4. Properties of Ti1xAlxN coatings

Ti1xAlxN coatings are now extensively used for machining applications, where good mechanical and chemical properties are needed in order to ensure the stability of the coatings in inservice conditions. These properties are strongly dependent on the chemical composition of the coating and the working temperature, which could lead to a change in structure of the coating. Properties as hardness and oxidation resistance are then reviewed with respect to the aluminum content in the fcc-Ti1xAlxN phase and deposition temperature.

### 4.1. Mechanical properties

### 4.1.1. Hardness

### 4.1.1.1. Al-content dependence of hardness

Hardness of the coating is a key property for increasing the wear resistance and improving the tool lifetime of the cutting tools. The hardness of coatings is generally characterized by mean of nanohardness tests. Among the mechanical properties of Ti1xAlxN coatings, the hardness is only indirectly related to the aluminum content via the crystallographic structure. Thus, an increase of the aluminum content in the fcc-Ti1xAlxN phase leads to an increase of the hardness. The main mechanism is solid solution hardening: the substitution of Ti by Al atoms in the fcc-TiN leads to a lattice distortion and thus limits the motion of dislocations [28]. Above an aluminum content threshold, a mixture of two phases, the cubic and hexagonal structures, is deposited. The presence of the hcp phase, observed for about x = 0.7 in Ti1xAlxN, leads to a drop in hardness of the coatings [8, 23], as it can be seen in Figure 9. Highest aluminum contents lead, as specified earlier, to the deposition of hcp-Ti1xAlxN single-phased coatings, when a PVD method is used, associated with lower hardness. The hardness of some coatings deposited by PVD, CVD, and PECVD methods related to the structure and process parameters is summarized in the table given in Appendix A. It can be seen that PVD Ti1xAlxN monolayer coatings generally show hardness between 25 and 30 GPa. Concerning CVD processes, for which the aluminum content in the fcc-Ti1xAlxN solid solution can reach values up to x = 0.9, hardness up to 32.4 GPa was reported. Nanocomposite structures consisting of an alternate of fcc and hcp nanolamellae were found to have a relatively high hardness (≈26 GPa) compared to other hcp-AlNcontaining coatings [45]. Higher hardness (36 GPa) was then obtained for fcc-Ti1xAlxN/Alrich fcc-Ti1xAlxN nanocomposites [47]. PECVD techniques, due to their high out-of-equilibrium nature, also allow to deposit coatings which are fcc-Ti1xAlxN single-phased with x values up to ≈0.91, as reported by Prange et al. [28]. The highest hardness of single-phased coatings in this study was found for the fcc-Ti0.17Al0.83N with value of 38.7 GPa. Coatings with aluminum content higher than 0.91 were deposited and hardness up to 5000 HV (≈49 GPa) was reached. However, this huge increase in hardness is mainly attributed to the biphased nanocrystalline structure associated with a significant grain refinement. The large range of hardness observed for these different deposition processes can be explained by the high dependence of hardness to many parameters as the grain size (according to the Hall-Petch law), residual stresses, morphology, and structure of coatings [28, 84–87].

4.1.1.2. Temperature dependence of hardness

when Al content increases. This behavior is shown Figure 10.

unbalanced magnetron sputtering [23].

The nitrides, like TiN, show a sharp decrease in hardness at temperatures above 400C, because of mechanisms as restauration and recrystallization [19]. However, alloying with elements theoretically immiscible in TiN like aluminum allows changing completely the temperature dependence of the hardness. The evolution of the hardness as a function of temperature is related to the thermal stability of the material and to its decomposition process. The coherent spinodal nature of the decomposition of fcc-Ti1xAlxN can only lead to the formation of fcc-TiN and fcc-AlN with coherent interfaces, where the lattice mismatch allows hindering the dislocations' movement and so resulting in an increase of hardness. This behavior is responsible of the socalled age-hardening effect observed in Ti1xAlxN coatings. However, the peak of hardness is not observed at the same temperature for all coatings but depends also on the aluminum content in the films. Thus, an fcc-Ti1xAlxN coating with aluminum content close to the apparent solubility limit (for example, 0.7 for PVD coatings) has higher demixing energy than coatings with lower aluminum contents and has a trend to decompose more easily, as predicted by thermodynamical and ab initio calculations [29]. These calculations are in agreement with experimental studies showing that the age-hardening occurs earlier for Al-rich coatings. Chen et al. [23] have thus led a study of the mechanical properties and thermal stability of Ti1xAlxN coatings deposited by unbalanced magnetron sputtering process at 500C, 0.4 Pa, and bias voltage of 60 V. They clearly show that the temperature corresponding to the maximum hardness decreases

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Figure 10. Evolution of hardness vs. annealing temperature for coatings with different aluminum contents deposited by

Figure 9. Evolution of hardness vs. aluminum content in coatings deposited by cathodic arc ion plating [25].

### 4.1.1.2. Temperature dependence of hardness

dislocations [28]. Above an aluminum content threshold, a mixture of two phases, the cubic and hexagonal structures, is deposited. The presence of the hcp phase, observed for about x = 0.7 in Ti1xAlxN, leads to a drop in hardness of the coatings [8, 23], as it can be seen in Figure 9. Highest aluminum contents lead, as specified earlier, to the deposition of hcp-Ti1xAlxN single-phased coatings, when a PVD method is used, associated with lower hardness. The hardness of some coatings deposited by PVD, CVD, and PECVD methods related to the structure and process parameters is summarized in the table given in Appendix A. It can be seen that PVD Ti1xAlxN monolayer coatings generally show hardness between 25 and 30 GPa. Concerning CVD processes, for which the aluminum content in the fcc-Ti1xAlxN solid solution can reach values up to x = 0.9, hardness up to 32.4 GPa was reported. Nanocomposite structures consisting of an alternate of fcc and hcp nanolamellae were found to have a relatively high hardness (≈26 GPa) compared to other hcp-AlNcontaining coatings [45]. Higher hardness (36 GPa) was then obtained for fcc-Ti1xAlxN/Alrich fcc-Ti1xAlxN nanocomposites [47]. PECVD techniques, due to their high out-of-equilibrium nature, also allow to deposit coatings which are fcc-Ti1xAlxN single-phased with x values up to ≈0.91, as reported by Prange et al. [28]. The highest hardness of single-phased coatings in this study was found for the fcc-Ti0.17Al0.83N with value of 38.7 GPa. Coatings with aluminum content higher than 0.91 were deposited and hardness up to 5000 HV (≈49 GPa) was reached. However, this huge increase in hardness is mainly attributed to the biphased nanocrystalline structure associated with a significant grain refinement. The large range of hardness observed for these different deposition processes can be explained by the high dependence of hardness to many parameters as the grain size (according to the Hall-

186 Coatings and Thin-Film Technologies

Petch law), residual stresses, morphology, and structure of coatings [28, 84–87].

Figure 9. Evolution of hardness vs. aluminum content in coatings deposited by cathodic arc ion plating [25].

The nitrides, like TiN, show a sharp decrease in hardness at temperatures above 400C, because of mechanisms as restauration and recrystallization [19]. However, alloying with elements theoretically immiscible in TiN like aluminum allows changing completely the temperature dependence of the hardness. The evolution of the hardness as a function of temperature is related to the thermal stability of the material and to its decomposition process. The coherent spinodal nature of the decomposition of fcc-Ti1xAlxN can only lead to the formation of fcc-TiN and fcc-AlN with coherent interfaces, where the lattice mismatch allows hindering the dislocations' movement and so resulting in an increase of hardness. This behavior is responsible of the socalled age-hardening effect observed in Ti1xAlxN coatings. However, the peak of hardness is not observed at the same temperature for all coatings but depends also on the aluminum content in the films. Thus, an fcc-Ti1xAlxN coating with aluminum content close to the apparent solubility limit (for example, 0.7 for PVD coatings) has higher demixing energy than coatings with lower aluminum contents and has a trend to decompose more easily, as predicted by thermodynamical and ab initio calculations [29]. These calculations are in agreement with experimental studies showing that the age-hardening occurs earlier for Al-rich coatings. Chen et al. [23] have thus led a study of the mechanical properties and thermal stability of Ti1xAlxN coatings deposited by unbalanced magnetron sputtering process at 500C, 0.4 Pa, and bias voltage of 60 V. They clearly show that the temperature corresponding to the maximum hardness decreases when Al content increases. This behavior is shown Figure 10.

Figure 10. Evolution of hardness vs. annealing temperature for coatings with different aluminum contents deposited by unbalanced magnetron sputtering [23].

It is also important to note that an enhancement of the hardness with temperature was also observed in this study for an hcp-Ti0.25Al0.75N single-phased film [23]. Based on XRD analysis, the authors proposed that this hardness enhancement is related to the formation of fcc-TiN precipitates. This age-hardening effect can also be observed for nanocomposite coatings with high hardness stable up to 1100C, while the "standard" fcc-Ti1xAlxN solid solution coatings show a significant drop of hardness for temperatures above 1000C [45, 47].

### 4.1.2. Residual stresses

Residual stresses act strongly on the mechanical properties of coatings. Compressive stresses lead to better wear resistance, especially for adhesive wear at high temperature, by reducing diffusion in the coating. However, a too high level of compressive stress can lead to cohesive and/or adhesive spalling. For this reason, studies aim generally at obtaining coatings with the smallest residual stresses. CVD coatings show generally lower residual stresses than PVD coatings because of the atomic peening of the growing film, which leads to local densification and so, to additional compressive stresses in PVD coatings.

Residual stresses are decomposed into two contributions: the thermal stresses, related to the difference of coefficient of thermal expansion between the substrate and the coating; and the intrinsic stresses, related to the process parameters and to morphology. The flux and kinetic energy of species in the PVD processes, particularly those involving high ionization discharges, can generate a local densification of the coating (atomic peening), leading to compressive stresses, while, under low bombardment conditions, a columnar growth is associated to tensile or low compressive stresses, depending on the intercolumnar distance, according to the Hoffman model [88].

PVD coatings deposited on WC/Co carbides show compressive stresses between 1 and 5 GPa [26, 90] and values above 5 GPa are reported to cause delamination of the film deposited by cathodic arc deposition [90]. Concerning thermal CVD coatings, Endler et al., who deposited fcc-Ti1xAlxN single-phased films with x = 0.82 on WC/Co obtained compressive stress less than 1 Gpa, partly due to the presence of a TiN sublayer [27]. The residual stresses measured on coatings obtained by different deposition techniques are summarized in the table

Figure 11. Evolution of the residual stresses vs. aluminum content in Ti1xAlxN thin films deposited by reactive unbal-

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During machining operations, cutting tools are subjected, in addition to mechanical stresses, to high temperatures, that can reach 1000C [92], which can lead to severe oxidation and consequently to altered properties of the film. The oxidation resistance of Ti1xAlxN coatings is generally associated to the formation of a dense Al2O3 diffusion barrier layer that protects Ti1xAlxN against further oxidation. However, oxidation mechanisms are complex and the evolution of the structures as a function of the aluminum content has a strong influence on the oxidation behavior of the coatings. Nevertheless, although oxidation mechanisms are not now completely defined, it is widely accepted that the formation of the oxide phase is the consequence of the outward diffusion of Al associated with the inward diffusion of O [7, 93–95]. Moreover, analysis performed by RBS (and later by NRA [95]) shows that oxidation is associated to a nitrogen loss. This suggests that, in the first steps of oxidation, N is substituted by O, and oxides and/or oxinitrides such as (Ti, Al)NxOy are formed. With further oxidation, the composition of the system evolves progressively toward the pure Al2O3 and TiO2 layers as a function of oxidation parameters. With increasing oxidation time or temperature, Ti begins to oxidize to form rutile r-TiO2. The formation of this oxide can lead to crack formation in the Al2O3 protective layer, acting as diffusion paths for oxygen and leading to the complete oxidation of the coatings. Titanium atoms, by passing through these cracks, lead to the formation

in Appendix A.

4.2. Oxidation resistance

anced magnetron sputtering [91].

4.2.1. Overall mechanisms

Residual stresses are generally measured using the substrates' curvature technique and calculated by the Stoney's formula. Values obtained in different studies are difficult to compare because of the strong dependence to many factors. First, the "material" factors such as thickness of coating, texture, presence of underlayer as TiN [27], TiAl, Ti1xAlxN with gradual Al content [89], or nature of substrate influence strongly the residual stresses after deposition. Moreover, process parameters also play a role in the evolution of stresses, as deposition temperature or bias voltage [90]. Ahlgren et al. [90] have thus shown that increasing the bias voltage during the deposition of arc deposited Ti1xAlxN leads to a significant increase of the residual stresses in the coating. Stresses of 1.69 GPa were then obtained for a bias voltage of 40 V, while values of 5.59 GPa were found at bias voltage of 200 V.

Shum et al. led a structural and mechanical study of Ti1xAlxN films with respect to the Al content [91]. The films were deposited by reactive close-field unbalanced magnetron sputtering and high compressive stresses (≈ 3 GPa) were measured for TiN. This result was related to the low deposition pressure of 0.27 Pa, which favors an intense bombardment of the growing film. They found that Al incorporation up to x = 0.41 leads to a significant decrease of the residual stresses (Figure 11). For this Al content, the lowest compressive stresses of about 0.35 GPa were measured and they attributed these results to the presence of amorphous AlN around Ti1xAlxN grains. When aluminum content is further increased, the residual stresses rise up to value of ≈ 1.5 GPa for pure AlN (Figure 11). This behavior was linked to the increasing fraction of hcp-AlN.

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Figure 11. Evolution of the residual stresses vs. aluminum content in Ti1xAlxN thin films deposited by reactive unbalanced magnetron sputtering [91].

PVD coatings deposited on WC/Co carbides show compressive stresses between 1 and 5 GPa [26, 90] and values above 5 GPa are reported to cause delamination of the film deposited by cathodic arc deposition [90]. Concerning thermal CVD coatings, Endler et al., who deposited fcc-Ti1xAlxN single-phased films with x = 0.82 on WC/Co obtained compressive stress less than 1 Gpa, partly due to the presence of a TiN sublayer [27]. The residual stresses measured on coatings obtained by different deposition techniques are summarized in the table in Appendix A.

