate precursor sol [21]. The Chelating agent citric acid, acting as a reducing agent simultaneously in the exothermic redox reaction, lowered the annealing temperature required by the formation of crystalline Cu1.5Mn1.5O4

optimizing the withdrawal rates and annealing temperatures, coating with optical parameter values as good as *α*<sup>s</sup> = 0.876 and *ε*<sup>100</sup> = 0.057 is achieved after only one dipping/annealing cycle. Furthermore, the recycling experiment should be implemented to certify the reproducibility and stability of the metallic precursor sol. After reserved for 20 days, the precursor sol was deposited on aluminum substrate to obtain the CSS coating. **Figure 7** shows the typical x-ray diffraction (XRD) patterns and field emission scanning electron microscopy (FE-SEM) images of the CSS coating, which is deposited at 120 mm/min and annealed at 500°C for 1 h after recycling experiment. As can been seen from the XRD diffraction spectra, the diffraction peaks of the sample at 2*θ* values of 30.46°, 35.85°, 37.52°, and 57.82° correspond to (2 2 0), (3 1 1),

coating indicates the presence of jagged and uneven pores, which can be attrib-

, and other NO*<sup>x</sup>*

(2 2 2), and (3 3 3) crystal planes of cubic spinel structure Cu1.5Mn1.5O4

be refracted consecutively and enters into the absorber layer again.

O, CO2

, O2 , NO2

process. One interesting thing worthy to discuss is that the pores are conducive to enhance solar absorber of the CSS coatings. The pores look like light traps where the reflected light can

, and the upper is a spinel having the composition of Cu1.4Mn1.6O4

value up to 0.93. Mahallawy et al. [17] also successfully synthesized CuCoMnO*<sup>x</sup>*

coatings were deposited on aluminum substrate

/CuMnSiO*<sup>x</sup>*


/SiO2

. The morphology of

during the heat treating

black film spi-

. The films

mono-

. By

three-layer

tal structure. Kaluza et al. [18] have succeeded in synthesizing CuFeMnO4

is amorphous SiO2

with a *α*<sup>s</sup>

Cu1.5Mn1.5O4

layer coating, CuMnO*<sup>x</sup>*

ricate the precursor sol, and thus CuCoMnO*<sup>x</sup>*

280 Magnetic Spinels- Synthesis, Properties and Applications

selectivity (CSS) coatings. Cu1.5Mn1.5O4

uted to the liberation of abundant H2

/SiO2

Black-colored transition metal oxides with spinel structure are easy to synthesize via sol-gel methods, and most of them show the high spectral selectivity and thermal stability. However, most of the coatings based on those materials are relatively lower spectral selectivity than the commercial absorber surfaces. More explorations on precursor's combinations, absorber stack configuration and compositions, as well as the application of superior antireflection layer are needed to improve their spectral selectivity. Furthermore, there are still problems associated with the reproducibility of the sol, and few studies have been done on this type of absorber coatings using sol-gel methods. Some problems such as the border effect and the heterogeneity of the film, which have a negative effect on the practical application of spinel films, arise when sol-gel dip-coating method is used for large-scale deposition. The use of highly soluble raw materials and the avoidance of compounds that easily settle in precursor's preparation are the robust ways to solve the reproducibility problem. Additionally, other pertinent factors such as the thickness of silica (especially if used as a matrix), abrasion, corrosion resistance, and the durability of the spinel absorber film should also be examined more extensively in future research.

## **Author details**

Pengjun Ma1 , Qingfen Geng2 and Gang Liu1 \*

\*Address all correspondence to: gangliu@licp.cas.cn

1 Research & Development Center for Eco-material and Eco-chemistry, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou, China

2 CEPREI Certification Body, China CEPREI Laboratory, Guangzhou, China

## **References**


[11] T. Boström, G. Westin, E. Wäckelgård, Optimization of a solution-chemically derived solar absorbing spectrally selective surface, Sol. Energy Mater. Sol. Cells, 2007, 91, 38–43.

**Author details**

, Qingfen Geng2

282 Magnetic Spinels- Synthesis, Properties and Applications

\*Address all correspondence to: gangliu@licp.cas.cn

and Gang Liu1

2 CEPREI Certification Body, China CEPREI Laboratory, Guangzhou, China

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## **Anti‐Corrosive Properties and Physical Resistance of Alkyd Resin–Based Coatings Containing Mg‐Zn‐Fe Spinels Anti‐Corrosive Properties and Physical Resistance of Alkyd Resin–Based Coatings Containing Mg‐Zn‐Fe Spinels**

Kateřina Nechvílová, Andrea Kalendová and Miroslav Kohl Miroslav Kohl

Kateřina Nechvílová, Andrea Kalendová and

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

http://dx.doi.org/10.5772/65545

#### **Abstract**

The anti‐corrosive and physical properties of organic coatings containing spinel pigments Mg0.2Zn0.8Fe2O4 wer studied. Pigments exhibiting different particle morphol‐ ogies were synthesized by high‐temperature solid phase reactions. Core‐shell pigments containing Fe‐Mg‐Zn ferrite shells deposited on non‐isometric particles of mineral cores consisting of layered silicates were also prepared. The pigments were used in paints, the pigment volume concentrations in the binder being 5, 10 and 15%. Anti‐corrosive efficiency was investigated for paint films containing one of three ferro‐spinel (Mg0.2Zn0.8Fe2O4)‐based pigments or one of two core‐shell pigments consisting of Fe‐Mg‐ Zn shell and lamellar silicate‐based cores. The paint properties were examined by accelerated corrosion tests and by physico‐mechanical tests. The relationships between the pigment particle shape and the paint properties were examined. The effect of the pigment particle morphology on the mechanical properties of the paint films was also investigated. The dependence of the paint film properties on the pigment volume concentration was studied and the optimum concentrations providing the most efficient anti‐corrosive protection were determined for each pigment.

**Keywords:** anti‐corrosive pigment, organic coating, spinel, ferrite, core‐shell pigment, accelerated corrosion test

© 2017 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. © 2017 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.

## **1. Introduction**

Metals are routinelyprotectedagainst atmospheric corrosionby coating withpigmentedpaints possessing anti‐corrosive properties [1]. Development of non‐toxic anti‐corrosive pigments for paints protecting the substrate efficiently against corrosion is a major trend in the science and technology of organic coating materials. The goal of research in this area is to develop non‐toxic pigmentspossessingahighanti‐corrosiveefficiencywithaviewtoreplacingleadandchromium (VI)‐based pigments that used to be applied in the past [2–4]. One direction in this area focuses on chemical or physical modification of known pigments, typically of the phosphate type [3, 5]. New inorganic pigments or organic corrosion inhibitor are also developed [6]. Among rather new anti‐corrosive pigments are spinel‐type pigments, finding application as thermally stable colour pigments and also as anti‐corrosive pigments to protect the substrate against corrosion [7, 8]. Owing to their stable structure, spinel pigments exhibit a high resistance against high temperatures and chemical effects [9]. Mixed oxides in the spinel form, particularly ferrites containingCa,Zn,Mgandtheir solidsolutions, appliedinpaints alreadyassume afirmposition in anti‐corrosion protection of metals [10, 11]. More efficient than the spinel‐structure ferrite‐ type pigments with a combination of two cations in the lattice structure (ZnFe2O4, MgFe2O4) are second‐generation spinel pigments with three cations (Mg1‐xZnxFe2O4 and Ca1‐xZnxFe2O4) [12]. Ferrite pigments are reactive particles on the surface in which there occur chemical reactions leading to the formation ofless aggressive substances and a protective film during the diffusion of the corrosion substances through the polymeric film [13].