#### 4.2. Oxidation resistance

It is also important to note that an enhancement of the hardness with temperature was also observed in this study for an hcp-Ti0.25Al0.75N single-phased film [23]. Based on XRD analysis, the authors proposed that this hardness enhancement is related to the formation of fcc-TiN precipitates. This age-hardening effect can also be observed for nanocomposite coatings with high hardness stable up to 1100C, while the "standard" fcc-Ti1xAlxN solid solution coatings

Residual stresses act strongly on the mechanical properties of coatings. Compressive stresses lead to better wear resistance, especially for adhesive wear at high temperature, by reducing diffusion in the coating. However, a too high level of compressive stress can lead to cohesive and/or adhesive spalling. For this reason, studies aim generally at obtaining coatings with the smallest residual stresses. CVD coatings show generally lower residual stresses than PVD coatings because of the atomic peening of the growing film, which leads to local densification

Residual stresses are decomposed into two contributions: the thermal stresses, related to the difference of coefficient of thermal expansion between the substrate and the coating; and the intrinsic stresses, related to the process parameters and to morphology. The flux and kinetic energy of species in the PVD processes, particularly those involving high ionization discharges, can generate a local densification of the coating (atomic peening), leading to compressive stresses, while, under low bombardment conditions, a columnar growth is associated to tensile or low compressive stresses, depending on the intercolumnar distance, according to the

Residual stresses are generally measured using the substrates' curvature technique and calculated by the Stoney's formula. Values obtained in different studies are difficult to compare because of the strong dependence to many factors. First, the "material" factors such as thickness of coating, texture, presence of underlayer as TiN [27], TiAl, Ti1xAlxN with gradual Al content [89], or nature of substrate influence strongly the residual stresses after deposition. Moreover, process parameters also play a role in the evolution of stresses, as deposition temperature or bias voltage [90]. Ahlgren et al. [90] have thus shown that increasing the bias voltage during the deposition of arc deposited Ti1xAlxN leads to a significant increase of the residual stresses in the coating. Stresses of 1.69 GPa were then obtained for a bias voltage of

Shum et al. led a structural and mechanical study of Ti1xAlxN films with respect to the Al content [91]. The films were deposited by reactive close-field unbalanced magnetron sputtering and high compressive stresses (≈ 3 GPa) were measured for TiN. This result was related to the low deposition pressure of 0.27 Pa, which favors an intense bombardment of the growing film. They found that Al incorporation up to x = 0.41 leads to a significant decrease of the residual stresses (Figure 11). For this Al content, the lowest compressive stresses of about 0.35 GPa were measured and they attributed these results to the presence of amorphous AlN around Ti1xAlxN grains. When aluminum content is further increased, the residual stresses rise up to value of ≈ 1.5 GPa for pure AlN (Figure 11). This behavior was linked to the

40 V, while values of 5.59 GPa were found at bias voltage of 200 V.

show a significant drop of hardness for temperatures above 1000C [45, 47].

and so, to additional compressive stresses in PVD coatings.

4.1.2. Residual stresses

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Hoffman model [88].

increasing fraction of hcp-AlN.

#### 4.2.1. Overall mechanisms

During machining operations, cutting tools are subjected, in addition to mechanical stresses, to high temperatures, that can reach 1000C [92], which can lead to severe oxidation and consequently to altered properties of the film. The oxidation resistance of Ti1xAlxN coatings is generally associated to the formation of a dense Al2O3 diffusion barrier layer that protects Ti1xAlxN against further oxidation. However, oxidation mechanisms are complex and the evolution of the structures as a function of the aluminum content has a strong influence on the oxidation behavior of the coatings. Nevertheless, although oxidation mechanisms are not now completely defined, it is widely accepted that the formation of the oxide phase is the consequence of the outward diffusion of Al associated with the inward diffusion of O [7, 93–95]. Moreover, analysis performed by RBS (and later by NRA [95]) shows that oxidation is associated to a nitrogen loss. This suggests that, in the first steps of oxidation, N is substituted by O, and oxides and/or oxinitrides such as (Ti, Al)NxOy are formed. With further oxidation, the composition of the system evolves progressively toward the pure Al2O3 and TiO2 layers as a function of oxidation parameters. With increasing oxidation time or temperature, Ti begins to oxidize to form rutile r-TiO2. The formation of this oxide can lead to crack formation in the Al2O3 protective layer, acting as diffusion paths for oxygen and leading to the complete oxidation of the coatings. Titanium atoms, by passing through these cracks, lead to the formation of a porous and detrimental r-TiO2 top layer [40]. The oxidation behavior is thus mainly dependent on the aluminum content in the coating and on the oxidation temperature.

are assumed to be almost "pure" Al2O3 and TiO2. However, with increasing x in the asdeposited fcc-Ti1xAlxN coatings, higher aluminum contents were found in the Ti-rich interlayer. Thus, this Ti-rich layer cannot be attributed to TiO2 but rather to a mixed aluminum-titanium oxide such as (Ti0.9Al0.1)Ox oxide. Other studies also report that the oxide layer formed during oxidation at high temperature (above 800C) is mainly composed of an AlxTiyOz oxide [7, 23]. These results are consistent with the higher oxidation resistance observed for Al-rich fcc-Ti1xAlxN and it suggests that higher aluminum concentration, by the formation of a dense AlO3 top layer, hinders oxygen diffusion and limits the formation of

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Increasing aluminum content leads to the formation of the fcc/hcp biphased structure. Many studies show that this material has a poor oxidation resistance [24, 25]. A further increase of the aluminum content leads to the stabilization of the hcp-Ti1xAlxN or hcp-AlN (depending on the deposition temperature). These coatings were found by Vaz et al. to be totally composed of the AlxTiyOz oxide at 850C. Although Ti1xAlxN coatings with hcp single-phased structure are generally found to have the worse oxidation resistance [24, 95], Chen et al. found an unexpected high oxidation resistance for hcp-Ti0.25Al0.75N single-phased coatings [23]. CVD or PECVD coatings with fcc single-phased structure obtained for x above 0.7 seem to have a better resistance than PVD coatings (with lower aluminum content). Thus, Prange et al. found no presence of the detrimental TiO2 for Ti0.21Al0.79N coatings oxidized at 800C during 1 h [28], while Ti0.18Al0.82N coatings deposited by LPCVD were found to oxidize only at temperatures above 900C [27]. Similar results were obtained for hcp/fcc Ti0.05Al0.95N nanocomposite coat-

According to the results of Vaz et al. [24], coatings oxidized at lowest temperatures (below 600C) are only composed of titanium-aluminum oxinitrides, while, for coatings oxidized at higher temperatures, Ti1xAlxOy oxides are observed [96]. These experiments confirm that, at the lowest temperatures, only a part of the nitrogen atoms are "substituted" by oxygen and escapes through the gas phase by forming N2. The Ti and Al contents were also found to be the same than that of the as-deposited coatings, annealed at low temperature. However, high temperatures favor the formation of Al2O3 and the Al and Ti contents in this oxide layer tend

The wear aspects take a large place in the research on hard coatings. During machining, wear is the consequence of a combination of mechanical and chemical solicitations. The wear mechanisms are also very dependent on the type of machining, tool geometry, material to be machined, materials used for tool and coatings. Thus, this part is focused only on the description of the general wear mechanisms needed for the understanding of the wear properties of

Ti1xAlxN coatings rather than an exhaustive description of these mechanisms.

ings, with the first signs of Al2O3 and TiO2 at a temperature of 950C [45].

obviously to be different than that of the as-deposited coatings.

4.2.3. Temperature dependence of oxidation resistance

4.3. Wear resistance

4.3.1. General wear mechanisms

the porous and detrimental r-TiO2.

#### 4.2.2. Aluminum content dependence of oxidation resistance

Most of the studies on Ti1xAlxN coatings focus on the necessity to increase aluminum incorporation in the coatings in order to promote better oxidation resistance. Actually, a continuous decrease of the oxide thickness is observed with increasing aluminum content in the coating [23], which indicates that the oxide layer prevents the coating from further oxidation. However, the oxidation resistance seems to be more related to the crystallographic structure (fcc, hcp, or mixed fcc/hcp), which is linked to the aluminum content, rather than a direct dependence to aluminum content.

Vaz et al. have summarized the evolution of the formed oxide layer regarding the temperature and aluminum content (Figure 12). It was shown that oxidation resistance increases with increasing the aluminum content in the coating as long as the films remain fcc single-phased (up to x = 0.65 in this study) [24]. These results are in accordance with later experiments where oxide thickness decreased with increasing aluminum content [23, 95]. According to the results of Vaz et al., although a small amount of Ti and Al remains in the Al-rich top layer and Ti-rich sublayer, respectively, the oxides formed at high temperatures and for low aluminum content

Figure 12. Schematic diagram summarized the temperature and aluminum content dependences of Ti1xAlxN coatings oxidation [24].

are assumed to be almost "pure" Al2O3 and TiO2. However, with increasing x in the asdeposited fcc-Ti1xAlxN coatings, higher aluminum contents were found in the Ti-rich interlayer. Thus, this Ti-rich layer cannot be attributed to TiO2 but rather to a mixed aluminum-titanium oxide such as (Ti0.9Al0.1)Ox oxide. Other studies also report that the oxide layer formed during oxidation at high temperature (above 800C) is mainly composed of an AlxTiyOz oxide [7, 23]. These results are consistent with the higher oxidation resistance observed for Al-rich fcc-Ti1xAlxN and it suggests that higher aluminum concentration, by the formation of a dense AlO3 top layer, hinders oxygen diffusion and limits the formation of the porous and detrimental r-TiO2.

Increasing aluminum content leads to the formation of the fcc/hcp biphased structure. Many studies show that this material has a poor oxidation resistance [24, 25]. A further increase of the aluminum content leads to the stabilization of the hcp-Ti1xAlxN or hcp-AlN (depending on the deposition temperature). These coatings were found by Vaz et al. to be totally composed of the AlxTiyOz oxide at 850C. Although Ti1xAlxN coatings with hcp single-phased structure are generally found to have the worse oxidation resistance [24, 95], Chen et al. found an unexpected high oxidation resistance for hcp-Ti0.25Al0.75N single-phased coatings [23]. CVD or PECVD coatings with fcc single-phased structure obtained for x above 0.7 seem to have a better resistance than PVD coatings (with lower aluminum content). Thus, Prange et al. found no presence of the detrimental TiO2 for Ti0.21Al0.79N coatings oxidized at 800C during 1 h [28], while Ti0.18Al0.82N coatings deposited by LPCVD were found to oxidize only at temperatures above 900C [27]. Similar results were obtained for hcp/fcc Ti0.05Al0.95N nanocomposite coatings, with the first signs of Al2O3 and TiO2 at a temperature of 950C [45].

### 4.2.3. Temperature dependence of oxidation resistance

According to the results of Vaz et al. [24], coatings oxidized at lowest temperatures (below 600C) are only composed of titanium-aluminum oxinitrides, while, for coatings oxidized at higher temperatures, Ti1xAlxOy oxides are observed [96]. These experiments confirm that, at the lowest temperatures, only a part of the nitrogen atoms are "substituted" by oxygen and escapes through the gas phase by forming N2. The Ti and Al contents were also found to be the same than that of the as-deposited coatings, annealed at low temperature. However, high temperatures favor the formation of Al2O3 and the Al and Ti contents in this oxide layer tend obviously to be different than that of the as-deposited coatings.

#### 4.3. Wear resistance

of a porous and detrimental r-TiO2 top layer [40]. The oxidation behavior is thus mainly

Most of the studies on Ti1xAlxN coatings focus on the necessity to increase aluminum incorporation in the coatings in order to promote better oxidation resistance. Actually, a continuous decrease of the oxide thickness is observed with increasing aluminum content in the coating [23], which indicates that the oxide layer prevents the coating from further oxidation. However, the oxidation resistance seems to be more related to the crystallographic structure (fcc, hcp, or mixed fcc/hcp), which is linked to the aluminum content, rather than a direct depen-

Vaz et al. have summarized the evolution of the formed oxide layer regarding the temperature and aluminum content (Figure 12). It was shown that oxidation resistance increases with increasing the aluminum content in the coating as long as the films remain fcc single-phased (up to x = 0.65 in this study) [24]. These results are in accordance with later experiments where oxide thickness decreased with increasing aluminum content [23, 95]. According to the results of Vaz et al., although a small amount of Ti and Al remains in the Al-rich top layer and Ti-rich sublayer, respectively, the oxides formed at high temperatures and for low aluminum content

Figure 12. Schematic diagram summarized the temperature and aluminum content dependences of Ti1xAlxN coatings

dependent on the aluminum content in the coating and on the oxidation temperature.

4.2.2. Aluminum content dependence of oxidation resistance

dence to aluminum content.

190 Coatings and Thin-Film Technologies

oxidation [24].

### 4.3.1. General wear mechanisms

The wear aspects take a large place in the research on hard coatings. During machining, wear is the consequence of a combination of mechanical and chemical solicitations. The wear mechanisms are also very dependent on the type of machining, tool geometry, material to be machined, materials used for tool and coatings. Thus, this part is focused only on the description of the general wear mechanisms needed for the understanding of the wear properties of Ti1xAlxN coatings rather than an exhaustive description of these mechanisms.

lower values are reported to be due to the very thin oxide layer of several nanometers resulting from the elevation of the temperature at the contact interface between tool and workpiece.

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Regarding the work of Shum et al. [91], addition of a small amount of aluminum in TiN leads to a huge increase of the COF (up to x ≈ 0.6). A further increase of the aluminum content does not seem to strongly affect the COF, which stabilizes around 0.5 for x = 0.6. These results are confirmed for high aluminum contents in PACVD films [28]. In this study, both monophased

The temperature dependence of the COF does not seem to be clearly defined. Ohnuma et al. thus studied friction of Ti0.6Al0.4N, Ti0.42Al0.58N, and Ti0.3Al0.7N coatings with fcc, fcc + hcp, and hcp structures, respectively, against AISI 304 material [109]. Measurements were performed at temperatures between 300 and 873 K. All coatings showed a COF between 0.6 and 0.9. Although a minimum value of 0.6 was found at 473 K for all coatings, no clear correlation between COF and temperature was found and the evolution of the COF vs.

Wear behavior is generally determined by studying the wear track obtained by pin-on-disc method [109] or multipass scratch test [91]. The results are highly dependent on many factors such as material characteristics (substrate, counterbody), deposition parameters, and testing parameters (critical loading force). It is thus difficult to make an exhaustive state of the art of

Tests at room temperature seem to be highly related to the structure of the coatings. Studies show generally an increase of the wear resistance with increasing aluminum content in the fcc-Ti1xAlxN single-phased coatings, while the coatings with an fcc/hcp biphased structure, or only with the hcp structure, have a lower wear resistance [26, 64, 109, 110]. These results are consistent with the evolution of the hardness as a function of the aluminum content in the Ti1xAlxN coatings, influencing strongly the abrasive wear resistance of the coating. Tests performed at high temperatures show that Ti1xAlxN coatings have a better wear resistance than that of TiN. Ohnuma found thus a diminution of the wear track depth for temperatures above 600C for Ti1xAlxN coatings whatever the structure, as shown in Figure 14. This behavior was related to the formation of the Al2O3 oxide layer, known to prevent adhesive

=E<sup>∗</sup><sup>2</sup>

0.1 [87] for Ti1xAlxN coatings, allowing to have better erosion resistance than other common

Even though the previously defined wear tests allow to estimate the wear behavior of Ti1xAlxN coatings, machining tests are generally performed to determine accurately the inservice behavior of the coatings. The combined effects of several factors as stress and temperature, as well as tool geometry, can act strongly on the wear behavior of the coated tool. In the case of Ti1xAlxN coatings, only few literature is available on the influence of aluminum

, measurements show generally values above

the wear behavior. However, some overall results or trends can be identified.

films (x = 0.79) and biphased coatings (x = 0.96) have a COF around 0.5.

temperature could not be explained.

4.3.3. Wear properties of Ti1xAlxN coatings

wear [109]. Regarding the plasticity index H<sup>3</sup>

nitrides such as TiN, CrN, or CrAlN [101].

4.3.4. In-service behavior of Ti1xAlxN coatings

Figure 13. Three-step evolution of the flank wear of drills during drilling operations [97].

Wear can be divided into three steps, illustrated in Figure 13 and showing the evolution of the flank wear of twist drills. Firstly, roughness can cause abrasive wear and the asperities are "removed" after welding to the cutting tool. This step is related to a high friction coefficient [64] associated with high wear rates [97]. The second step is characterized by both adhesive and abrasive wears. With increasing the cutting time, the continuous removal of workpiece particles by adhesive wear acts as a third-body, leading to abrasive wear. In this stage, the wear slowly increases linearly with time. The last stage is characterized by high wear rates leading to failure of the cutting tool [98]. The use of coatings as Ti1�xAlxN, because of their high hardness and diffusion barrier properties, allows to limit the abrasive and adhesive wears. The Al2O3 dense layer formed during cutting operations is also considered as the main factor responsible for the wear resistance of Ti1�xAlxN coatings.

However, in addition to the wear mechanisms previously defined, coated tools are subjected to peening due to mechanical and thermal stresses, leading to the appearance of fatigue cracks that grow until coating failure [99]. Resistance to cracking of a coating can be estimated by the plastic deformation ratio, given by H<sup>3</sup> <sup>=</sup>E<sup>∗</sup><sup>²</sup> with <sup>E</sup><sup>∗</sup> <sup>¼</sup> <sup>E</sup><sup>=</sup> <sup>1</sup> � <sup>ν</sup><sup>2</sup> , where <sup>H</sup> is the hardness, E\*, the effective Young modulus, E, the young modulus, and ν, the Poisson ratio of the coating. Better resistance to cracking could be obtained for the highest H<sup>3</sup> =E<sup>∗</sup><sup>²</sup> values [78, 87, 100, 101]. Further details on wear mechanisms are given in [99, 102–105].

### 4.3.2. Coefficient of friction (COF)

The friction coefficient plays a key role in the first step of wear mechanisms, where a high wear rate is observed. However, the friction coefficient is dependent of many parameters such as the nature of the workpiece material, loading force, or growth direction of the coating [106, 107]. For these reasons, results available in literature show a high dispersion. Then, values of 1.4 were found for arc-deposited TiN sliding against M2 steel [108]. High values of about 1 were also obtained for CVD TiN coatings sliding against Al2O3. However, lower values of 0.3 to 0.45 were found for unbalance magnetron sputtering TiN against M42 steel [91]. These lower values are reported to be due to the very thin oxide layer of several nanometers resulting from the elevation of the temperature at the contact interface between tool and workpiece.

Regarding the work of Shum et al. [91], addition of a small amount of aluminum in TiN leads to a huge increase of the COF (up to x ≈ 0.6). A further increase of the aluminum content does not seem to strongly affect the COF, which stabilizes around 0.5 for x = 0.6. These results are confirmed for high aluminum contents in PACVD films [28]. In this study, both monophased films (x = 0.79) and biphased coatings (x = 0.96) have a COF around 0.5.

The temperature dependence of the COF does not seem to be clearly defined. Ohnuma et al. thus studied friction of Ti0.6Al0.4N, Ti0.42Al0.58N, and Ti0.3Al0.7N coatings with fcc, fcc + hcp, and hcp structures, respectively, against AISI 304 material [109]. Measurements were performed at temperatures between 300 and 873 K. All coatings showed a COF between 0.6 and 0.9. Although a minimum value of 0.6 was found at 473 K for all coatings, no clear correlation between COF and temperature was found and the evolution of the COF vs. temperature could not be explained.

## 4.3.3. Wear properties of Ti1xAlxN coatings

Wear can be divided into three steps, illustrated in Figure 13 and showing the evolution of the flank wear of twist drills. Firstly, roughness can cause abrasive wear and the asperities are "removed" after welding to the cutting tool. This step is related to a high friction coefficient [64] associated with high wear rates [97]. The second step is characterized by both adhesive and abrasive wears. With increasing the cutting time, the continuous removal of workpiece particles by adhesive wear acts as a third-body, leading to abrasive wear. In this stage, the wear slowly increases linearly with time. The last stage is characterized by high wear rates leading to failure of the cutting tool [98]. The use of coatings as Ti1�xAlxN, because of their high hardness and diffusion barrier properties, allows to limit the abrasive and adhesive wears. The Al2O3 dense layer formed during cutting operations is also considered as the main factor

Figure 13. Three-step evolution of the flank wear of drills during drilling operations [97].

However, in addition to the wear mechanisms previously defined, coated tools are subjected to peening due to mechanical and thermal stresses, leading to the appearance of fatigue cracks that grow until coating failure [99]. Resistance to cracking of a coating can be estimated by the

the effective Young modulus, E, the young modulus, and ν, the Poisson ratio of the coating.

The friction coefficient plays a key role in the first step of wear mechanisms, where a high wear rate is observed. However, the friction coefficient is dependent of many parameters such as the nature of the workpiece material, loading force, or growth direction of the coating [106, 107]. For these reasons, results available in literature show a high dispersion. Then, values of 1.4 were found for arc-deposited TiN sliding against M2 steel [108]. High values of about 1 were also obtained for CVD TiN coatings sliding against Al2O3. However, lower values of 0.3 to 0.45 were found for unbalance magnetron sputtering TiN against M42 steel [91]. These

<sup>=</sup>E<sup>∗</sup><sup>²</sup> with <sup>E</sup><sup>∗</sup> <sup>¼</sup> <sup>E</sup><sup>=</sup> <sup>1</sup> � <sup>ν</sup><sup>2</sup> , where <sup>H</sup> is the hardness, E\*,

=E<sup>∗</sup><sup>²</sup> values [78, 87, 100, 101].

responsible for the wear resistance of Ti1�xAlxN coatings.

Better resistance to cracking could be obtained for the highest H<sup>3</sup>

Further details on wear mechanisms are given in [99, 102–105].

plastic deformation ratio, given by H<sup>3</sup>

192 Coatings and Thin-Film Technologies

4.3.2. Coefficient of friction (COF)

Wear behavior is generally determined by studying the wear track obtained by pin-on-disc method [109] or multipass scratch test [91]. The results are highly dependent on many factors such as material characteristics (substrate, counterbody), deposition parameters, and testing parameters (critical loading force). It is thus difficult to make an exhaustive state of the art of the wear behavior. However, some overall results or trends can be identified.

Tests at room temperature seem to be highly related to the structure of the coatings. Studies show generally an increase of the wear resistance with increasing aluminum content in the fcc-Ti1xAlxN single-phased coatings, while the coatings with an fcc/hcp biphased structure, or only with the hcp structure, have a lower wear resistance [26, 64, 109, 110]. These results are consistent with the evolution of the hardness as a function of the aluminum content in the Ti1xAlxN coatings, influencing strongly the abrasive wear resistance of the coating. Tests performed at high temperatures show that Ti1xAlxN coatings have a better wear resistance than that of TiN. Ohnuma found thus a diminution of the wear track depth for temperatures above 600C for Ti1xAlxN coatings whatever the structure, as shown in Figure 14. This behavior was related to the formation of the Al2O3 oxide layer, known to prevent adhesive wear [109]. Regarding the plasticity index H<sup>3</sup> =E<sup>∗</sup><sup>2</sup> , measurements show generally values above 0.1 [87] for Ti1xAlxN coatings, allowing to have better erosion resistance than other common nitrides such as TiN, CrN, or CrAlN [101].

### 4.3.4. In-service behavior of Ti1xAlxN coatings

Even though the previously defined wear tests allow to estimate the wear behavior of Ti1xAlxN coatings, machining tests are generally performed to determine accurately the inservice behavior of the coatings. The combined effects of several factors as stress and temperature, as well as tool geometry, can act strongly on the wear behavior of the coated tool. In the case of Ti1xAlxN coatings, only few literature is available on the influence of aluminum

machining of high-speed steel, Ni-based superalloys, aluminum and titanium alloys are

Ti-Al-N-Based Hard Coatings: Thermodynamical Background, CVD Deposition, and Properties. A Review

http://dx.doi.org/10.5772/intechopen.79747

195

In this chapter, the thermodynamical basis, deposition, and some mechanical and chemical properties of coatings based on Ti1xAlxN and Ti1xAlxSiyN have been reviewed. The general mechanisms (hardening effect, oxidation mechanisms, and wear mechanisms) associated to these properties were also discussed by pointing out the influence of the temperature and the aluminum content in the films. Increasing aluminum content induces structural changes that are directly associated with the properties. The high thermal stability and the age-hardening effect, observed in titanium-aluminum nitride-based coatings, are related to the coherent spinodal decomposition in the Ti-Al-N system. Concerning elaboration by means of CVD based processes, the stabilization of the fcc-Ti1xAlxN metastable solid solution can be achieved by a combination of relatively low deposition temperatures, high partial pressures of reactive precursors, and low total pressure. As-deposited monolayer coatings show a higher hardness and a better oxidation resistance when aluminum content is increased. Hardness up to about 30 GPa and stability in air up to 800C have been measured. These values outperform the typical values of about 25 GPa for hardness and 500C for the temperature of starting oxidation found for usual TiN coatings. These properties are related to the B1-structure observed for Ti1xAlxN coatings with x up to 0.7 for PVD and 0.9 for PACVD and CVD processes. In in-service conditions, an increase of hardness and oxidation resistance is also observed. The increased hardness is related to the so-called age-hardening effect observed at elevated temperature (≈800C) in the Ti-Al-N system and the enhancement of oxidation resistance is associated to the formation of a dense and protective Al2O3 top-layer. The adjustment of the aluminum content and the control of the cutting parameters lead to better wear resistance and tool lifetime in machining operations, for

The authors would like to acknowledge the CEA Saclay, Direction du Programme Matériaux Avancés and Groupement d'Intérêt Public Haute-Marne (GIP 52) for their financial support.

Summary of mechanical properties depending on Aluminum content and structure of

Ti1xAlxN Ti1xAlxN coatings obtained by different deposition processes.

which both high mechanical and chemical stabilities are required.

Acknowledgements

A. Appendix

reported in [99, 112–114], respectively.

5. Summary

Figure 14. Evolution of wear track depth vs. temperature for coatings with different crystallographic structures (singlephase cubic for x = 0.4, two-phase cubic/hexagonal for 0.58, and single-phase hexagonal for x = 0.7) [109].

Figure 15. Evolution of the tool lifetime vs. Al content for Ti1xAlxN coatings [26].

content in Ti1xAlxN during machining tests like milling or drilling. However, it appears that even small amounts of aluminum in TiN coatings lead to a strong increase of the tool lifetime [50, 64, 111]. Hörling et al. also showed that increasing aluminum content in the cubic Ti1xAlxN single-phased coatings led to a continuous increase of the tool lifetime until x ≈ 0.66 (Figure 15). This behavior should be related to the higher hardness as well as a better oxidation resistance [26, 110]. A further increase of the aluminum content leads to a drop of the tool lifetime due to the presence of hcp-AlN with poor mechanical properties.

More details about the influence of machining parameters (cutting speed, cutting tool material, and geometry) on the wear and breakage mechanisms of Ti1xAlxN-coated tools during machining of high-speed steel, Ni-based superalloys, aluminum and titanium alloys are reported in [99, 112–114], respectively.

### 5. Summary

In this chapter, the thermodynamical basis, deposition, and some mechanical and chemical properties of coatings based on Ti1xAlxN and Ti1xAlxSiyN have been reviewed. The general mechanisms (hardening effect, oxidation mechanisms, and wear mechanisms) associated to these properties were also discussed by pointing out the influence of the temperature and the aluminum content in the films. Increasing aluminum content induces structural changes that are directly associated with the properties. The high thermal stability and the age-hardening effect, observed in titanium-aluminum nitride-based coatings, are related to the coherent spinodal decomposition in the Ti-Al-N system. Concerning elaboration by means of CVD based processes, the stabilization of the fcc-Ti1xAlxN metastable solid solution can be achieved by a combination of relatively low deposition temperatures, high partial pressures of reactive precursors, and low total pressure. As-deposited monolayer coatings show a higher hardness and a better oxidation resistance when aluminum content is increased. Hardness up to about 30 GPa and stability in air up to 800C have been measured. These values outperform the typical values of about 25 GPa for hardness and 500C for the temperature of starting oxidation found for usual TiN coatings. These properties are related to the B1-structure observed for Ti1xAlxN coatings with x up to 0.7 for PVD and 0.9 for PACVD and CVD processes. In in-service conditions, an increase of hardness and oxidation resistance is also observed. The increased hardness is related to the so-called age-hardening effect observed at elevated temperature (≈800C) in the Ti-Al-N system and the enhancement of oxidation resistance is associated to the formation of a dense and protective Al2O3 top-layer. The adjustment of the aluminum content and the control of the cutting parameters lead to better wear resistance and tool lifetime in machining operations, for which both high mechanical and chemical stabilities are required.

### Acknowledgements

The authors would like to acknowledge the CEA Saclay, Direction du Programme Matériaux Avancés and Groupement d'Intérêt Public Haute-Marne (GIP 52) for their financial support.

### A. Appendix

content in Ti1xAlxN during machining tests like milling or drilling. However, it appears that even small amounts of aluminum in TiN coatings lead to a strong increase of the tool lifetime [50, 64, 111]. Hörling et al. also showed that increasing aluminum content in the cubic Ti1xAlxN single-phased coatings led to a continuous increase of the tool lifetime until x ≈ 0.66 (Figure 15). This behavior should be related to the higher hardness as well as a better oxidation resistance [26, 110]. A further increase of the aluminum content leads to a drop of the

Figure 14. Evolution of wear track depth vs. temperature for coatings with different crystallographic structures (single-

phase cubic for x = 0.4, two-phase cubic/hexagonal for 0.58, and single-phase hexagonal for x = 0.7) [109].

194 Coatings and Thin-Film Technologies

More details about the influence of machining parameters (cutting speed, cutting tool material, and geometry) on the wear and breakage mechanisms of Ti1xAlxN-coated tools during

tool lifetime due to the presence of hcp-AlN with poor mechanical properties.