The protective properties of organic coatings are dependent, among other factors, on their impermeability to liquids and gases [14–16]. Protective film impermeability to liquids and gases can be improved by using non‐isometric fillers [17, 18], e.g. micaceous iron oxide (specularite) [19–21]. **Figure 1** shows schematically the effect of non‐isometric pigment particles on the physical properties of the paint films, consisting in an improvement of the mechanical properties of the films, reduction of the rate of diffusion through the films and protection of the binder against UV radiation [22, 23].

**Figure 1.** Pigment particles with a lamellar morphology in the paint film [29, 30].

Among pigments exhibiting non‐isometric particle morphology are also some core‐shell pigments [24]. The composition of both the core and the shell can be very diverse. Inorganic substances or metal powders are frequently used [25, 26]. Among assets of core‐shell pigments is their economy, because raw materials from natural sources can often be used as the cores [27]. Control of the particle morphology, owing to which the product particles retain the shape of the initial core particles, is another reason why such pigments are used [28].

**1. Introduction**

286 Magnetic Spinels- Synthesis, Properties and Applications

Metals are routinelyprotectedagainst atmospheric corrosionby coating withpigmentedpaints possessing anti‐corrosive properties [1]. Development of non‐toxic anti‐corrosive pigments for paints protecting the substrate efficiently against corrosion is a major trend in the science and technology of organic coating materials. The goal of research in this area is to develop non‐toxic pigmentspossessingahighanti‐corrosiveefficiencywithaviewtoreplacingleadandchromium (VI)‐based pigments that used to be applied in the past [2–4]. One direction in this area focuses on chemical or physical modification of known pigments, typically of the phosphate type [3, 5]. New inorganic pigments or organic corrosion inhibitor are also developed [6]. Among rather new anti‐corrosive pigments are spinel‐type pigments, finding application as thermally stable colour pigments and also as anti‐corrosive pigments to protect the substrate against corrosion [7, 8]. Owing to their stable structure, spinel pigments exhibit a high resistance against high temperatures and chemical effects [9]. Mixed oxides in the spinel form, particularly ferrites containingCa,Zn,Mgandtheir solidsolutions, appliedinpaints alreadyassume afirmposition in anti‐corrosion protection of metals [10, 11]. More efficient than the spinel‐structure ferrite‐ type pigments with a combination of two cations in the lattice structure (ZnFe2O4, MgFe2O4) are second‐generation spinel pigments with three cations (Mg1‐xZnxFe2O4 and Ca1‐xZnxFe2O4) [12]. Ferrite pigments are reactive particles on the surface in which there occur chemical reactions leading to the formation ofless aggressive substances and a protective film during the diffusion

The protective properties of organic coatings are dependent, among other factors, on their impermeability to liquids and gases [14–16]. Protective film impermeability to liquids and gases can be improved by using non‐isometric fillers [17, 18], e.g. micaceous iron oxide (specularite) [19–21]. **Figure 1** shows schematically the effect of non‐isometric pigment particles on the physical properties of the paint films, consisting in an improvement of the mechanical properties of the films, reduction of the rate of diffusion through the films and

Among pigments exhibiting non‐isometric particle morphology are also some core‐shell pigments [24]. The composition of both the core and the shell can be very diverse. Inorganic

of the corrosion substances through the polymeric film [13].

protection of the binder against UV radiation [22, 23].

**Figure 1.** Pigment particles with a lamellar morphology in the paint film [29, 30].

An ideal anti‐corrosive pigment will not only possess good corrosion‐inhibition properties but it will also have a favourable effect on the paint film's mechanical resistance [31, 32]. Among candidates that are studied in this context are pigments based on mixed oxides, specifically ferrites. The development of pigments whose primary particles possess a non‐isometric, lamellar shape represents another direction in the development of anti‐corrosion pigments. Such pigments, with their alkaline properties, can enhance both the barrier and inhibiting mechanisms of the protective paint system [33]. Examples include Zn, Ca, Mg ferrites with a lamellar, non‐isometric particle shape which, apart from their anti‐corrosion properties, improve the paint film's mechanical properties such as bending strength, elasticity and adhesion to the substrate [34]. The resulting ferrite spinels were found to retain particle morphology of the initial iron oxides [35]. The most efficient anti‐corrosive ferrites include ferro‐spinels containing zinc and magnesium in specific ratios [9, 36]. It has been found that ferrites with Mg2+ and Zn2+ as the divalent cations exhibit a very good anti‐corrosive efficiency, especially if their molar ratio is 0.2:0.8 [37]. The studies are continued in this work, where one of the phenomena investigated was the effect of the particle size on the properties of the protective paint films based on a modified alkyd resin.

The aim of this study was to synthesize and investigate the properties of Mg‐Zn‐Fe ferro‐ spinel‐based pigment possessing different particle morphologies with the goal to prepare such paints pigmented with them as they exhibit very good physical and anticorrosive properties. Core‐shell pigments containing silicate cores and Mg‐Zn‐Fe shell were also prepared and studied with focus on their properties in protective coatings in a simulated corrosive environ‐ ment. The particle morphology of such pigments is dictated by the core particle's shape and size, affecting the physical properties and barrier properties of the organic coating in corrosive environments. In order to get a deeper insight into the role of the pigment in the paint, the effect of the core material structure and chemical composition as well as the effect of the pigment concentration on the barrier capacity and the physical and anti‐corrosive properties of the paint in the corrosive environment should be investigated [38, 39]. Layered silicates, specifically talc and calcined kaolin, exhibited differences in the composition, physical properties of their surface and particle texture, which may affect the pigment properties in the paints. Talc, Mg3(OH)2 (Si4O10), is a soft hydrophobic mineral (hardness 1 on the Mosh scale) with lamellar particles. Calcined kaolin obtained by calcination of the mineral kaolinite Al4(OH)8 (Si4O10) at 600°C contains morphologically diverse particles, largely lamellar, is harder than talc and contains more structural phases (largely mullite and quartz) [40–42]. All that may play a role in the overall anti‐corrosive properties of the core‐shell pigments [43, 44]. Core‐shell pigments represent an economical solution for the synthesis of anti‐corrosive pigments [27] obtained by the high‐temperature reaction. Moreover, talc and kaolin are natural materials which are reasonably well available and are environmentally and toxicologically acceptable.

## **2. Experimental methods and procedures**

#### **2.1. Pigment synthesis**

The pigments were formulated as the oxides Mg0.2Zn0.8Fe2O4 possessing a spinel structure. Efforts were made during the synthesis of the ferrite Mg0.2Zn0.8Fe2O4 ("pigment *A*") to achieve (create) the spinel structure and isometric particles. Therefore, α‐Fe2O3 (hematite), whose particles are regularly nodular, was used as the starting material. The aim of the synthesis of the next Mg0.2Zn0.8Fe2O4 ferrite ("pigment *B*") was to obtain a with a non‐isometric—specifi‐ cally needle‐shaped—particle morphology. Goethite α‐FeO (OH), with needle‐shaped particles, was used as the ferric oxide source. The third Mg0.2Zn0.8Fe2O4 ferrite ("pigment *C*") was prepared by an identical procedure with the aim to prepare a non‐isometric spinel pigment with a lamellar particle shape. Specularite, a ferric oxide Fe2O3 with a lamellar particle structure, was used as the starting material. The aim of the synthesis of the core‐shell pigments with Mg‐Zn‐Fe shell was to coat the mineral core with a functional layer of mixed Mg‐Zn‐Fe oxides in the 0.2:0.8:2.0 ratios. Calcined Kaolin and talc Mg3(OH)2(Si4O10), i.e. materials exhibiting a lamellar particle shape, were used as the mineral carriers (cores). The core‐shell pigments were synthesized by applying the matrix core (kaolin, talc)‐to‐mixed oxide (Mg‐Zn‐ Fe) shell weight ratio 1:1.