Figure 15. Evolution of the tool lifetime vs. Al content for Ti1xAlxN coatings [26].

Summary of mechanical properties depending on Aluminum content and structure of Ti1xAlxN Ti1xAlxN coatings obtained by different deposition processes.


Ref. x Structure

0.85 hcp

[25] 0 fcc

 Cathodic arc ion

(Mo)

 400C

0.40

100

 /

/

≈20.6

 /

plating

0.6 fcc 0.7 fcc/hcp

0.85 hcp

[110] 0 fcc

0.5 fcc 0.6 fcc/hcp

[116] 0.5 fcc

[64] 0 fcc

0.64 fcc [28] 0.02 fcc

0.83 fcc

0.91 fcc/hcp

[27] 0.82 fcc

0.9 fcc

 LPCVD

 LPCVD

 PACVD

 H11 steel or

510C

150

 /

530

/

≈49

 /

WC/Co

WC/6%Co-TiN

800C 800C

<10.000

 /

/

/

32.4

 539 197

<10.000

 /

/

/

29.2

 444

 PACVD

 H11 steel or

510C

150

 /

480

/

38.7

 /

http://dx.doi.org/10.5772/intechopen.79747

WC/Co

 PACVD

 PACVD

 H11 steel or

WC/Co

 Z85 WCDV

 500C 510C

150

 /

430

/

≈23.5

 /

 PACVD

 Z85 WCDV

 500C

50–100

50–100

 /

300

3.2

≈25.7

≈365

 /

150

+0.3

≈29

≈490

 Cathodic arc

WC/Co

 450C

3.20

40

 /

2.48

27.1

 515

evaporation

 Reactive DC

SKH51

 350C

0.80

80

 /

/

≈21.6

 /

Ti-Al-N-Based Hard Coatings: Thermodynamical Background, CVD Deposition, and Properties. A Review

magnetron

 Reactive DC

SKH51

 350C

0.80

80

 /

/

≈37.2

 /

magnetron

 Reactive DC

SKH51

 350C

0.80

80

 /

/

≈23.5

 /

magnetron

 Cathodic arc ion

(Mo)

 400C

0.40

100

 /

/

≈19.6

 /

plating

 Cathodic arc ion

(Mo)

 400C

0.40

100

 /

/

≈27

 /

plating

 Cathodic arc ion

(Mo)

 400C

0.40

100

 /

/

≈31.4

 /

plating

 Cathodic arc ion

WC/Co

 450C

 2–5

50

 /

/

≈18.6

 /

plating

 Process

Substrate

 Deposition

Pressure

Bias

Discharge

Residual

Hardness

Young

modulus (Gpa)

(Gpa)

stresses (GPa)

temperature

 (C)

(Pa)

voltage

voltage (V)

(V)


Ref. x Structure

[23] 0 fcc

 Magnetron

Si

500C

0.40

60

 /

1356

22.8

 378

196 Coatings and Thin-Film Technologies

sputtering

0.62 fcc 0.67 fcc/hcp

0.75 hcp

[26] 0 fcc

0.66 fcc 0.74 fcc/hcp

[90] ≈0.5 fcc

[3] 0 fcc

 Close-field

M42 steel or

Room

temperature

 0.27

50

 /

≈ 3

23

 245.3

Si

M42 steel or

Room

temperature

 0.27

50

 /

≈ 0.35

 31.4

 315.2

Si

M42 steel or

Room

temperature

 0.27

50

 /

≈ 0.45

 28.5

 304.8

Si

M42 steel or

Room

temperature

 0.27

/

 /

/

/

≈37.5

 /

50

 /

≈ 1.5

13.8

 175.2

Si

WC/Co

magnetron

0.41 fcc 0.48 fcc/hcp

1 hcp

[19] 0.66 fcc

[115] 0.48 fcc

[98] 0 fcc

 Cathodic arc ion

WC/Co

 450C

 2–5

50

 /

/

≈20

 /

plating

0.6 fcc

 Cathodic arc ion

WC/Co

 450C

 2–5

50

 /

/

≈30

 /

plating

 Cathodic arc

WC/Co

 550C

2.00

100

 /

/

31.2

 /

evaporation

 Cathodic arc

evaporation

 Close-field magnetron

 Close-field magnetron

 Close-field magnetron

 Arc evaporation

 Arc evaporation

 WC/Co-6%

 WC/Co-6%

 500C

 Arc evaporation

 WC/Co-6%

 500C

 //

/

/

40 to

/

1.7 to

5

200

 /

/

 Arc evaporation

 WC/Co-6%

 500C

 //

 /

 /

2.9 3.1

32.6

32.3

 595

 627

28.5

 617

 Magnetron

Si

500C

0.40

60

 /

+0.186

22.9

≈255

sputtering

 Magnetron

Si

500C

0.40

60

 /

+0.102

27.2

≈320

sputtering

 Magnetron

Si

500C

0.40

60

 /

0.883

31.3

≈380

sputtering

 Process

Substrate

 Deposition

Pressure

Bias

Discharge

Residual

Hardness

Young

modulus (Gpa)

(Gpa)

stresses (GPa)

temperature

 (C)

(Pa)

voltage

voltage (V)

(V)


Author details

Florent Uny1,2\*, Elisabeth Blanquet

4 Centre CEA de Saclay, France

References

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2012;60(5):2091-2096

2017;127:182-185

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2 Nogent International Center for CVD Innovation, LRC CEA-ICD-LASMIS, UTT, Antenne de

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'un réacteur

<sup>2</sup> absorber for solar selec-

### Author details

Florent Uny1,2\*, Elisabeth Blanquet3 , Frédéric Schuster<sup>4</sup> and Frédéric Sanchette1,2

\*Address all correspondence to: florent.uny@utt.fr

1 ICD-LASMIS, Université de Technologie de Troyes, CNRS, Antenne de Nogent, Pôle Technologique de Haute-Champagne, Nogent, France

2 Nogent International Center for CVD Innovation, LRC CEA-ICD-LASMIS, UTT, Antenne de Nogent, Pôle Technologique de Haute-Champagne, Nogent, France


### References

Ref. x Structure

 Process

Substrate

 Deposition

Pressure

Bias

Discharge

Residual

Hardness

Young

modulus (Gpa)

(Gpa)

stresses (GPa)

temperature

WC/6%Co-TiN

> 0.82 fcc

0.82 fcc

 LPCVD

WC/6%Co-TiN

800C

<10.000

 /

/

0.92

/

/

 LPCVD

WC/6%Co-TiN

850C

<10.000

 /

/

0.45

/

/

198 Coatings and Thin-Film Technologies

 (C)

(Pa)

voltage

voltage (V)

(V)


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**Chapter 10**

**Provisional chapter**

**New Materials for Thin Film Solar Cells**

**New Materials for Thin Film Solar Cells**

DOI: 10.5772/intechopen.81393

Thin film technology has a world-wide reputation in the field of thin film deposition process and also it paves a way for innovative techniques in large scale applications. Modern thin film technology has evolved into a sophisticated way to increase the performance and esthetic value for making new functional devices. One such application is search of new materials for thin film solar cells as it provides the solution for the today's concern of energy crisis. Depending on the processing technology solar cells are of various types. Among them, silicon wafer solar cells and thin film solar cells are most promising. Thin film technology has made solar cells more feasible to be employed in terms of device design and fabrication. The efficiencies produced by these solar cells still need to be improved. For this many investigations for further improvement from CIGS (copper indium gallium selenide) solar cell to dye sensitized solar cells and perovskite solar cells. Due to toxic nature and environmental impact the use of lead in perovskite solar cells are replaced by tin or some materials which would equalize the achieved efficiency of lead. Hence the developments in search of innovative materials continue its path in thin film

solar cells to develop the photovoltaic field by enhancing its efficiency.

**Keywords:** solar cell, physical vapor deposition, drop casting, interfacial impact,

Energy is the key factor for any living creature to exist in the universe. With the advent of industrialization and increase in population led to a surge in the crisis for energy. The reduction of our dependence on fossil fuels (oil, coal and natural gas), as well as the evolution towards a cleaner future requires the large deployment of sustainable renewable energy sources. Among them solar energy is the most abundant and is available throughout the year.

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

© 2019 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use,

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

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

Senthil T.S and Kalaiselvi C.R

Senthil T.S and Kalaiselvi C.R

**Abstract**

efficiency

**1. Introduction**

http://dx.doi.org/10.5772/intechopen.81393


#### **New Materials for Thin Film Solar Cells New Materials for Thin Film Solar Cells**

Senthil T.S and Kalaiselvi C.R Senthil T.S and Kalaiselvi C.R

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

http://dx.doi.org/10.5772/intechopen.81393

#### **Abstract**

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speed steel. Surface and Coating Technology. 1998;108:369-376

face and Coating Technology. 2004;177-178:623-626

and Manufacture. 2005;45(12-13):1436-1442

and Hard Materials. 2017;62:47-57

593-596

surface design. Tribology International. 1998;31(1):107-120

Mines de Paris; 2009

206 Coatings and Thin-Film Technologies

2015;41(9):10349-10379

and Films. 1992;10(4):1749

Thin film technology has a world-wide reputation in the field of thin film deposition process and also it paves a way for innovative techniques in large scale applications. Modern thin film technology has evolved into a sophisticated way to increase the performance and esthetic value for making new functional devices. One such application is search of new materials for thin film solar cells as it provides the solution for the today's concern of energy crisis. Depending on the processing technology solar cells are of various types. Among them, silicon wafer solar cells and thin film solar cells are most promising. Thin film technology has made solar cells more feasible to be employed in terms of device design and fabrication. The efficiencies produced by these solar cells still need to be improved. For this many investigations for further improvement from CIGS (copper indium gallium selenide) solar cell to dye sensitized solar cells and perovskite solar cells. Due to toxic nature and environmental impact the use of lead in perovskite solar cells are replaced by tin or some materials which would equalize the achieved efficiency of lead. Hence the developments in search of innovative materials continue its path in thin film solar cells to develop the photovoltaic field by enhancing its efficiency.

DOI: 10.5772/intechopen.81393

**Keywords:** solar cell, physical vapor deposition, drop casting, interfacial impact, efficiency

### **1. Introduction**

Energy is the key factor for any living creature to exist in the universe. With the advent of industrialization and increase in population led to a surge in the crisis for energy. The reduction of our dependence on fossil fuels (oil, coal and natural gas), as well as the evolution towards a cleaner future requires the large deployment of sustainable renewable energy sources. Among them solar energy is the most abundant and is available throughout the year.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2019 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Moreover the solar energy has the greatest potential to fulfill the thirst for energy and the need for innovation of clean and eco-friendly technologies. In this perspective developing solar cells is one of the best approaches to convert solar energy into electrical energy based on photovoltaic effect. Depending on the cell material and the processing techniques solar cells are of two kinds based on wafer and thin film. Thin film solar cells have the key advantage of their dimensionality, having thickness a fraction of other types of solar cells and it is attractive in terms of cost factor with minimum material usage [1]. However thin film solar cells promise to achieve the goal with the low cost and high efficiency from CIGS to dye sensitized solar cells and recently with the advent of perovskite based materials for solar cells.

**2.** Introducing Se vapor at any time during the process to avoid potential damage to the bot-

New Materials for Thin Film Solar Cells http://dx.doi.org/10.5772/intechopen.81393 209

Deposition of CIGS junction on a soda lime glass substrate covered by Mo layer at 1 μm thick [1]. The best back contact material for CIGS solar cells was found to be molybdenum since it is having high conductivity and relative stability at the high processing temperature. The deposition of Mo layer is by means of DC magnetron sputtering technique. The grown films play a crucial role in the performance of the CIGS device. During the deposition process,

The two points to be considered in the CIGS thin films are: (1) the Mo back contact which cannot obviously be used in a tandem cell due to its opacity, (2) the CdS layer deposited by the chemical bath deposition (CBD) technique, which can impact bottom cell properties. The schematic view of the CIGS thin film by sputtering process is shown in **Figure 1**. The bottom Mo contact uses rather a common bilayer structure, a first adhesion layer is magnetron DC sputtered at high pressure and a second one is deposited at low pressure to increase the conductivity. A standard CdS buffer layer is deposited by chemical bath deposition. Ultimately, it will be replaced by a Zn layer deposited using the sputtering technique at room temperature. The top electrode is composed of an ZnO/aluminum doped zinc oxide (AZO) bi-layer. The absorber layer process is deposited at room temperature using pulsed DC magnetron sputtering and argon plasma [1]. In order to obtain the chemical composition of CIGS absorber, energy dispersive spectroscopy measurements (EDS) were carried out under 10 kV mode and it was shown in **Figure 2**. Atomic composition of the obtained film was calculated as: Cu—14.8 at%, In—18.35 at%, Ga—2.86 at% and Se—25.95 at%. On the basis of measured values the relative ratios of Cu/(In + Ga) and Ga/(In + Ga) were found to be 0.7 and 0.13, respectively [5].

**3.** Cd free buffer layer also targeting an environmentally friendly process [1].

parameters such as power, time and pressure should be carefully chosen [5].

tom Si cell by invoking unwanted species.

**Figure 1.** Preparation of CIGS thin films by sputtering process.

### **2. Preparation of thin film solar cells**

To prepare solar cell there are variety of methods and materials are used, among them thin film solar cells are unique. In this chapter the historical background and the emergence of new techniques in the growth of thin film solar cells such as CIGS, dye sensitized solar cells and perovskite solar cells are presented.

### **2.1. CIGS thin film solar cells**

Copper indium gallium selenide (CIGS) based solar cells are receiving worldwide attraction for solar power generation. These materials absorb light at a rate of 10–100 times more efficient compared with silicon-based solar cells, thus the thickness of the films obtained in the order of a few microns. The major advantage of this technology is attributed as of low raw material usage with less complex procedure for manufacturing. CIGS solar cells exhibit high radiation resistance, making them suitable for space applications [2].

CIGS is a promising absorber material and received considerable attention due to its direct band gap, high absorption co-efficient and less material wastage. It is an efficient thin film solar cell with the efficiency of 22.8% comparable to crystalline silicon (c-Si) wafer based solar cells [3]. CIGS thin films can be fabricated by various methods among them physical vapor deposition is important.

Thin film deposition by PVD is regarded as a vacuum coating method and is classified into two techniques, evaporation and sputtering. The particles are capable of moving in a straight path as the system is kept in vacuum and the films are coated by physical means that are commonly directional, rather than conformal in nature [4].

### *2.1.1. Sputtering method*

Sputtering deposition is a vacuum based process that is used for CIGS cell fabrication. The fabrication process will involve the following steps:

**1.** Deposition at room temperature for attaining low temperature process and hence reduced the cost of equipment.


Moreover the solar energy has the greatest potential to fulfill the thirst for energy and the need for innovation of clean and eco-friendly technologies. In this perspective developing solar cells is one of the best approaches to convert solar energy into electrical energy based on photovoltaic effect. Depending on the cell material and the processing techniques solar cells are of two kinds based on wafer and thin film. Thin film solar cells have the key advantage of their dimensionality, having thickness a fraction of other types of solar cells and it is attractive in terms of cost factor with minimum material usage [1]. However thin film solar cells promise to achieve the goal with the low cost and high efficiency from CIGS to dye sensitized

solar cells and recently with the advent of perovskite based materials for solar cells.

To prepare solar cell there are variety of methods and materials are used, among them thin film solar cells are unique. In this chapter the historical background and the emergence of new techniques in the growth of thin film solar cells such as CIGS, dye sensitized solar cells and

Copper indium gallium selenide (CIGS) based solar cells are receiving worldwide attraction for solar power generation. These materials absorb light at a rate of 10–100 times more efficient compared with silicon-based solar cells, thus the thickness of the films obtained in the order of a few microns. The major advantage of this technology is attributed as of low raw material usage with less complex procedure for manufacturing. CIGS solar cells exhibit high

CIGS is a promising absorber material and received considerable attention due to its direct band gap, high absorption co-efficient and less material wastage. It is an efficient thin film solar cell with the efficiency of 22.8% comparable to crystalline silicon (c-Si) wafer based solar cells [3]. CIGS thin films can be fabricated by various methods among them physical vapor

Thin film deposition by PVD is regarded as a vacuum coating method and is classified into two techniques, evaporation and sputtering. The particles are capable of moving in a straight path as the system is kept in vacuum and the films are coated by physical means that are

Sputtering deposition is a vacuum based process that is used for CIGS cell fabrication. The

**1.** Deposition at room temperature for attaining low temperature process and hence reduced

radiation resistance, making them suitable for space applications [2].

commonly directional, rather than conformal in nature [4].

fabrication process will involve the following steps:

**2. Preparation of thin film solar cells**

perovskite solar cells are presented.

**2.1. CIGS thin film solar cells**

208 Coatings and Thin-Film Technologies

deposition is important.

*2.1.1. Sputtering method*

the cost of equipment.

Deposition of CIGS junction on a soda lime glass substrate covered by Mo layer at 1 μm thick [1]. The best back contact material for CIGS solar cells was found to be molybdenum since it is having high conductivity and relative stability at the high processing temperature. The deposition of Mo layer is by means of DC magnetron sputtering technique. The grown films play a crucial role in the performance of the CIGS device. During the deposition process, parameters such as power, time and pressure should be carefully chosen [5].

The two points to be considered in the CIGS thin films are: (1) the Mo back contact which cannot obviously be used in a tandem cell due to its opacity, (2) the CdS layer deposited by the chemical bath deposition (CBD) technique, which can impact bottom cell properties. The schematic view of the CIGS thin film by sputtering process is shown in **Figure 1**. The bottom Mo contact uses rather a common bilayer structure, a first adhesion layer is magnetron DC sputtered at high pressure and a second one is deposited at low pressure to increase the conductivity. A standard CdS buffer layer is deposited by chemical bath deposition. Ultimately, it will be replaced by a Zn layer deposited using the sputtering technique at room temperature. The top electrode is composed of an ZnO/aluminum doped zinc oxide (AZO) bi-layer. The absorber layer process is deposited at room temperature using pulsed DC magnetron sputtering and argon plasma [1].

In order to obtain the chemical composition of CIGS absorber, energy dispersive spectroscopy measurements (EDS) were carried out under 10 kV mode and it was shown in **Figure 2**. Atomic composition of the obtained film was calculated as: Cu—14.8 at%, In—18.35 at%, Ga—2.86 at% and Se—25.95 at%. On the basis of measured values the relative ratios of Cu/(In + Ga) and Ga/(In + Ga) were found to be 0.7 and 0.13, respectively [5].

**Figure 1.** Preparation of CIGS thin films by sputtering process.

**2.2. Dye sensitized solar cell**

absorbed on a thin TiO<sup>2</sup>

*2.2.1. Doctor blade technique*

method the resulting TiO<sup>2</sup>

**Figure 4.** Doctor blade technique.

compressed for 30 s. The derived TiO<sup>2</sup>

method, TiO<sup>2</sup>

free TiO<sup>2</sup>

ner: TiO<sup>2</sup>

simple and low cost techniques [10].

The existence of an optimized thickness of TiO<sup>2</sup>

For the modified doctor blade method, a dense TiO<sup>2</sup>

Dye-sensitized solar cell (DSSC) has been known as a promising photovoltaic device to achieve moderate efficiency at low cost [7]. The principle of DSSC imitates natures novel effect of photosynthesis. In DSSCs, the photo-sensitizer captured the incident photons that are

transfer ions to a counter electrode thus electric current is produced [8]. DSSCs have attained efficiencies of upto 12.3% [9] and attained benefit from their versatile nature and low cost for

ally prepared by various methods. Among them doctor blade and spin coating methods are

Doctor blade or tape casting is one of the widely used techniques for producing thin films on large area surfaces. In the doctor blading process, a well-mixed slurry consisting of a suspension of ceramic particles along with other additives is placed on a substrate and a constant relative movement is established between the blade and the substrate resulting in the formation of a gel-layer when dried [11]. The diagrammatic view of this technique is shown in (**Figure 4**).

and transfer to the film so that in DSSC the highest efficiency could be achieved [12]. In this

thin film. The improved performance of DSSC could be ascribed to the compact and crack-

a plastic bag for 10 min under rolling of a steel pipe. After that, by the way of doctor blade

air for 30 min. Then the film surface was covered by a drop of lubricating oil and stabilized at 100°C in an oven for reducing cracks. Finally, the film was covered with a flat glass and

novel method is also observed to substantially improve the overall conversion efficiency of the resulting DSSC which presents strong potential for DSSC module construction [10].

thin film prepared by the modified doctor blade method [10].

paste usually employed as a surfactant material to increase the porosity of TiO<sup>2</sup>

powder was blended in a mixture of polyethylene glycol and de-ionized water in

paste was coated on TCO glass [13]. Then the film was dried in

thin film was annealed at 450°C for 2 h [10]. This

the manufacturing process. The main component of DSSC is TiO<sup>2</sup>

layer placed on the anode. DSSCs make use of liquid electrolytes to

thin film that is convention-

New Materials for Thin Film Solar Cells http://dx.doi.org/10.5772/intechopen.81393 211

thin film is inherent for the charge storage

paste was prepared in a following man-

**Figure 2.** EDS spectral analysis of the top surface of CIGS absorber layer.

A sputtering technique using low temperature processes and without any hazardous gas diffusion at high temperature can then lead to functional devices. Hence sputtering technique can be successfully used for the fabrication of CIGS thin films solar cells [5]. SEM cross sectional view of CIGS thin film solar cell is shown in **Figure 3**.

The stoichiometry of CIGS is very complex thereby making uniform and large-area deposition is very difficult. It also indeed of encapsulation which is expensive as it prone to moisture and oxygen easily. Also, their reliance on rare elements of tellurium & indium and recycling of toxic element like cadmium may limit their potential for large-scale production and it would be replaced by some other elements [6].

**Figure 3.** SEM image of CIGS thin film solar cell.

### **2.2. Dye sensitized solar cell**

Dye-sensitized solar cell (DSSC) has been known as a promising photovoltaic device to achieve moderate efficiency at low cost [7]. The principle of DSSC imitates natures novel effect of photosynthesis. In DSSCs, the photo-sensitizer captured the incident photons that are absorbed on a thin TiO<sup>2</sup> layer placed on the anode. DSSCs make use of liquid electrolytes to transfer ions to a counter electrode thus electric current is produced [8]. DSSCs have attained efficiencies of upto 12.3% [9] and attained benefit from their versatile nature and low cost for the manufacturing process. The main component of DSSC is TiO<sup>2</sup> thin film that is conventionally prepared by various methods. Among them doctor blade and spin coating methods are simple and low cost techniques [10].

### *2.2.1. Doctor blade technique*

A sputtering technique using low temperature processes and without any hazardous gas diffusion at high temperature can then lead to functional devices. Hence sputtering technique can be successfully used for the fabrication of CIGS thin films solar cells [5]. SEM cross sec-

The stoichiometry of CIGS is very complex thereby making uniform and large-area deposition is very difficult. It also indeed of encapsulation which is expensive as it prone to moisture and oxygen easily. Also, their reliance on rare elements of tellurium & indium and recycling of toxic element like cadmium may limit their potential for large-scale production and it would

tional view of CIGS thin film solar cell is shown in **Figure 3**.

**Figure 2.** EDS spectral analysis of the top surface of CIGS absorber layer.

be replaced by some other elements [6].

210 Coatings and Thin-Film Technologies

**Figure 3.** SEM image of CIGS thin film solar cell.

Doctor blade or tape casting is one of the widely used techniques for producing thin films on large area surfaces. In the doctor blading process, a well-mixed slurry consisting of a suspension of ceramic particles along with other additives is placed on a substrate and a constant relative movement is established between the blade and the substrate resulting in the formation of a gel-layer when dried [11]. The diagrammatic view of this technique is shown in (**Figure 4**).

The existence of an optimized thickness of TiO<sup>2</sup> thin film is inherent for the charge storage and transfer to the film so that in DSSC the highest efficiency could be achieved [12]. In this method, TiO<sup>2</sup> paste usually employed as a surfactant material to increase the porosity of TiO<sup>2</sup> thin film. The improved performance of DSSC could be ascribed to the compact and crackfree TiO<sup>2</sup> thin film prepared by the modified doctor blade method [10].

For the modified doctor blade method, a dense TiO<sup>2</sup> paste was prepared in a following manner: TiO<sup>2</sup> powder was blended in a mixture of polyethylene glycol and de-ionized water in a plastic bag for 10 min under rolling of a steel pipe. After that, by the way of doctor blade method the resulting TiO<sup>2</sup> paste was coated on TCO glass [13]. Then the film was dried in air for 30 min. Then the film surface was covered by a drop of lubricating oil and stabilized at 100°C in an oven for reducing cracks. Finally, the film was covered with a flat glass and compressed for 30 s. The derived TiO<sup>2</sup> thin film was annealed at 450°C for 2 h [10]. This novel method is also observed to substantially improve the overall conversion efficiency of the resulting DSSC which presents strong potential for DSSC module construction [10].

**Figure 4.** Doctor blade technique.

**Figure 5.** FE-SEM images of surface of TiO<sup>2</sup> thin film prepared by doctor blade method. Ref. [14].

Surface morphology of two kinds of TiO<sup>2</sup> thin films are shown in **Figure 5**. Even though they are looking similar, the variation in the thickness is larger for modified doctor blade thin film than the conventional one.

**2.3. Perovskite solar cells (PSCs)**

**Figure 6.** Flow diagram of spin coating process.

inorganic halide perovskites, CH<sup>3</sup>

*2.3.2. Interfacial impact on the performance*

and HC(NH<sup>2</sup>

excitons whereas organic photovoltaic and DSSCs are excitonic.

)+

NH<sup>3</sup>

PbBr<sup>3</sup>

perovskite material acts as a light absorber and electron transporting layer (TiO<sup>2</sup>

Perovskite solar cells are in focus of the solar cell development research in recent years due to their high efficiency, cost effective fabrication and band gap tuning ability. Perovskite compound was first discovered by Gustav Rose in 1839, named after a Russian mineralogist L.A. Perovski. The specific crystal structure was first found in an inorganic mineral CaTiO<sup>3</sup>

cation B with divalent metal ions such as Pb2+, Sn2+ or Cu2+ and the X anions are halides (Cl−, Br−, I−) [19]. Perovskite are non-excitonic as no external force is needed for the generation of

Miyasaka and his co-workers replaced the dye pigment in DSSCs with two hybrid organic-

3.81% obtained was not a considerable one [20]. A major breakthrough happened in 2012 when Gratzel and Park et al. used Spiro-MeOTAD as the hole transport material (HTM) with efficiency of 9.7%. In 2016 the efficiency of perovskite solar cells was improved to be 22.1% [21].

The perovskite solar cells employs a perovskite absorber between an electron transporting layer on a conducting glass substrate (FTO) and a hole transporting layer with a metal back contact on top. The working principle of this device was similar to that of DSSCs. Here

charge separation and electron transport whereas the holes are transferred to hole transporting layer. An ideal selective contact does not deteriorate the light absorbing layer and also does not induce degradation of the device. There are also no energy losses when photo-generated

and CH<sup>3</sup>

NH<sup>3</sup> PbI<sup>3</sup>

) with a cubic unit cell [18]. The cation A is replaced by a small organic cations such as

to create organic-inorganic hybrid materials while the

and the efficiency of 3.13 and

New Materials for Thin Film Solar Cells http://dx.doi.org/10.5772/intechopen.81393 213

) takes part in

*2.3.1. Emergence of perovskite*

(ABX<sup>3</sup>

CH<sup>3</sup> NH<sup>3</sup> + , C<sup>2</sup> H5 NH<sup>3</sup> +

### *2.2.2. Spin coating method*

Spin coating method is one of the most common techniques for the preparation of thin films on substrates. Spin coating method is widely used in micro-fabrication, where it can be used to create thin films with thicknesses below 10 nm [15]. The advantage of spin coating method is its ability to quickly and easily produce uniform films, ranging from a few nanometres to few microns in thickness.

The spin coating procedure includes deposition, spin up, spin off, and evaporation. The substrate being covered by depositing the solution which is rotated at high speeds to coat the substrate using centrifugal force. The volatile solvent easily evaporates and hence the desired film's thickness is dependent on the concentration of solution, solvent, and spin speeds [16]. The film thickness of the order less than 10 nm is useful in the field of micro fabrication and photolithography [4]. The flow diagram for the spin coating process is shown in the (**Figure 6**).

Titanium (IV) isopropoxide (TTIP), ethyl alcohol, nitric acid (HNO<sup>3</sup> ) and distilled water were used as received without further purification. The synthesis procedure of nanocrystalline TiO<sup>2</sup> thin film can be obtained by mixing titanium isopropoxide with ethanol and the distilled water was added drop by drop by continuous stirring about 1 h. The resulting solution was peptized using nitric acid and refluxed at 80°C. After that TiO<sup>2</sup> sol had been prepared and coated on conductive glass substrate with the spin rate of 3000 rpm for 30 s. Then followed by annealing the TiO<sup>2</sup> thin films at 450°C [17].

The major challenge of DSSC is that it should withstand its stability under high temperature and its low absorption coefficients due to its interfacial recombination. Moreover the cost of the ruthenium dye used sets its intrinsic drawback limits its scope to small scale applications. Then search for new innovative material pertains to continue its thirst for perovskite related materials by the replacement of liquid electrolyte.