The incoming materials were homogenized for 45 minutes, not only to obtain intimate mixtures but also to achieve mechanical‐chemical activation of the materials, enhance contact of the particles in the powder mixture and increase the contact area for the reaction [4]. The high‐ temperature reaction was achieved by reaction mixture calcination in an electric furnace. The process was conducted as a two‐stage procedure: the pigments were first calcined at 1000°C for 2 hours and then at 1180°C. The calcination temperatures were selected based on previous X‐ray analysis of the products of tentative experiments. The resulting calcined material was multiply rinsed with distilled water. In order to obtain the right pigment particle size for use in paints, the calcinate was subjected to wet milling in a planetary ball mill (Pulverisette 6, Nietzsche, Germany). Accommodated in a milling vessel made of a zirconium silicate ceramics, the material was milled at 480 rpm for 420 minutes. The milling bodies were rollers made of the same ceramic material. The finely ground pigments were rinsed with water again and dried at 105°C in a laboratory‐scale electric furnace. The core‐shell pigments (or the mixed oxides forming the functional layer on the core) were also synthesized by high temperature solid‐ phase reaction, viz. by two‐stage calcination like the non‐core ferrites [43, 44].

Structural purity of the products was checked and their X‐ray diffraction spectra were measured on a D8 Advance Diffractometer (Bruker AXS). The pigment particle surface and shape were examined on a JEOL‐JSM 5600 LV scanning electron microscope (JEOL, Japan) [6].

#### **2.2. Measuring the anticorrosion efficiency of the pigments in paints**

In order to assess their potential anti‐corrosion efficiency, the pigments were added to a solution of an alkyd resin modified with soy oil (density 1.1 g/cm3 , dry matter fraction 58.9%). The pigment volume concentrations (PVC) selected for the paints were 5, 10 and 15%. The PVC/CPVC ratio (CPVC = critical PVC) was invariably adjusted to 0.35 by using limestone (CaCO3) as an anti‐corrosive‐neutral filler. Cobalt octoate at a fraction of 0.3 wt.% was used as the siccative. The paints were prepared by dispersing the powdered pigments in the liquid binder in a Dispermat CV pearl mill. Test samples were prepared by coating test steel panels (deep‐drawn cold‐rolled steel manufactured by Q‐panel, UK) 150 mm × 100 mm × 0.9 mm size with the paints by means of a box‐type application ruler 200 μm slot width as per ISO 1514. The dry film thickness (DFT) was measured as per ISO 2808. Ten panels were prepared for each paint. A thin line 7 cm long and penetrating as deep as the substrate was cut into the paint film by using a sharp blade. Paint films were also prepared on polyethylene sheets, allowed to dry and peeled off and cut to pieces, 1 mm × 1 mm size. Such unsupported films were used to prepare 10% aqueous suspensions of the paint films in redistilled water [45].

**2. Experimental methods and procedures**

288 Magnetic Spinels- Synthesis, Properties and Applications

The pigments were formulated as the oxides Mg0.2Zn0.8Fe2O4 possessing a spinel structure. Efforts were made during the synthesis of the ferrite Mg0.2Zn0.8Fe2O4 ("pigment *A*") to achieve (create) the spinel structure and isometric particles. Therefore, α‐Fe2O3 (hematite), whose particles are regularly nodular, was used as the starting material. The aim of the synthesis of the next Mg0.2Zn0.8Fe2O4 ferrite ("pigment *B*") was to obtain a with a non‐isometric—specifi‐ cally needle‐shaped—particle morphology. Goethite α‐FeO (OH), with needle‐shaped particles, was used as the ferric oxide source. The third Mg0.2Zn0.8Fe2O4 ferrite ("pigment *C*") was prepared by an identical procedure with the aim to prepare a non‐isometric spinel pigment with a lamellar particle shape. Specularite, a ferric oxide Fe2O3 with a lamellar particle structure, was used as the starting material. The aim of the synthesis of the core‐shell pigments with Mg‐Zn‐Fe shell was to coat the mineral core with a functional layer of mixed Mg‐Zn‐Fe oxides in the 0.2:0.8:2.0 ratios. Calcined Kaolin and talc Mg3(OH)2(Si4O10), i.e. materials exhibiting a lamellar particle shape, were used as the mineral carriers (cores). The core‐shell pigments were synthesized by applying the matrix core (kaolin, talc)‐to‐mixed oxide (Mg‐Zn‐

The incoming materials were homogenized for 45 minutes, not only to obtain intimate mixtures but also to achieve mechanical‐chemical activation of the materials, enhance contact of the particles in the powder mixture and increase the contact area for the reaction [4]. The high‐ temperature reaction was achieved by reaction mixture calcination in an electric furnace. The process was conducted as a two‐stage procedure: the pigments were first calcined at 1000°C for 2 hours and then at 1180°C. The calcination temperatures were selected based on previous X‐ray analysis of the products of tentative experiments. The resulting calcined material was multiply rinsed with distilled water. In order to obtain the right pigment particle size for use in paints, the calcinate was subjected to wet milling in a planetary ball mill (Pulverisette 6, Nietzsche, Germany). Accommodated in a milling vessel made of a zirconium silicate ceramics, the material was milled at 480 rpm for 420 minutes. The milling bodies were rollers made of the same ceramic material. The finely ground pigments were rinsed with water again and dried at 105°C in a laboratory‐scale electric furnace. The core‐shell pigments (or the mixed oxides forming the functional layer on the core) were also synthesized by high temperature solid‐

phase reaction, viz. by two‐stage calcination like the non‐core ferrites [43, 44].

**2.2. Measuring the anticorrosion efficiency of the pigments in paints**

solution of an alkyd resin modified with soy oil (density 1.1 g/cm3

Structural purity of the products was checked and their X‐ray diffraction spectra were measured on a D8 Advance Diffractometer (Bruker AXS). The pigment particle surface and shape were examined on a JEOL‐JSM 5600 LV scanning electron microscope (JEOL, Japan) [6].

In order to assess their potential anti‐corrosion efficiency, the pigments were added to a

The pigment volume concentrations (PVC) selected for the paints were 5, 10 and 15%. The

,

dry matter fraction 58.9%).

**2.1. Pigment synthesis**

Fe) shell weight ratio 1:1.

The cyclic corrosion test in an environment with condensing water and SO2 was performed as per CSN EN ISO 3231. This test consisted of 24‐hour cycles comprising condensation of distilled water with SO2 (0.2 mg/l) at 36°C for 8 hours followed by drying at 23 ± 2°C for 16 hours. The samples were evaluated after 1392 test hours. The corrosion tests were evaluated in accordance with the standards ASTM D 714‐87, CSN ISO 2409, ASTM D 610‐85 and ASTM D 1654‐92. The following corrosion effects were assessed: formation (= size and frequency of occurrence) of blisters on the film surface and near the test cut made in the film; extent of substrate metal surface corrosion (corrosion‐affected surface area fraction in %); and propa‐ gation (in mm) of corrosion in the vicinity of the test cut. Anti‐corrosion efficiency on the 100−0 scale (100 = excellent anti‐corrosion efficiency, 0 = poor anti‐corrosion efficiency) was assigned to the corrosion effects in an environment with SO2 [46].

#### *2.2.1. Determination of corrosion‐induced steel mass loss in aqueous extracts of the pigments*

In this test, steel panels of defined size and known weight were submerged for 10 days in filtrates of 10 wt.% aqueous suspensions of the powdered pigments tested. The observed corrosion weight losses were converted to relative data (%) with respect to the values in water. The relative data obtained from the suspensions of the powdered pigments are denoted *X*corrp. Relative corrosion losses (*X*corrf) of steel panels exposed to 10% aqueous extracts of the loose pigmented paint films during 10 days were also measured [47, 48].