**Figure 6.** Flow diagram of spin coating process.

### **2.3. Perovskite solar cells (PSCs)**

#### *2.3.1. Emergence of perovskite*

Surface morphology of two kinds of TiO<sup>2</sup>

**Figure 5.** FE-SEM images of surface of TiO<sup>2</sup>

than the conventional one.

212 Coatings and Thin-Film Technologies

*2.2.2. Spin coating method*

few microns in thickness.

TiO<sup>2</sup>

annealing the TiO<sup>2</sup>

thin films are shown in **Figure 5**. Even though they

) and distilled water were

sol had been prepared and

are looking similar, the variation in the thickness is larger for modified doctor blade thin film

thin film prepared by doctor blade method. Ref. [14].

Spin coating method is one of the most common techniques for the preparation of thin films on substrates. Spin coating method is widely used in micro-fabrication, where it can be used to create thin films with thicknesses below 10 nm [15]. The advantage of spin coating method is its ability to quickly and easily produce uniform films, ranging from a few nanometres to

The spin coating procedure includes deposition, spin up, spin off, and evaporation. The substrate being covered by depositing the solution which is rotated at high speeds to coat the substrate using centrifugal force. The volatile solvent easily evaporates and hence the desired film's thickness is dependent on the concentration of solution, solvent, and spin speeds [16]. The film thickness of the order less than 10 nm is useful in the field of micro fabrication and photolithography [4]. The flow diagram for the spin coating process is shown in the (**Figure 6**).

used as received without further purification. The synthesis procedure of nanocrystalline

coated on conductive glass substrate with the spin rate of 3000 rpm for 30 s. Then followed by

The major challenge of DSSC is that it should withstand its stability under high temperature and its low absorption coefficients due to its interfacial recombination. Moreover the cost of the ruthenium dye used sets its intrinsic drawback limits its scope to small scale applications. Then search for new innovative material pertains to continue its thirst for perovskite related

 thin film can be obtained by mixing titanium isopropoxide with ethanol and the distilled water was added drop by drop by continuous stirring about 1 h. The resulting solution was

Titanium (IV) isopropoxide (TTIP), ethyl alcohol, nitric acid (HNO<sup>3</sup>

peptized using nitric acid and refluxed at 80°C. After that TiO<sup>2</sup>

thin films at 450°C [17].

materials by the replacement of liquid electrolyte.

Perovskite solar cells are in focus of the solar cell development research in recent years due to their high efficiency, cost effective fabrication and band gap tuning ability. Perovskite compound was first discovered by Gustav Rose in 1839, named after a Russian mineralogist L.A. Perovski. The specific crystal structure was first found in an inorganic mineral CaTiO<sup>3</sup> (ABX<sup>3</sup> ) with a cubic unit cell [18]. The cation A is replaced by a small organic cations such as CH<sup>3</sup> NH<sup>3</sup> + , C<sup>2</sup> H5 NH<sup>3</sup> + and HC(NH<sup>2</sup> )+ to create organic-inorganic hybrid materials while the cation B with divalent metal ions such as Pb2+, Sn2+ or Cu2+ and the X anions are halides (Cl−, Br−, I−) [19]. Perovskite are non-excitonic as no external force is needed for the generation of excitons whereas organic photovoltaic and DSSCs are excitonic.

Miyasaka and his co-workers replaced the dye pigment in DSSCs with two hybrid organicinorganic halide perovskites, CH<sup>3</sup> NH<sup>3</sup> PbBr<sup>3</sup> and CH<sup>3</sup> NH<sup>3</sup> PbI<sup>3</sup> and the efficiency of 3.13 and 3.81% obtained was not a considerable one [20]. A major breakthrough happened in 2012 when Gratzel and Park et al. used Spiro-MeOTAD as the hole transport material (HTM) with efficiency of 9.7%. In 2016 the efficiency of perovskite solar cells was improved to be 22.1% [21].

### *2.3.2. Interfacial impact on the performance*

The perovskite solar cells employs a perovskite absorber between an electron transporting layer on a conducting glass substrate (FTO) and a hole transporting layer with a metal back contact on top. The working principle of this device was similar to that of DSSCs. Here perovskite material acts as a light absorber and electron transporting layer (TiO<sup>2</sup> ) takes part in charge separation and electron transport whereas the holes are transferred to hole transporting layer. An ideal selective contact does not deteriorate the light absorbing layer and also does not induce degradation of the device. There are also no energy losses when photo-generated carriers are injected from the light absorbing material into the selective contact, hence no recombination at the interface, and the Fermi level of its corresponding carrier is maintained at the interface without any drop. The contact layer must also be balanced with respect to perovskite layer as otherwise it would lead to charge accumulation at selective contact and leads to interfacial charge recombination [22].

The primary role of a selective contact is to reduce the interfacial recombination between perovskite light absorbing layer and the FTO and Au extracting contacts. It should be as thin as possible in order to reduce the transport resistance but thick enough to avoid pinholes, hindering effective charge recombination. Obviously the goodness of an interfacing material for an efficient electron transport and performance in PSC will lies on the nature and interaction of the chosen selecting contact with the light absorbing perovskite layer. Tan et al. reported that using chlorine-capped TiO<sup>2</sup> colloidal nanocrystal film that mitigates interfacial recombination and improves interface binding in low-temperature planar solar cells [23].

Hole transporting material (HTM) also plays a vital role in both extracting and transporting the holes from the perovskite materials to the back electrodes thus minimizing undesired recombination losses at the interfaces. Spiro-OMeTAD is the widely used HTM results in better performance of the device. As of high cost and tedious synthesis inhibits its commercial application looking forward for inorganic semiconductors. Recently CuSCN has been used as HTM in perovskite solar cells and considerable high power conversion efficiency has been achieved [24].

There are numerous techniques employed to deposit different layers of perovskite solar cells like drop casting, spin coating, slot die coating, screen printing, ink-jet printing, etc. Among them drop casting is a simple and low cost method for the deposition of perovskite films [25].

### *2.3.3. Drop casting method*

Drop casting is a simple and potentially scalable casting method proposed for the fabrication of micro and nanocrystalline thin films. The impingement of a solution drop onto a substrate in a simple process called drop casting, usually results in spreading of the liquid solution and the formation of a non-uniform thin solid film after solvent evaporation [26]. Drop-casting usually happens by the way of releasing large droplets in a controlled manner that spreads and wet the surface upon impact as far better than spray coating, although its application is limited to small-area films and coatings [27]. Schematic view of perovskite solar cell with different layers is shown in the (**Figure 7**).

Drop casting is similar to spin coating method but the major difference is no substrate spinning is required and also less material wastage. Film thickness depends on the volume of dispersion used and the particle concentration, both of which can be easily varied. Generally, it is desirable to use solvents that are volatile, wet the substrate, and are not susceptible to thin film instabilities. However, water is not a recommended solvent for depositing some materials because the nanoparticles oxidize as the water evaporates from the sample [29]. In some cases alcohols can replace water, a good choice relies on the use of organic solvents such as hexane or toluene for nanoparticles with hydrophobic capping

New Materials for Thin Film Solar Cells http://dx.doi.org/10.5772/intechopen.81393 215

**Figure 7.** Schematic diagram of perovskite solar cell with different layers.

**Figure 8.** SEM image of perovskite solar cell with different layers.

The demerits of drop-casting are that even under ideal conditions, the rate of evaporation differs across the substrate thus leads to variations in film thickness or internal structure. However, it serves as a quick and accessible method to generate thin films even on a small substrates. **Figure 8** shows the SEM image of the perovskite thin films prepared with different layers.

ligands.

The formation of homogeneous thin perovskite layer is extremely important, and it can be developed by solution process. The process is employed for the most common structure of methyl ammonium lead iodide. The initial components taken in the appropriate ratio is spread over the entire surface of the substrate. Then, the spin-coater is accelerated to the desired rotational speed to evaporate the solvent. If the solvent does not dissolve the perovskite materials and is miscible with DMSO (dimethyl sulfoxide), a toluene or chloroform solution is dripped on the substrate during spinning. At last, all constituents are frozen into a uniform layer on the removal of the residual DMSO and a new complex as an intermediate phase is obtained. Films of different thicknesses were achieved by varying solution concentration [28].

**Figure 7.** Schematic diagram of perovskite solar cell with different layers.

carriers are injected from the light absorbing material into the selective contact, hence no recombination at the interface, and the Fermi level of its corresponding carrier is maintained at the interface without any drop. The contact layer must also be balanced with respect to perovskite layer as otherwise it would lead to charge accumulation at selective contact and

The primary role of a selective contact is to reduce the interfacial recombination between perovskite light absorbing layer and the FTO and Au extracting contacts. It should be as thin as possible in order to reduce the transport resistance but thick enough to avoid pinholes, hindering effective charge recombination. Obviously the goodness of an interfacing material for an efficient electron transport and performance in PSC will lies on the nature and interaction of the chosen selecting contact with the light absorbing perovskite layer. Tan et al.

recombination and improves interface binding in low-temperature planar solar cells [23].

Hole transporting material (HTM) also plays a vital role in both extracting and transporting the holes from the perovskite materials to the back electrodes thus minimizing undesired recombination losses at the interfaces. Spiro-OMeTAD is the widely used HTM results in better performance of the device. As of high cost and tedious synthesis inhibits its commercial application looking forward for inorganic semiconductors. Recently CuSCN has been used as HTM in perovskite solar cells and considerable high power conversion efficiency has been achieved [24].

There are numerous techniques employed to deposit different layers of perovskite solar cells like drop casting, spin coating, slot die coating, screen printing, ink-jet printing, etc. Among them drop casting is a simple and low cost method for the deposition of perovskite films [25].

Drop casting is a simple and potentially scalable casting method proposed for the fabrication of micro and nanocrystalline thin films. The impingement of a solution drop onto a substrate in a simple process called drop casting, usually results in spreading of the liquid solution and the formation of a non-uniform thin solid film after solvent evaporation [26]. Drop-casting usually happens by the way of releasing large droplets in a controlled manner that spreads and wet the surface upon impact as far better than spray coating, although its application is limited to small-area films and coatings [27]. Schematic view of perovskite solar cell with

The formation of homogeneous thin perovskite layer is extremely important, and it can be developed by solution process. The process is employed for the most common structure of methyl ammonium lead iodide. The initial components taken in the appropriate ratio is spread over the entire surface of the substrate. Then, the spin-coater is accelerated to the desired rotational speed to evaporate the solvent. If the solvent does not dissolve the perovskite materials and is miscible with DMSO (dimethyl sulfoxide), a toluene or chloroform solution is dripped on the substrate during spinning. At last, all constituents are frozen into a uniform layer on the removal of the residual DMSO and a new complex as an intermediate phase is obtained.

Films of different thicknesses were achieved by varying solution concentration [28].

colloidal nanocrystal film that mitigates interfacial

leads to interfacial charge recombination [22].

214 Coatings and Thin-Film Technologies

reported that using chlorine-capped TiO<sup>2</sup>

different layers is shown in the (**Figure 7**).

*2.3.3. Drop casting method*


**Figure 8.** SEM image of perovskite solar cell with different layers.

Drop casting is similar to spin coating method but the major difference is no substrate spinning is required and also less material wastage. Film thickness depends on the volume of dispersion used and the particle concentration, both of which can be easily varied. Generally, it is desirable to use solvents that are volatile, wet the substrate, and are not susceptible to thin film instabilities. However, water is not a recommended solvent for depositing some materials because the nanoparticles oxidize as the water evaporates from the sample [29]. In some cases alcohols can replace water, a good choice relies on the use of organic solvents such as hexane or toluene for nanoparticles with hydrophobic capping ligands.

The demerits of drop-casting are that even under ideal conditions, the rate of evaporation differs across the substrate thus leads to variations in film thickness or internal structure. However, it serves as a quick and accessible method to generate thin films even on a small substrates. **Figure 8** shows the SEM image of the perovskite thin films prepared with different layers.

### **3. Recent trends in solar cells**

### **3.1. Flexible perovskite solar cells (F-PSCs)**

Now research has been moved towards a flexible PSCs from rigid PSCs and it achieved notable milestones in terms of its efficiency with light weight [30]. Unlike rigid PSCs employ glass substrate for its processing, F-PSCs used flexible substrates like poly ethylene terephthalate (PET) which could not endure at high processing temperature significantly higher than 200°C. Hence various low-temperature techniques have been developed to fabricate electronic selective layers (ESLs) with great progress in attaining high power conversion efficiency (PCE) [31]. The two significant aspects to attain high efficiency in F-PSCs are the low temperature ESLs and the high quality perovskite absorbers that include grain size, trap density, charge transport, carrier lifetime, etc. The F-PSCs give much lower PCE compared with rigid counterparts though they are having similar process used to deposit perovskite absorber film. It requires different deposition methods to achieve desirable power conversion efficiency of perovskite films especially for larger area F-PSCs. The quality of the perovskite film was significantly improved by the inclusion of an additive dimethyl sulfide (DS) thereby enhancing both the grain size and crystallinity. When DS was introduced into the perovskite precursor solution, it chelates with Pb2+ to form an intermediate complex resulting in smaller Gibbs free energy and slower the rate of growth for perovskite to crystallize. Further the trap density of the perovskite film is reduced as the chelation interaction efficiently retards transformation kinetics during the thin film crystallization process, hence PCE of F-PSCs is increased to as high as 18.40%, the highest reported value so far for the F-PSCs. It is also expected that large area F-PSCs may also give improved efficiency as 13.35% [32].

devices requires the development of suitable transparent conducting electrodes which have high transparency, conductivity and process compatibility with prior device fabrication steps

Deposition of transparent conductive oxides (TCO), such as indium tin oxide (ITO), aluminum doped zinc oxide (AZO), or indium zinc oxide (IZO) acts as an electrode. Halide perovskites are known to have solvent and temperature instabilities hence physical vapor deposition processes are preferred rather than sputtering and spin coating. Thermally evaporated molybdenum oxide (MoOx) and thin Ag are the two existing material systems that have been reported as potential buffer layers for semitransparent devices [35]. PCE of 16.0% can be achieved and an average transparency of 54% in the near infrared region using thin Ag. We can apply the same semi-transparent cells in a 4-terminal (4T) tandem configuration with Cu(In,Ga)Se (CIGS) cell and attained a tandem efficiency of 20.7%. Although a tandem with Si cells would yield higher efficiency, both perovskite and CIGS are thin film technologies which are lightweight and can be deposited on flexible substrate. Further they exhibit radiation hardness but capable to withstand radiation levels several orders higher than crystalline

The recent surge for interest towards the ultra-high band gap absorbers for tandem solar cells from the oldest material selenium with a band gap of 1.95 eV. Moreover, Se devices are air-stable, non-toxic, and extremely simple to fabricate. Se solar cell utilized n-type TiO<sup>2</sup>

ciency of 5% lasts same for more than 30 years. This structure is similar to hybrid perovskite solar cells without the advanced hole transporting layer [37]. Now the redesigned Se device for its improvement in terms of its efficiency can be modified in three aspects. First, introducing a reliable inorganic MoOx (molybdenum oxide) with high-work-function as a holeselective layer between selenium and the gold back contact in order to reduce recombination and improve collection as it is been in both CdTe solar cells as well as other inorganic and organic photovoltaics [38]. Second method by reducing the thickness of the selenium absorber to only 100 nm—20 times less than the previous Se champion cell as well as typical chalcogenide absorbers such as copper indium gallium selenide (CIGS). Finally optimizing the buffer layer to reduce the cliff at interface should significantly improve Voc (open circuit voltage) as well as fill factor. Though Se an attractive alternative candidate for a high-band-gap absorber they are prone to practical applications by simple and inexpensive fabrication process, lack of highly toxic elements such as Cd and Pb, and stability upon prolonged storage and air exposure for these devices. Thus now Se solar cells find its new outlook after three decades

IEC Technical Committee TC82 was established in 1981. It is the most important International body regarding photovoltaic related standardization. The main task of TC82 is to prepare

(FTO) coated glass and p-type Se followed by gold contact with its effi-

,

New Materials for Thin Film Solar Cells http://dx.doi.org/10.5772/intechopen.81393 217

Si, lending them suitable for high altitudes and space applications [36].

with slight modification which suits well for our future energy crisis.