#### *2.2.2. Effect of the pigments on the physico‐mechanical properties of the organic coatings*

The following tests were performed on the paint films: resistance to bending over a cylindrical mandrel; resistance to the cupping test (Erichsen test); resistance to a weight dropping on the reverse side of the test panels and adhesion (cross‐cut test, or lattice test, knife spacing 1 mm). Degree of adhesion of the paints (ISO 2409) was performed by the cross‐cut test. Impact resistance (ISO 6272) measured the maximum height of free drop of a weight (1000 g) at which the paint film still resisted damage. Resistance of the paint film against cupping (ISO 1520) was made in an Erichsen cupping tester. Resistance of the coating during bending over a cylindrical mandrel (ISO 1519) provides the largest diameter of the mandrel (in mm) causing disturbance of the paint film when the test panel is bent over it [49]. The results of the tests were used to calculate the overall physical‐mechanical efficiency, i.e. overall paint film resistance to mechanical effects. A resistance score on the 100−0 scale (100 = excellentresistance, 0 = poorresistance) was assigned to each testresult. The overall physical‐mechanicalresistance of the paints (*ME*) was calculated as the arithmetic mean of the cohesion score from the bending test on the cylindrical mandrel, the resilience score from the impact test, the degree of adhesion score and the strength score from the cupping [50].

The starting hematite was also examined (pigment *F*). The non‐pigmented organic coating material was subjected to the mechanical resistance tests and corrosion tests as well.

## **3. Results and discussion**

#### **3.1. Pigment morphology and structure**

Morphology of the pigment particles synthesized is shown in the SEM photographs in **Figure 2**. The aim of the synthesis of the pigments *A*, *B* and *C* was to obtain pigments possessing the spinel structure whose chemical formula is *Mg0.2Zn0.8Fe2O4.* Mixed spinel ferrites containing zinc and magnesium cations constituted the majority phases. The magnesium atoms can be iso‐morphically substituted by zinc atoms, owing to which the mixed spinels can be formed across an unlimited range of concentrations. The magnesium‐to‐zinc cation ratio could not be determined precisely from the analysis because the lattice parameters in MgFe2O4 and ZnFe2O4 approach each other and the diffraction lines may overlap. The X‐ray diffraction analysis of the pigments *A*, *B* and *C* did not exhibit any crystal phases of the starting materials; in other words, the starting materials had completely reacted to the products. These results are consistent with [37]. Furthermore, the three spinel types were synthesized with a view to obtaining pigments *Mg0.2Zn0.8Fe2O4* possessing different particle shapes, making use of the different particle shapes of the starting ferric oxide types: the particles of pigment *A* were isometric, the particles of pigment *B* were needle‐shaped (acicular) and the particles of pigment *C* were lamellar. The three pigments are also referred to as isometric *Mg0.2Zn0.8Fe2O*, acicular *Mg0.2Zn0.8Fe2O4* and lamellar *Mg0.2Zn0.8Fe2O4*, respectively, throughout this text.

The core‐shell pigment *D* (*Mg0.2Zn0.8Fe2O4/kaolin*) was prepared as a pigment whose non‐ isometric (lamellar) core particles of calcined kaolin are coated with a functional shell of the mixed oxide *Mg0.2Zn0.8Fe2O4*. For this, the starting materials for the shell were mixed in proportions providing the Mg‐Zn‐Fe cation ratio 0.2:0.8:2. Mixed oxides of aluminium, iron and magnesium were identified (the mixed oxide hercynite). The analysis gave evidence that chemical reactions had occurred between the kaolin surface and the remaining starting materials: indeed, the mixed oxide hercynite (FeMg3Al2O4) contains the aluminium cation, which initially was present in the starting kaolin. In other words, a chemically bonded functional layer of mixed oxides had formed on the kaolin particles. The X‐ray diffraction analysis showed that the pigment also included silicates (MgSiO3) and crystalline silicon oxide (SiO2) phases (cristobalite) from the starting kaolin. The resulting pigment *D* (simplified referred to as *Mg0.2Zn0.8Fe2O4/kaolin* or *Mg0.2Zn0.8Fe2O4/Al6Si2O13*) exhibited a lamellar particle shape.

Anti‐Corrosive Properties and Physical Resistance of Alkyd Resin–Based Coatings Containing Mg‐Zn‐Fe Spinels http://dx.doi.org/10.5772/65545 291

resistance to mechanical effects. A resistance score on the 100−0 scale (100 = excellentresistance, 0 = poorresistance) was assigned to each testresult. The overall physical‐mechanicalresistance of the paints (*ME*) was calculated as the arithmetic mean of the cohesion score from the bending test on the cylindrical mandrel, the resilience score from the impact test, the degree of adhesion

The starting hematite was also examined (pigment *F*). The non‐pigmented organic coating

Morphology of the pigment particles synthesized is shown in the SEM photographs in **Figure 2**. The aim of the synthesis of the pigments *A*, *B* and *C* was to obtain pigments possessing the spinel structure whose chemical formula is *Mg0.2Zn0.8Fe2O4.* Mixed spinel ferrites containing zinc and magnesium cations constituted the majority phases. The magnesium atoms can be iso‐morphically substituted by zinc atoms, owing to which the mixed spinels can be formed across an unlimited range of concentrations. The magnesium‐to‐zinc cation ratio could not be determined precisely from the analysis because the lattice parameters in MgFe2O4 and ZnFe2O4 approach each other and the diffraction lines may overlap. The X‐ray diffraction analysis of the pigments *A*, *B* and *C* did not exhibit any crystal phases of the starting materials; in other words, the starting materials had completely reacted to the products. These results are consistent with [37]. Furthermore, the three spinel types were synthesized with a view to obtaining pigments *Mg0.2Zn0.8Fe2O4* possessing different particle shapes, making use of the different particle shapes of the starting ferric oxide types: the particles of pigment *A* were isometric, the particles of pigment *B* were needle‐shaped (acicular) and the particles of pigment *C* were lamellar. The three pigments are also referred to as isometric *Mg0.2Zn0.8Fe2O*, acicular

*Mg0.2Zn0.8Fe2O4* and lamellar *Mg0.2Zn0.8Fe2O4*, respectively, throughout this text.

The core‐shell pigment *D* (*Mg0.2Zn0.8Fe2O4/kaolin*) was prepared as a pigment whose non‐ isometric (lamellar) core particles of calcined kaolin are coated with a functional shell of the mixed oxide *Mg0.2Zn0.8Fe2O4*. For this, the starting materials for the shell were mixed in proportions providing the Mg‐Zn‐Fe cation ratio 0.2:0.8:2. Mixed oxides of aluminium, iron and magnesium were identified (the mixed oxide hercynite). The analysis gave evidence that chemical reactions had occurred between the kaolin surface and the remaining starting materials: indeed, the mixed oxide hercynite (FeMg3Al2O4) contains the aluminium cation, which initially was present in the starting kaolin. In other words, a chemically bonded functional layer of mixed oxides had formed on the kaolin particles. The X‐ray diffraction analysis showed that the pigment also included silicates (MgSiO3) and crystalline silicon oxide (SiO2) phases (cristobalite) from the starting kaolin. The resulting pigment *D* (simplified referred to as *Mg0.2Zn0.8Fe2O4/kaolin* or *Mg0.2Zn0.8Fe2O4/Al6Si2O13*) exhibited a lamellar particle

material was subjected to the mechanical resistance tests and corrosion tests as well.

score and the strength score from the cupping [50].

**3. Results and discussion**

shape.

**3.1. Pigment morphology and structure**

290 Magnetic Spinels- Synthesis, Properties and Applications

**Figure 2.** Morphology of the particles of tested pigments (SEM): (a) Mg0.2Zn0.8Fe2O4 isometric; (b) Mg0.2Zn 0.8Fe2O4 acicu‐ lar; (c) Mg0.2Zn0.8Fe2O4 lamellar; (d) mixed oxide Mg‐Zn‐Fe/kaoline; (e) mixed oxide Mg‐Zn‐Fe/talc.

Pigment *E* was synthesized with a view to obtaining a core‐shell pigment where the mixed Mg‐Zn‐Fe oxide shell covers a mineral core possessing a lamellar particle structure, specifi‐ cally talc *(Mg3(OH)2Si4O10)*. The highest diffraction line belonged to the crystalline ferrite phase. As in the remaining pigments, isomorphic replacement of the zinc cation by a magnesium cation had occurred in this ferrite layer and so their precise ratio could not be

determined. Components of minor importance in this pigment included magnetite (Fe3O4) from the starting hematite and the silicon oxide cristobalite from talc in the core. The resulting pigment *E* (simply referred to as *Mg0.2Zn0.8Fe2O4/talc* or *Mg0.2Zn0.8Fe2O4/Mg3*(*Si4O10*)(*OH*)*2*) exhibited a lamellar particle shape.

## **3.2. Steel panel mass loss in aqueous extracts of the powdered pigments and of the loose pigmented films, and pH and conductivity of the pigmented aqueous extracts**

**Table 1** contains the calculated relative corrosion losses of the steel panels in aqueous extracts of the powdered pigments (*X*corrp) and in aqueous extracts of the paint films (*X*corrf), and pH values (pH*<sup>f</sup>* ) and specific electric conductivities (*χf* ) of aqueous extracts of the loose paint films.


\* Parameters are given as arithmetic averages within three measured values.

a pH was measured with an accuracy of ±0.01.

b Conductivity was measured with an accuracy of ±0.5%.

**Table 1.** Relative steel mass losses due to corrosion in aqueous extracts of the powdered pigments (*X*corrp) and in aqueous extracts of the loose paint films (*X*corrf,) and pH and specific electric conductivity values (pH*<sup>f</sup>* , *χf* ) of the aqueous extracts of the loose paint films (PVC = 10%).

The corrosion loss data (*X*corp) characterize the pigment's ability to affect the resistance of the metal to corrosion in the pigment extract, where ions passivating the metal surface are present to a larger or lesser extent. The data demonstrate that the corrosion phenomena are partly inhibited by the presence of the pigments [18]. Low steel losses by corrosion were observed with the core‐shell pigments *D* (*mixed Mg‐Zn‐Fe oxide/kaolin*), and *E* (mixed *Mg‐Zn‐Fe oxide/ talc*) and with the lamellar ferrite *C* (lamellar *Mg0.2Zn0.8Fe2O4*), viz. 36, 39 and 44%, respectively. Those (comparable) corrosion losses were related to the extracts' specific conductivities or pH levels [29]. Lower corrosion losses were observed with the core‐shell pigments, where the silicate core and the silicate phases contributed favourably to corrosion protection of the steel panels [44].

determined. Components of minor importance in this pigment included magnetite (Fe3O4) from the starting hematite and the silicon oxide cristobalite from talc in the core. The resulting pigment *E* (simply referred to as *Mg0.2Zn0.8Fe2O4/talc* or *Mg0.2Zn0.8Fe2O4/Mg3*(*Si4O10*)(*OH*)*2*)

**3.2. Steel panel mass loss in aqueous extracts of the powdered pigments and of the loose**

**Table 1** contains the calculated relative corrosion losses of the steel panels in aqueous extracts of the powdered pigments (*X*corrp) and in aqueous extracts of the paint films (*X*corrf), and pH

> **pH***<sup>f</sup>*  **(‐)**

*X***corrf (%)** *X***corrp. (%) pH28** *χ***<sup>28</sup>**

19.66 78.82 6.2 ± 0.01 190

15.99 76.76 6.90 ± 0.01 195

4.29 33.31 7.90 ± 0.01 199

9.40 36.09 6.10 ± 0.01 200

9.98 43.93 6.70 ± 0.01 180

28.15 89.87 5.10 ± 0.01 170

**Table 1.** Relative steel mass losses due to corrosion in aqueous extracts of the powdered pigments (*X*corrp) and in

aqueous extracts of the loose paint films (*X*corrf,) and pH and specific electric conductivity values (pH*<sup>f</sup>*

) of aqueous extracts of the loose paint films.

*χ* **(μS/cm)**

**Specific conductivityb**

, *χf* ) of the

**pigmented films, and pH and conductivity of the pigmented aqueous extracts**

) and specific electric conductivities (*χf*

Water ‐ 100 ‐ ‐ **Non‐pigmented paint film** 100 ‐ 3.04 ± 0.01 216

Parameters are given as arithmetic averages within three measured values.

pH was measured with an accuracy of ±0.01.

Conductivity was measured with an accuracy of ±0.5%.

aqueous extracts of the loose paint films (PVC = 10%).

**Pigment Corrosion losses\* pHa**

exhibited a lamellar particle shape.

292 Magnetic Spinels- Synthesis, Properties and Applications

values (pH*<sup>f</sup>*

(Mg0.2Zn0.8Fe2O4) isometric

(Mg0.2Zn0.8Fe2O4) acicular

(Mg0.2Zn0.8Fe2O4) lamellar

(Mg‐Zn‐Fe) mixed oxide/kaolin

(Mg‐Zn‐Fe) mixed oxide/talc

(hematite) Fe2O3

**A**

**B**

**C**

**D**

**E**

**F**

\*

a

b

The pigments also affected the pH and specific conductivity levels of the paint films pigmented with them. The alkaline nature of the pigments induced changes in the pH of the pigmented films compared to the non‐pigmented film. The pigments shifted the pH values upwards due to the neutralization of the acid components of the binder [5, 51]. This effect was most pro‐ nounced for pigment *C* (lamellar *Mg0.2Zn0.8Fe2O4*), where pH*<sup>f</sup>* was 7.9 (**Table 1**). The core‐shell pigment *E* (mixed *Mg‐Zn‐Fe oxide/talc*) was a next pigment imparting an alkaline nature to the paint. The effects were reasonable, the pigments themselves inducing alkalinity in their aqueous extracts. Most acid (pH*<sup>f</sup>* = 3.0) was the aqueous extract of the non‐pigmented film. The conductivities of the aqueous extracts of the loose paint films were also affected by the pigments. The conductivity of the extract of the non‐pigmented film was *χf* = 216 μS/cm. The conductivity was lower if a pigment was present. This was due to neutralization of the binder's acid components (R‐COOH) giving rise to metallic soaps, where dissociation was suppressed [52]. This effect was marked, once again, with pigment *C* (lamellar *Mg0.2Zn0.8Fe2O4*) and, to a lesser extent, with the core‐shell pigment *E* (*mixed Mg‐Zn‐Fe oxide/talc*) [43, 44].

The anti‐corrosive efficiency of a paint film depends on the pigment's ability to release inhibiting components that are involved in the reactions inside the film and affect the diffusing environment. The corrosion losses in the suspensions of the organic films (*X*corrf) containing anti‐corrosive pigments provide information on the potential reactions occurring inside the paint in the liquid state, during the film formation and, to some extent, during the ageing of the cured film [48, 53]. The binder was an alkyd resin, a polyester modified by a fatty acid, which is well suited to the indirect examination of the behaviour of the ferrite and of the paint film. The carboxy groups in the binder were responsible for the acid nature of the aqueous extract of the film: the observed value was pH 3.1 when no pigment was present. However, if a pigment was present, the pH levels were largely as high as pH 6−8, due to the formation of metallic soaps inside the film [54]. Alkaline anti‐corrosive pigments neutralized the acid groups in the binder [55]. This effect is basically similar to the inhibiting effect of minimum. Metallic soaps may also exhibit inhibiting properties at the protected metal–organic coating interface [6]. The alkyd binder contains both acid ‐COOH groups and an alkaline Zn0.8Mg0.2Fe2O4 pigment, as shown schematically in **Figure 3** [56]. Neutralization of the carboxy groups inside the film results in more or less neutral products: neutral metallic soaps M*n+* ( ‐ OOCR)*n* acidic metallic soaps M*n+* (OOCR)*n*/RCOOH, and alkaline metallic soaps M*n+* (O2‐)*x*(OH‐ )*y*( ‐ OOCR)*z* (*x*/2 + *y* + *z* = *n*). In the formulas, R is a hydrocarbon and M is a metal (e.g. Zn, Mg) whose oxidation state [36]. The acid components of the binder that were released into the aqueous medium may also be neutralized by the alkaline extract of the anti‐corrosive pigment [54, 57]. The low corrosion losses in the extracts of the loose paint films corroborated the results obtained with the extracts of the powdered pigments. The most marked active chemical protection was observed with the core‐shell pigments *D* (*mixed Mg‐Zn‐Fe oxide/ kaolin*) (*X*corrf = 36.09%) and **C** (*lamellar Mg0.2Zn0.8Fe2O4*) (*X*corrf = 39.31%).