**4. International standards for solar panels**

is a major perspective [34].

**3.3. Ultra high band gap solar cells**

deposited on SnO<sup>2</sup>

Deployment of flexible PV technology is not only motivated by the quest for high-throughput and low-cost manufacturing but on the view for marketing it would be able to access with its eminent properties as it is being flexible, thin and lightweight, which would make it easy to integrate or apply on any surface or structure (either rigid, curved or flexible) [33] and even have its applications in portable and indoor electronics.

#### **3.2. CIGS perovskite tandem solar cells**

The development of high efficiency semi-transparent perovskite solar cells is necessary for the application in integrated photovoltaics and tandem solar cells. Tandem solar cells allow higher efficiencies than single-junction solar cells by better utilizing the energy of shortwavelength photons in the spectrum of sunlight. Top cells comprising a high-bandgap semiconductor to generate photocurrent at high voltage from the short-wavelength part of the solar spectrum. Longer-wavelength light, beyond the bandgap of the top cell, is transmitted to an underlying bottom cell comprising a lower-bandgap semiconductor with broad absorption coefficient. Tandem cells that are unique to inorganic-organic metal-halide-perovskite materials, particularly due to their bandgap-tuneability and luminescence efficiency thus an excellent candidates for applications in building integrated photovoltaics (BIPV) and tandem solar cells. Rapid growth in perovskite solar cells already reached its highest efficiency of 22.1%. However the translation of such high efficiencies to semitransparent perovskite devices requires the development of suitable transparent conducting electrodes which have high transparency, conductivity and process compatibility with prior device fabrication steps is a major perspective [34].

Deposition of transparent conductive oxides (TCO), such as indium tin oxide (ITO), aluminum doped zinc oxide (AZO), or indium zinc oxide (IZO) acts as an electrode. Halide perovskites are known to have solvent and temperature instabilities hence physical vapor deposition processes are preferred rather than sputtering and spin coating. Thermally evaporated molybdenum oxide (MoOx) and thin Ag are the two existing material systems that have been reported as potential buffer layers for semitransparent devices [35]. PCE of 16.0% can be achieved and an average transparency of 54% in the near infrared region using thin Ag. We can apply the same semi-transparent cells in a 4-terminal (4T) tandem configuration with Cu(In,Ga)Se (CIGS) cell and attained a tandem efficiency of 20.7%. Although a tandem with Si cells would yield higher efficiency, both perovskite and CIGS are thin film technologies which are lightweight and can be deposited on flexible substrate. Further they exhibit radiation hardness but capable to withstand radiation levels several orders higher than crystalline Si, lending them suitable for high altitudes and space applications [36].

### **3.3. Ultra high band gap solar cells**

**3. Recent trends in solar cells**

216 Coatings and Thin-Film Technologies

**3.1. Flexible perovskite solar cells (F-PSCs)**

Now research has been moved towards a flexible PSCs from rigid PSCs and it achieved notable milestones in terms of its efficiency with light weight [30]. Unlike rigid PSCs employ glass substrate for its processing, F-PSCs used flexible substrates like poly ethylene terephthalate (PET) which could not endure at high processing temperature significantly higher than 200°C. Hence various low-temperature techniques have been developed to fabricate electronic selective layers (ESLs) with great progress in attaining high power conversion efficiency (PCE) [31]. The two significant aspects to attain high efficiency in F-PSCs are the low temperature ESLs and the high quality perovskite absorbers that include grain size, trap density, charge transport, carrier lifetime, etc. The F-PSCs give much lower PCE compared with rigid counterparts though they are having similar process used to deposit perovskite absorber film. It requires different deposition methods to achieve desirable power conversion efficiency of perovskite films especially for larger area F-PSCs. The quality of the perovskite film was significantly improved by the inclusion of an additive dimethyl sulfide (DS) thereby enhancing both the grain size and crystallinity. When DS was introduced into the perovskite precursor solution, it chelates with Pb2+ to form an intermediate complex resulting in smaller Gibbs free energy and slower the rate of growth for perovskite to crystallize. Further the trap density of the perovskite film is reduced as the chelation interaction efficiently retards transformation kinetics during the thin film crystallization process, hence PCE of F-PSCs is increased to as high as 18.40%, the highest reported value so far for the F-PSCs. It is also

expected that large area F-PSCs may also give improved efficiency as 13.35% [32].

have its applications in portable and indoor electronics.

**3.2. CIGS perovskite tandem solar cells**

Deployment of flexible PV technology is not only motivated by the quest for high-throughput and low-cost manufacturing but on the view for marketing it would be able to access with its eminent properties as it is being flexible, thin and lightweight, which would make it easy to integrate or apply on any surface or structure (either rigid, curved or flexible) [33] and even

The development of high efficiency semi-transparent perovskite solar cells is necessary for the application in integrated photovoltaics and tandem solar cells. Tandem solar cells allow higher efficiencies than single-junction solar cells by better utilizing the energy of shortwavelength photons in the spectrum of sunlight. Top cells comprising a high-bandgap semiconductor to generate photocurrent at high voltage from the short-wavelength part of the solar spectrum. Longer-wavelength light, beyond the bandgap of the top cell, is transmitted to an underlying bottom cell comprising a lower-bandgap semiconductor with broad absorption coefficient. Tandem cells that are unique to inorganic-organic metal-halide-perovskite materials, particularly due to their bandgap-tuneability and luminescence efficiency thus an excellent candidates for applications in building integrated photovoltaics (BIPV) and tandem solar cells. Rapid growth in perovskite solar cells already reached its highest efficiency of 22.1%. However the translation of such high efficiencies to semitransparent perovskite The recent surge for interest towards the ultra-high band gap absorbers for tandem solar cells from the oldest material selenium with a band gap of 1.95 eV. Moreover, Se devices are air-stable, non-toxic, and extremely simple to fabricate. Se solar cell utilized n-type TiO<sup>2</sup> , deposited on SnO<sup>2</sup> (FTO) coated glass and p-type Se followed by gold contact with its efficiency of 5% lasts same for more than 30 years. This structure is similar to hybrid perovskite solar cells without the advanced hole transporting layer [37]. Now the redesigned Se device for its improvement in terms of its efficiency can be modified in three aspects. First, introducing a reliable inorganic MoOx (molybdenum oxide) with high-work-function as a holeselective layer between selenium and the gold back contact in order to reduce recombination and improve collection as it is been in both CdTe solar cells as well as other inorganic and organic photovoltaics [38]. Second method by reducing the thickness of the selenium absorber to only 100 nm—20 times less than the previous Se champion cell as well as typical chalcogenide absorbers such as copper indium gallium selenide (CIGS). Finally optimizing the buffer layer to reduce the cliff at interface should significantly improve Voc (open circuit voltage) as well as fill factor. Though Se an attractive alternative candidate for a high-band-gap absorber they are prone to practical applications by simple and inexpensive fabrication process, lack of highly toxic elements such as Cd and Pb, and stability upon prolonged storage and air exposure for these devices. Thus now Se solar cells find its new outlook after three decades with slight modification which suits well for our future energy crisis.

### **4. International standards for solar panels**

IEC Technical Committee TC82 was established in 1981. It is the most important International body regarding photovoltaic related standardization. The main task of TC82 is to prepare international standards for systems of photovoltaic conversion of solar energy in to electrical energy and for all the elements in the entire photovoltaic energy system.

[2] Jean J, Brown PR, Jaffe RL, Buonassisi T, Bulović V. Pathways for solar photovoltaics.

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### **5. Conclusion**

Thin film technologies made a remarkable advent in the photovoltaic industry for the manufacturing of solar cells. Thin film solar cells offer a most promising options for reducing the cost of photovoltaic systems. Enormous progress in device performance has been made in dye sensitized to perovskite solar cells. Perovskite solar cells open the door for novel applications in the production of solar cells with increase in stability and reliability. Major concern is about the toxicity of lead and its impact in the environment. Then a major task is finding a material which would replace the lead. Tin can replace lead but the stability issues still pertain and the efficiency achieved was low compared with lead. The device performance suits best in N<sup>2</sup> gas atmosphere to avoid quick degradation of tin. Therefore a thirst for new material without lead such as BiFeO<sup>3</sup> , BiFe<sup>2</sup> CrO6 and BiMnO<sup>3</sup> . The overall power conversion efficiency is lower than metal halides but they are more stable. Hence it been able to investigate bismuth based photovoltaic materials. Cesium with bismuth, for example, Cs<sup>3</sup> Bi2 I 9 will give the highest solar cell performance with low toxicity and environment impact. It also proves to be a predominant in the fulfillment of our future energy crisis in terms of its simplicity in manufacturing process and low cost. Moreover there is also an opportunity to develop high-performance tandem cell technology that uses both perovskite and existing technologies and this may allow market introduction as a new premium product.

### **Author details**

Senthil T.S\* and Kalaiselvi C.R

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

Department of Physics, Erode Sengunthar Engineering College, Perundurai, India

### **References**

[1] Vilot JP, Ayachi B, Aviles T, Miska P. Full sputtering deposition of thin film solar cells: A way of achieving high efficiency sustainable tandem cells. Journal of Electronic Materials. 2017;**46**(11):6523-6527

[2] Jean J, Brown PR, Jaffe RL, Buonassisi T, Bulović V. Pathways for solar photovoltaics. Energy & Environmental Science. 2015;**8**:1200-1219

international standards for systems of photovoltaic conversion of solar energy in to electrical

• ASTM E44-ASTM Committee E44 on Solar, geothermal and other alternative energy.

• IEEE SCC21-IEEE SCC21 standards Coordinating Committee on fuel cells, photo-voltaic,

Thin film technologies made a remarkable advent in the photovoltaic industry for the manufacturing of solar cells. Thin film solar cells offer a most promising options for reducing the cost of photovoltaic systems. Enormous progress in device performance has been made in dye sensitized to perovskite solar cells. Perovskite solar cells open the door for novel applications in the production of solar cells with increase in stability and reliability. Major concern is about the toxicity of lead and its impact in the environment. Then a major task is finding a material which would replace the lead. Tin can replace lead but the stability issues still pertain and the efficiency achieved was

more stable. Hence it been able to investigate bismuth based photovoltaic materials. Cesium with

environment impact. It also proves to be a predominant in the fulfillment of our future energy crisis in terms of its simplicity in manufacturing process and low cost. Moreover there is also an opportunity to develop high-performance tandem cell technology that uses both perovskite and existing technologies and this may allow market introduction as a new premium product.

. The overall power conversion efficiency is lower than metal halides but they are

will give the highest solar cell performance with low toxicity and

gas atmosphere to avoid quick

, BiFe<sup>2</sup>

CrO6

energy and for all the elements in the entire photovoltaic energy system.

• ISO TC 180-ISO Technical Committee 180, solar energy.

low compared with lead. The device performance suits best in N<sup>2</sup>

Bi2 I 9

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

Materials. 2017;**46**(11):6523-6527

degradation of tin. Therefore a thirst for new material without lead such as BiFeO<sup>3</sup>

Department of Physics, Erode Sengunthar Engineering College, Perundurai, India

[1] Vilot JP, Ayachi B, Aviles T, Miska P. Full sputtering deposition of thin film solar cells: A way of achieving high efficiency sustainable tandem cells. Journal of Electronic

dispersed generation and energy storage.

**5. Conclusion**

218 Coatings and Thin-Film Technologies

and BiMnO<sup>3</sup>

bismuth, for example, Cs<sup>3</sup>

**Author details**

**References**

Senthil T.S\* and Kalaiselvi C.R

• IEC TC82-IEC Technical Committee 82, solar photovoltaic energy system.


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**Chapter 11**

**Provisional chapter**

**Organometal Halide Perovskites Thin Film and Their**

**Organometal Halide Perovskites Thin Film and Their** 

The organometal halide perovskite solar cells (PSCs) have attracted attention and achieved efficiencies compared with traditional solar cells. There are several ways to develop perovskite solar cells like effect of moisture, degradation, and understanding the reason for instability of perovskite. In this chapter, we are specified how to make coating and film fabrication are affected by the existing methods. Improvement in the photovoltaic performance of PSCs can be achieved by enhanced processing technique. These

and crystal quality of the morphology for perovskite films. There is no doubt that film coating indicates that the crystallization and morphology of perovskite films affect the absorption intensity and obviously influence the short-circuit current density. This study points out an enhancement of the stability of perovskite films and solar cells by reducing

**Keywords:** perovskite solar cells, working mechanism, photovoltaic parameters,

Metal halide perovskite solar cells (PSCs) have emerged as a kind of encouraging alternative to existing photovoltaic technologies with both solution processability and superior photovoltaic performances. Fundamental studies on perovskite materials [1], device designs [2, 3], fabrication processes [4–10], and materials engineering [11–16] have boosted the rapid development of PSCs. Consequently, a certified power conversion efficiency (PCE) of 22.1%

solution controlling the substrate temperature

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

© 2018 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use,

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

DOI: 10.5772/intechopen.79678

**Impact on the Efficiency of Perovskite Solar Cells**

**Impact on the Efficiency of Perovskite Solar Cells**

Ahmed Mourtada Elseman

Ahmed Mourtada Elseman

**Abstract**

http://dx.doi.org/10.5772/intechopen.79678

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

techniques include the spin-coating PbI<sup>2</sup>

residual strains in perovskite films.

stability, low cost

**1. Introduction**

#### **Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite Solar Cells Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite Solar Cells**

DOI: 10.5772/intechopen.79678

Ahmed Mourtada Elseman Ahmed Mourtada Elseman

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

http://dx.doi.org/10.5772/intechopen.79678

#### **Abstract**

The organometal halide perovskite solar cells (PSCs) have attracted attention and achieved efficiencies compared with traditional solar cells. There are several ways to develop perovskite solar cells like effect of moisture, degradation, and understanding the reason for instability of perovskite. In this chapter, we are specified how to make coating and film fabrication are affected by the existing methods. Improvement in the photovoltaic performance of PSCs can be achieved by enhanced processing technique. These techniques include the spin-coating PbI<sup>2</sup> solution controlling the substrate temperature and crystal quality of the morphology for perovskite films. There is no doubt that film coating indicates that the crystallization and morphology of perovskite films affect the absorption intensity and obviously influence the short-circuit current density. This study points out an enhancement of the stability of perovskite films and solar cells by reducing residual strains in perovskite films.