**Figure 3.** Schematic illustration of the processes running on extraction of alkyd coatings containing Mg‐Zn‐Fe pigment in an aqueous medium [54].

#### **3.3. Mechanical resistance tests of the paint films containing the pigments**

Outstanding mechanical properties constitute a precondition for good anti‐corrosion efficiency of a paint film. **Table 2** presents the results of mechanical tests of the paints' films containing the pigments. The non‐pigmented binder and a paint containing the starting ferric pigment *F* (hematite *Fe2O3*) were also measured as reference materials. Resistance in the dropping weight test was 100 cm for all the paint films. Similarly, resistance to the bending test was invariably less than 4 cm. Since the values were identical for all the pigments, they are not included in **Table 2**. The pigmented paints films exhibited very good resistance against mechanical stress at any PVC. The paint film was never disturbed when bent over the mandrel 4 mm in diameter. No defects were observed on the films after the test hammer was dropped from the largest height, i.e. 100 cm. Good results were also obtained for all of the pigments in the cross‐cut test, any defects were negligible and nearly all of the films scored 1 in this test. The only pigment for which defects were observed on the lattice (particularly at higher PVC levels) was the core‐ shell pigment *D* (*mixed Mg‐Zn‐Fe oxide/kaolin*), and so the films scored 2. The films gave very good results also in the cupping test, no film disturbance was observed with a magnifying glass after the body penetration to 8 mm distance. The paint film resistance in this test was PVC‐dependent, viz. so that the resistance against penetration decreased slightly with increasing PVC. Pigment *A* (*isometric Mg0.2Zn0.8Fe2O4,*) was an exception in this respect. **Table 3** also demonstrates that the highest resistance against the test body penetration was obtained with pigment *B* having needle‐shaped particles (*acicular Mg0.2Zn0.8Fe2O4*), viz. 9.5 mm at PVC = 5%. Good results were also obtained with the core‐shell pigment *E* (*mixed Mg‐Zn‐Fe oxide/ talc*), where the film was disturbed at 9.4 mm.

(e.g. Zn, Mg) whose oxidation state [36]. The acid components of the binder that were released into the aqueous medium may also be neutralized by the alkaline extract of the anti‐corrosive pigment [54, 57]. The low corrosion losses in the extracts of the loose paint films corroborated the results obtained with the extracts of the powdered pigments. The most marked active chemical protection was observed with the core‐shell pigments *D* (*mixed Mg‐Zn‐Fe oxide/*

**Figure 3.** Schematic illustration of the processes running on extraction of alkyd coatings containing Mg‐Zn‐Fe pigment

Outstanding mechanical properties constitute a precondition for good anti‐corrosion efficiency of a paint film. **Table 2** presents the results of mechanical tests of the paints' films containing the pigments. The non‐pigmented binder and a paint containing the starting ferric pigment *F* (hematite *Fe2O3*) were also measured as reference materials. Resistance in the dropping weight test was 100 cm for all the paint films. Similarly, resistance to the bending test was invariably less than 4 cm. Since the values were identical for all the pigments, they are not included in **Table 2**. The pigmented paints films exhibited very good resistance against mechanical stress at any PVC. The paint film was never disturbed when bent over the mandrel 4 mm in diameter. No defects were observed on the films after the test hammer was dropped from the largest height, i.e. 100 cm. Good results were also obtained for all of the pigments in the cross‐cut test, any defects were negligible and nearly all of the films scored 1 in this test. The only pigment for which defects were observed on the lattice (particularly at higher PVC levels) was the core‐ shell pigment *D* (*mixed Mg‐Zn‐Fe oxide/kaolin*), and so the films scored 2. The films gave very good results also in the cupping test, no film disturbance was observed with a magnifying

**3.3. Mechanical resistance tests of the paint films containing the pigments**

in an aqueous medium [54].

*kaolin*) (*X*corrf = 36.09%) and **C** (*lamellar Mg0.2Zn0.8Fe2O4*) (*X*corrf = 39.31%).

294 Magnetic Spinels- Synthesis, Properties and Applications


\* Parameters are given as arithmetic averages within three measured values.

**Table 2.** Results of mechanical tests of the paints containing the pigments synthesized and the reference pigment and of the reference paint (dry film thickness DFT = (60 ± 10) μm.

Nearly all of the paint films exhibited very good overall physico‐mechanical resistance, their score was 95. Only the paints with the core‐shell pigment *D* (*mixed Mg‐Zn‐Fe oxide/kaolin*) provided poorer cross‐cut test at high PVCs and scored 91 only. The results give evidence that owing to their morphological properties, the pigments tested do not detract from the mechan‐ ical resistance of the alkyd resin–based paints. These results are very important for practical applications of these pigments in paints.


**Table 3.** Corrosion resistance of the paints with the pigments synthesized and with the reference pigment and of the reference paint, measured after 912 hours exposure of the steel panels coated with the paints in an atmosphere with SO2, DFT = (85 ± 10) μm.

#### **3.4. Corrosion tests of the pigmented paints**

**Table 3** lists the results of corrosion resistance measurements of the paint films in the atmosphere with condensed moisture and SO2, obtained after 912 hours of exposure. Nearly all of the paint films had osmotic blisters on their surface after the exposure. The paint films with pigment *C* (*lamellar Mg0.2Zn0.8Fe2O4*) were most resistant in this respect, especially at PVC =10%, where no blisters were found on the film surface and a few small blisters only (6F) were observed around the test cut. The results observed when the pigment concentration was PVC = 5% were also very good (6M). Equally good results were obtained with the paint containing pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*). This was true particularly at PVC = 5%, whereas a higher occurrence of blisters was observed at higher concentrations. So, the blisters observed around the test cut at PVC = 10% were categorized as 4MD, as against the 4F score at PVC = 5%.

owing to their morphological properties, the pigments tested do not detract from the mechan‐ ical resistance of the alkyd resin–based paints. These results are very important for practical

> **On the film surface**

**Degree of blistering ASTM D 714‐87**

**Paint film Steel panel Calculated anti‐**

**Metal surface corrosion (%)**

**Corrosion in the cut (mm) ASTMD 1654‐92**

**Around the ASTMD 610‐85**

5 2F ‐ 2.0−3.0 3 73 10 4M 6F 2.0−3.0 33 49 15 4MD 6F 1.0−2.0 50 42

5 4F ‐ 0.5−1.0 3 79 10 4MD ‐ 0.5−1.0 1 71 15 4MD 6F 0.5−1.0 16 54

5 6M ‐ 0.5−1.0 16 68 10 6F ‐ 0.5−1.0 16 68 15 2F 2F 1.0−2.0 33 45

5 6D 8M 0.5−1.0 33 48 10 6D 8M 0.5−1.0 >50 36 15 4D 8MD 0.5−1.0 >50 30

5 8D 8M 1.0−2.0 >50 35 10 8MD 8M 0.5−1.0 >50 43 15 6MD 8M 1.0−2.0 >50 39

5 4M 8M 1.0−2.0 10 58 10 6MD 8D 1.0−2.0 >50 29 15 6D 8D 0.5−1.0 >50 26

**Non‐pigmented paint film** 4MD 2F 1.0−2.0 33 45

**Table 3.** Corrosion resistance of the paints with the pigments synthesized and with the reference pigment and of the reference paint, measured after 912 hours exposure of the steel panels coated with the paints in an atmosphere with

**Table 3** lists the results of corrosion resistance measurements of the paint films in the atmosphere with condensed moisture and SO2, obtained after 912 hours of exposure. Nearly all of the paint films had osmotic blisters on their surface after the exposure. The paint films with pigment *C* (*lamellar Mg0.2Zn0.8Fe2O4*) were most resistant in this respect, especially at PVC =10%, where no blisters were found on the film surface and a few small blisters only (6F) were observed around the test cut. The results observed when the pigment concentration was PVC

**corrosive efficiency ESO2**

applications of these pigments in paints.

296 Magnetic Spinels- Synthesis, Properties and Applications

**(%)**

**cut**

**Pigment PVC** 

Mg0.2Zn0.8Fe2O4 isometric

Mg0.2Zn0.8Fe2O4 acicular

Mg0.2Zn0.8Fe2O4 lamellar

Mg‐Zn‐Fe mixed oxide/kaolin

Mg‐Zn‐Fe mixed oxide/talc

SO2, DFT = (85 ± 10) μm.

**3.4. Corrosion tests of the pigmented paints**

**A**

**B**

**C**

**D**

**E**

**F**

hematite Fe2O3

Similar behaviour was observed for the paint containing pigment *A* (*isometric Mg0.2Zn0.8*). Here, too, the resistance against blistering in the cut was poorer at higher concentrations: the scores at PVC = 5, 10 and 15% were 2F, 4M and 4MD, respectively.

The occurrence of blisters was more pronounced for the remaining paints. So, when the paint with the core‐shell pigment *E* (*mixed Mg‐Zn‐Fe oxide/talc*) was used, the paint film surface was speckled with very small blisters (frequency of occurrence 8M). This applies to the entire PVC range. The effects were very similar with the core‐shell pigment *D* (*mixed Mg‐Zn‐Fe oxide/ kaolin*), where, in addition, more abundant (score 6D) blisters were found near the test cut. The poorest results were obtained by using the paint pigmented with the reference pigment *F* (hematite *Fe2O3*) at PVC =10% and at PVC = 15%, where the paint film surface was densely covered with very small blisters (score 8D). The differences in the extent of blistering on the various paint films are apparent from **Figure 4**.

**Figure 4.** Photographs of the steel panels coated with paints containing: (a) pigment *F* (hematite *Fe2O3*) at PVC = 10%, (left); (b) pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*) at PVC = 5% (middle); and (c) pigment *C* (lamellar *Mg0.2Zn0.8Fe2O4*) at PVC = 10% (right).

The paint films were removed from the steel panels in order to examine corrosion on the metal surface (beneath the initially present films) and corrosion propagation from the test cut. Corrosion in the test cut was similar for all of the paints. Appreciable corrosion propagation from the cut, reaching a 3 mm distance, was only observed for the paint with the isometric pigment *A* (*isometric Mg0.2Zn0.8Fe2O4*) at PVC = 5% and PVC = 10%, followed by the paint with the isometric reference pigment *F* (1.0−2.0 mm). The degree of corrosion was higher in this atmosphere than in the atmosphere with condensed moisture for nearly all of the paints. The best protection was provided by the paint with pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*) at PVC = 10% and at PVC= 5%, the affected surface fraction being 1 and 3%, respectively. A comparable protection was provided by the paint with pigment *A* (*isometric Mg0.2Zn0.8Fe2O4*) at PVC = 5% (corroded surface fraction 3%). This pigment exhibited a sharp drop of its protective capacity against corrosion beneath the film with increasing PVC: the affected fraction was as high as 50% at PVC=10%.

A comparison of the steel panel corrosion beneath selected paint films is presented in **Figure 5**. The photograph on the right shows the steel panel initially protected by the paint with pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*) at PVC = 5%. The other two photographs show panels with appreciable degree of corrosion beneath the paint films.

**Figure 5.** Photographs of the steel panels after removing the paint films: (a) pigment *E* (mixed *Mg‐Zn‐Fe oxide/talc*) at PVC = 15% (left); (b) pigment *D* (mixed *Mg‐Zn‐Fe oxide/kaolin*) at PVC = 15% (middle); (c) pigment *B* (acicular *Mg0.2Zn0.8Fe2O4*) at PVC = 5% (right).

The paint that was found most efficient in protection against corrosion in the environment with condensed moisture and SO2 was that with the non‐isometric pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*). Paint films with this pigment were well resistant to blistering and efficiently protected the metal against corrosion. Favourably, this pigment is alkaline and can neutralize the strong acid aqueous medium acting on the paint film in the environment of the SO2 chamber whose acidity is about pH 2. In the optimum arrangement, the needle‐shaped pigment particles enhance cohesion of the paint and slow down diffusion of substances dissolved in the environment with condensing water [55]. The pigment is capable of forming soaps by reacting with the fatty acids (RCOOH) in the binder. Good resistance against blistering was also provided by the paint with pigment *C* (*lamellar Mg0.2Zn0.8Fe2O4*). At PVC = 15%, however, the protective capacity of this paint was as poor as that of the non‐pigmented film—the affected metal surface fraction was 33%. In the optimum arrangement, particularly at low PVCs, the lamellar pigment particles enhance paint adhesion to the metal and it is more difficult for the corrosive medium to diffuse through the paint film. Similar results were obtained with the paint containing pigment *A* (*isometric Mg0.2Zn0.8Fe2O4*).

The following conclusion can be drawn from the observations in this corrosive environment:


The overall anti‐corrosive efficiency was calculated for all the paint films tested. The most efficient in this respect was the paint with pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*) at PVC = 5% (ESO2 = 79) and at PVC = 10% (ESO2 =71), followed by the paint with pigment *C* (*lamellar Mg0.2Zn0.8Fe2O4*) at PVC= 5% and 10% (ESO2 = 68). Both pigment types exhibit a fairly uniform anti‐corrosive efficiency. The corrosive environment contains mobile SO2 ions [58] creating a strongly acid medium containing sulphite ions SO3 <sup>2</sup> <sup>−</sup> and sulphate ions SO4 2 − , and so a pigment with small particles exerting an alkaline effect is needed to attain a high corrosion resistance. Such a pigment will hinder diffusion through the paint film and neutralize the action of acid substances [52]. The paint with pigment *A* (*isometric Mg0.2Zn0.8Fe2O4*), was found efficient, its concentration, however, must be as low as PVC = 5% (ESO2 =73).

The paint properties are not so favourable, in other words, the paint is not so efficient, if a core‐ shell pigment with lamellar particles is used. Such a paint is even less efficient than the coating material containing no pigment at all. Ideal for a paint exposed to the environment in question is a pigment with needle‐shaped particles exerting an alkaline effect, which will also strengthen and reinforce the film.

## **4. Conclusions**

protection was provided by the paint with pigment *A* (*isometric Mg0.2Zn0.8Fe2O4*) at PVC = 5% (corroded surface fraction 3%). This pigment exhibited a sharp drop of its protective capacity against corrosion beneath the film with increasing PVC: the affected fraction was as high as

A comparison of the steel panel corrosion beneath selected paint films is presented in **Figure 5**. The photograph on the right shows the steel panel initially protected by the paint with pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*) at PVC = 5%. The other two photographs show panels

**Figure 5.** Photographs of the steel panels after removing the paint films: (a) pigment *E* (mixed *Mg‐Zn‐Fe oxide/talc*) at PVC = 15% (left); (b) pigment *D* (mixed *Mg‐Zn‐Fe oxide/kaolin*) at PVC = 15% (middle); (c) pigment *B* (acicular

The paint that was found most efficient in protection against corrosion in the environment with condensed moisture and SO2 was that with the non‐isometric pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*). Paint films with this pigment were well resistant to blistering and efficiently protected the metal against corrosion. Favourably, this pigment is alkaline and can neutralize the strong acid aqueous medium acting on the paint film in the environment of the SO2 chamber whose acidity is about pH 2. In the optimum arrangement, the needle‐shaped pigment particles enhance cohesion of the paint and slow down diffusion of substances dissolved in the environment with condensing water [55]. The pigment is capable of forming soaps by reacting with the fatty acids (RCOOH) in the binder. Good resistance against blistering was also provided by the paint with pigment *C* (*lamellar Mg0.2Zn0.8Fe2O4*). At PVC = 15%, however, the protective capacity of this paint was as poor as that of the non‐pigmented film—the affected metal surface fraction was 33%. In the optimum arrangement, particularly at low PVCs, the lamellar pigment particles enhance paint adhesion to the metal and it is more difficult for the corrosive medium to diffuse through the paint film. Similar results were obtained with the

with appreciable degree of corrosion beneath the paint films.

50% at PVC=10%.

298 Magnetic Spinels- Synthesis, Properties and Applications

*Mg0.2Zn0.8Fe2O4*) at PVC = 5% (right).

paint containing pigment *A* (*isometric Mg0.2Zn0.8Fe2O4*).

Ferro‐spinels (Zn‐Mg‐Fe) differing in the primary particle shapes were synthesized and added as pigments to a paint formula, and the paints containing them at various volume concentra‐ tions were investigated with respect to their anti‐corrosive properties. Every property of a pigmented paint is at its optimum at a specific pigment concentration; this applies particularly to the physical and anti‐corrosive properties [37]. So it is possible to identify a pigment concentration at which a specific property is at its best or at which the overall anti‐corrosive efficiency is at its maximum [59]. It is convenient that such pigment concentrations are usually not very high.

The spinels fall in the class of chemically acting pigments that help slow down corrosion processes on the metal surface beneath the paint film through their alkaline nature and by neutralization of the carboxy groups. The lamellar shape of the pigment particles enhances paint film adhesion to the substrate and its cohesion and reduces the formation of blisters on the paint film surface (and around the test cut). The pigmented paint films exhibited very good physico‐mechanical properties, commensurable with those of the alkyd resin alone.

The best anti‐corrosive efficiency in the accelerated corrosion test in the atmosphere with SO2 was found for the paint containing the non‐isometric pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*) at PVC = 5%.

It follows from the results that the morphology the pigment particles plays a major role in the paints' anti‐corrosive properties. Paints with non‐isometric pigments gave better results than paints with isometric pigments during nearly all measurements. This is primarily due to the barrier effect to substances penetrating through the protective film. Furthermore, the aggres‐ sive substances in the penetrating medium come in closer contact with the pigment particles, which may result in neutralization of the former. The best combined anti‐corrosive effect was found for pigment *C* with lamellar particles (lamellar *Mg0.2Zn0.8Fe2O4*). This pigment acted very favourably by the chemical inhibition mechanism; induced saponification in the alkyd resin owing to by its high alkalinity; and acted by the barrier effect. The core‐shell pigments *D* (mixed *Mg‐Zn‐Fe oxide/kaolin*) and *E* (mixed *Mg‐Zn‐Fe oxide/talc*), which also exhibit very good anti‐ corrosive efficiency in many parameters, may also be used as convenient variants: they are less expensive than pigment *C* and contain less zinc because the mineral core represents a consid‐ erable fraction of the pigment particle weight.

Paints with the pigments tested are also advantageous in that they protect the metal against corrosion. And they are environmentally friendly and might be favourably used to replace the toxic chromium (VI)‐based pigments.

## **Author details**

Kateřina Nechvílová, Andrea Kalendová\* and Miroslav Kohl

\*Address all correspondence to: andrea.kalendova@upce.cz

Faculty of Chemical Technology, University of Pardubice, Pardubice, Czech Republic

#### **References**

pigmented paint is at its optimum at a specific pigment concentration; this applies particularly to the physical and anti‐corrosive properties [37]. So it is possible to identify a pigment concentration at which a specific property is at its best or at which the overall anti‐corrosive efficiency is at its maximum [59]. It is convenient that such pigment concentrations are usually

The spinels fall in the class of chemically acting pigments that help slow down corrosion processes on the metal surface beneath the paint film through their alkaline nature and by neutralization of the carboxy groups. The lamellar shape of the pigment particles enhances paint film adhesion to the substrate and its cohesion and reduces the formation of blisters on the paint film surface (and around the test cut). The pigmented paint films exhibited very good

The best anti‐corrosive efficiency in the accelerated corrosion test in the atmosphere with SO2 was found for the paint containing the non‐isometric pigment *B* (*acicular Mg0.2Zn0.8Fe2O4*) at

It follows from the results that the morphology the pigment particles plays a major role in the paints' anti‐corrosive properties. Paints with non‐isometric pigments gave better results than paints with isometric pigments during nearly all measurements. This is primarily due to the barrier effect to substances penetrating through the protective film. Furthermore, the aggres‐ sive substances in the penetrating medium come in closer contact with the pigment particles, which may result in neutralization of the former. The best combined anti‐corrosive effect was found for pigment *C* with lamellar particles (lamellar *Mg0.2Zn0.8Fe2O4*). This pigment acted very favourably by the chemical inhibition mechanism; induced saponification in the alkyd resin owing to by its high alkalinity; and acted by the barrier effect. The core‐shell pigments *D* (mixed *Mg‐Zn‐Fe oxide/kaolin*) and *E* (mixed *Mg‐Zn‐Fe oxide/talc*), which also exhibit very good anti‐ corrosive efficiency in many parameters, may also be used as convenient variants: they are less expensive than pigment *C* and contain less zinc because the mineral core represents a consid‐

Paints with the pigments tested are also advantageous in that they protect the metal against corrosion. And they are environmentally friendly and might be favourably used to replace the

and Miroslav Kohl

Faculty of Chemical Technology, University of Pardubice, Pardubice, Czech Republic

physico‐mechanical properties, commensurable with those of the alkyd resin alone.

not very high.

300 Magnetic Spinels- Synthesis, Properties and Applications

PVC = 5%.

erable fraction of the pigment particle weight.

toxic chromium (VI)‐based pigments.

Kateřina Nechvílová, Andrea Kalendová\*

\*Address all correspondence to: andrea.kalendova@upce.cz

**Author details**


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## *Edited by Mohindar Singh Seehra*

Magnetic spinels including ferrites are insulating magnetic oxides and chalcogenides with strong coupling to microwave frequencies and low eddy current losses making them indispensable for applications in wireless communications. The 13 chapters and preface of this book discuss other potential applications of magnetic spinels along with various methods used for their synthesis and their varied properties resulting from substituting different metal ions at the A and B sites. These applications include ferrofluids, anticorrosion coatings, absorber coatings for photothermal conversion, biomedicine, and environmental applications such as oxidation of volatile organic compounds and removal of arsenic and heavy metals from water. Emphasis is placed on structure-property correlations and on the nature of magnetism in spinels and their nanoparticles with current information provided for future research.

Magnetic Spinels - Synthesis, Properties and Applications

Magnetic Spinels

Synthesis, Properties and Applications

*Edited by Mohindar Singh Seehra*

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