**Keywords:** perovskite solar cells, working mechanism, photovoltaic parameters, stability, low cost

### **1. Introduction**

Metal halide perovskite solar cells (PSCs) have emerged as a kind of encouraging alternative to existing photovoltaic technologies with both solution processability and superior photovoltaic performances. Fundamental studies on perovskite materials [1], device designs [2, 3], fabrication processes [4–10], and materials engineering [11–16] have boosted the rapid development of PSCs. Consequently, a certified power conversion efficiency (PCE) of 22.1%

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2018 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

has been obtained after the past several years of vigorous work. However, despite the overwhelming achievements in terms of performance of PSCs, the long-term stability and current– voltage hysteresis still remain critical [4].

**2.1. The mesoporous scaffold**

of uniform and dense MAPbI<sup>3</sup>

A simple binder-free colloid Al<sup>2</sup>

on top of the scaffold thin porous Al<sup>2</sup>

mesoporous film and (d) planar heterojunction structure.

**Figure 2.** (a) Schematic cross section of WO<sup>3</sup>

cells and (b) energy band diagram [34].

trated mesoporous TiO<sup>2</sup>

MAPbBr<sup>3</sup>

The best performance for PSCs have reported and reached a certified PCE of 22.1%, have depend on high temperature processing (450–550°C) and the mesoporous structure used TiO<sup>2</sup> as ETLs [29]. Moreover, in another work, a 16.2% efficiency was obtained from the combination

and perovskite composite layer (200 nm)/perovskite upper layer (300 nm)/PTAA (50 nm)/Au

Furthermore, it has been revealed that a solid thin film of the perovskite absorber was formed

both carrier species with an internal quantum efficiency approaching 100% [33]. Mahmood et al. reported two-dimensional (2D) nanosheets with enhanced absorber infiltration as compared with

nanostructured porous ETL to obtain highly efficient perovskite solar cells with PCE of 11.2% [34].

**Figure 1.** Development of device configuration in perovskite solar cells [28]. These figures reveal (a) a sensitization concept, (b) extremely thin layer of perovskite deposited on mesoporous scaffold layer, (c) perovskite infiltration into

bilayer architecture consisting of perovskite-infil-

nanoparticle meso-superstructured scaffold with anneal-

films. This supported charge separation and transport of

n-type semiconductor as a nanostructured porous ETL for perovskite solar

(70 nm)/mesoporous TiO<sup>2</sup>

n-type semiconductor as a

layer delivered PCEs of up to 12.3%.

http://dx.doi.org/10.5772/intechopen.79678


225

electrodes [30]. Similarly, the same group proceeded the FAPbI<sup>3</sup>

Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite...

and MAPbBr<sup>3</sup>

(100 nm) with a PCE of 18.4% at maximum power point condition [31, 32].

O3

O3

1D nanostructures as revealed in **Figure 2**. The author utilized WO<sup>3</sup>

ing temperature of 150°C deposited over compact TiO<sup>2</sup>

system with an architecture as follows: FTO/blocking TiO<sup>2</sup>

Perovskite solar cells, the most promising new technology in the academia and industry, have promised a highly competitive alternative to silicon solar cells and other commercial alternatives. Perovskite solar cells are high-performance photovoltaic devices which have the potential to enter in the market in the near future. Low processing costs and highly abundant raw materials may permit a short-energy payback time and low overall CO emissions. After an impressive increase in PCE from ~10% in 2012 to ~22.1% in early 2016, experts expect to discover further improvements in efficiency in the next several years [12]. Perovskite solar cell research is still in its infancy considering that the first work was only published in 2009 [1]. The commercialization of perovskite solar cell needs to address several fundamental issues in the near future: for example, control the growth of thin film and deposition, make scale and numerous process, achieve high stability and long lifetime, and low toxicity. To be competitive, cost will be a concern for manufacturing companies. Although the raw materials for making perovskite solar cells are inexpensive and abundant, recent analyses of costperformance and commercialization requirements are not entirely positive [17]. To make perovskite solar cells competitive, several goals are needed to be achieved. For example, the levelized cost of electricity (LCOE) for residential use is 9.0 cents per kWh by 2020 and is expected to decline to 5.0 cents/kWh by 2030 [18]. This is a huge challenge for perovskite technology at present.

In this review, we summarize recently developed perovskite film deposition techniques and evaluate their suitability for industrial production of perovskite solar cells and modules. Our discussion of stability and device lifetime focuses mainly on the measurement standards and issues relative to commercialization. Thereafter, we present techniques that used to fabricate the perovskite solar cells such as one-step spin-coating and two-step deposition techniques, solvent-solvent extraction, vapor-assisted solution processes, dual-source vacuum deposition, hybrid deposition, hybrid chemical vapor deposition, sequential vapor deposition, and flash evaporation. The control of the morphology for perovskite thin films has been observed by many efforts.

### **2. The common architecture of PSC**

Generally, PSC has three main types of device architectures: (i) mesoscopic structures using mesoporous semiconducting materials as electron-transporting layers (ETLs), i.e., TiO<sup>2</sup> [19], WO<sup>3</sup> [20], SrTiO<sup>3</sup> [21], ZnO [22], Zn<sup>2</sup> SnO<sup>4</sup> [23], and SnO<sup>2</sup> [24]); (ii) meso-superstructures employing mesoporous insulators such as Al<sup>2</sup> O3 [25] and ZrO<sup>2</sup> [26] as scaffolds, while perovskite itself acts as ETL; and (iii) planar structures implementing ultrathin compact layer materials for both hole-blocking and electron-conducting purposes. Among these types, the planar PSC was motivated by the requests of more simple process and lower cost for future applications [27]. **Figure 1** reveals the evolution of device configuration as sequence from mesostructure to planar heterojunction. The categories related to the types of device architecture were discussed here in summary.

### **2.1. The mesoporous scaffold**

has been obtained after the past several years of vigorous work. However, despite the overwhelming achievements in terms of performance of PSCs, the long-term stability and current–

Perovskite solar cells, the most promising new technology in the academia and industry, have promised a highly competitive alternative to silicon solar cells and other commercial alternatives. Perovskite solar cells are high-performance photovoltaic devices which have the potential to enter in the market in the near future. Low processing costs and highly abundant raw materials may permit a short-energy payback time and low overall CO emissions. After an impressive increase in PCE from ~10% in 2012 to ~22.1% in early 2016, experts expect to discover further improvements in efficiency in the next several years [12]. Perovskite solar cell research is still in its infancy considering that the first work was only published in 2009 [1]. The commercialization of perovskite solar cell needs to address several fundamental issues in the near future: for example, control the growth of thin film and deposition, make scale and numerous process, achieve high stability and long lifetime, and low toxicity. To be competitive, cost will be a concern for manufacturing companies. Although the raw materials for making perovskite solar cells are inexpensive and abundant, recent analyses of costperformance and commercialization requirements are not entirely positive [17]. To make perovskite solar cells competitive, several goals are needed to be achieved. For example, the levelized cost of electricity (LCOE) for residential use is 9.0 cents per kWh by 2020 and is expected to decline to 5.0 cents/kWh by 2030 [18]. This is a huge challenge for perovskite

In this review, we summarize recently developed perovskite film deposition techniques and evaluate their suitability for industrial production of perovskite solar cells and modules. Our discussion of stability and device lifetime focuses mainly on the measurement standards and issues relative to commercialization. Thereafter, we present techniques that used to fabricate the perovskite solar cells such as one-step spin-coating and two-step deposition techniques, solvent-solvent extraction, vapor-assisted solution processes, dual-source vacuum deposition, hybrid deposition, hybrid chemical vapor deposition, sequential vapor deposition, and flash evaporation. The control of the morphology for perovskite thin films has been observed

Generally, PSC has three main types of device architectures: (i) mesoscopic structures using mesoporous semiconducting materials as electron-transporting layers (ETLs), i.e., TiO<sup>2</sup>

perovskite itself acts as ETL; and (iii) planar structures implementing ultrathin compact layer materials for both hole-blocking and electron-conducting purposes. Among these types, the planar PSC was motivated by the requests of more simple process and lower cost for future applications [27]. **Figure 1** reveals the evolution of device configuration as sequence from mesostructure to planar heterojunction. The categories related to the types of device architec-

[23], and SnO<sup>2</sup>

[25] and ZrO<sup>2</sup>

O3

SnO<sup>4</sup>

[19],

[24]); (ii) meso-superstructures

[26] as scaffolds, while

voltage hysteresis still remain critical [4].

224 Coatings and Thin-Film Technologies

technology at present.

by many efforts.

[20], SrTiO<sup>3</sup>

WO<sup>3</sup>

**2. The common architecture of PSC**

[21], ZnO [22], Zn<sup>2</sup>

employing mesoporous insulators such as Al<sup>2</sup>

ture were discussed here in summary.

The best performance for PSCs have reported and reached a certified PCE of 22.1%, have depend on high temperature processing (450–550°C) and the mesoporous structure used TiO<sup>2</sup> as ETLs [29]. Moreover, in another work, a 16.2% efficiency was obtained from the combination of uniform and dense MAPbI<sup>3</sup> and MAPbBr<sup>3</sup> bilayer architecture consisting of perovskite-infiltrated mesoporous TiO<sup>2</sup> electrodes [30]. Similarly, the same group proceeded the FAPbI<sup>3</sup> - MAPbBr<sup>3</sup> system with an architecture as follows: FTO/blocking TiO<sup>2</sup> (70 nm)/mesoporous TiO<sup>2</sup> and perovskite composite layer (200 nm)/perovskite upper layer (300 nm)/PTAA (50 nm)/Au (100 nm) with a PCE of 18.4% at maximum power point condition [31, 32].

A simple binder-free colloid Al<sup>2</sup> O3 nanoparticle meso-superstructured scaffold with annealing temperature of 150°C deposited over compact TiO<sup>2</sup> layer delivered PCEs of up to 12.3%. Furthermore, it has been revealed that a solid thin film of the perovskite absorber was formed on top of the scaffold thin porous Al<sup>2</sup> O3 films. This supported charge separation and transport of both carrier species with an internal quantum efficiency approaching 100% [33]. Mahmood et al. reported two-dimensional (2D) nanosheets with enhanced absorber infiltration as compared with 1D nanostructures as revealed in **Figure 2**. The author utilized WO<sup>3</sup> n-type semiconductor as a nanostructured porous ETL to obtain highly efficient perovskite solar cells with PCE of 11.2% [34].

**Figure 1.** Development of device configuration in perovskite solar cells [28]. These figures reveal (a) a sensitization concept, (b) extremely thin layer of perovskite deposited on mesoporous scaffold layer, (c) perovskite infiltration into mesoporous film and (d) planar heterojunction structure.

**Figure 2.** (a) Schematic cross section of WO<sup>3</sup> n-type semiconductor as a nanostructured porous ETL for perovskite solar cells and (b) energy band diagram [34].

### **2.2. The planar heterojunction**

Planar PSCs have been demonstrated by Tan et al. with smooth and pinhole-free TiO<sup>2</sup> -Cl as the ETL. The film also exhibited negligible parasitic absorption loss over the entire visible to near-infrared spectrum. Solar cells fabricated on TiO<sup>2</sup> -Cl exhibit considerably better performance than those on TiO<sup>2</sup> for all PV parameters. Correspondingly, TiO<sup>2</sup> -Cl resulted in a higher average PCE (19.8%) than the Cl-free TiO<sup>2</sup> (15.8%) [35, 36]. The best-performing smallarea CsMAFA solar cell (0.049 cm<sup>2</sup> ) exhibited a high laboratory PCE of 21.4% without hysteresis in J-V sweeps. Similarly, large-area (1.1 cm<sup>2</sup> ) cells fabricated on TiO<sup>2</sup> -Cl showed a PCE value >20% with negligible hysteresis. Zhou et al. manipulated carrier behavior with planar heterojunction perovskite solar cells. Yttrium-doped TiO<sup>2</sup> (Y-TiO<sup>2</sup> ) was fabricated at <150°C annealing temperature as the ETL to enhance electron extraction and transport over reduced work function ITO treated with polyethylenimine ethoxylated (PEIE) solution. These treatments produced a PCE of 19.3% [37, 38]. The use of CdSe nanoparticles (solution processed at 150°C) has been investigated to replace the widely used TiO<sup>2</sup> as the ETLs for the conventional planar heterojunction PSCs. Devices with CdSe nanoparticle ETLs were performed well, with the PCE of 11.7% [39]. A planar device with high PCE was applicable through Hagfeldt et al. and Jiang et al., using SnO<sup>2</sup> [24, 40–42].

**3. Film formation of ETL**

ods have been discussed.

**3.1. Vacuum thermal co-evaporation**

**3.2. Layer-by-layer sequential vacuum sublimation**

**3.3. Vapor deposition by dual source**

There is no doubt that the crystallinity, thickness, material film morphology, and purity have impact on the efficiency of solar cell performance. The film formation has been relied on deposition techniques such as one-step spin-coating [7, 46–50] and two-step deposition techniques [51, 52], solvent-solvent extraction [53], vapor-assisted solution processes [54–57], dual-source vacuum deposition [58–61], hybrid deposition [62, 63], hybrid chemical vapor deposition [62, 64–66], sequential vapor deposition [67, 68], and flash evaporation [69]. The control of the morphology for perovskite thin films has been observed by many efforts. These efforts include the optimization of the annealing time, temperature [70, 71], selection of the underlayer material and thickness [3, 72–75], and the use of alternative deposition methods such as two-step deposition and vacuum sublimation [4, 76, 77]. Herein, some of these meth-

Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite...

http://dx.doi.org/10.5772/intechopen.79678

In this method, the authors used vacuum thermal co-evaporation of organic halide and metal halide to resulting perovskite thin films with homogeneous morphology and improved thinfilm coverage. These results achieved a high performance of 12–15% PCE [5, 78]. Despite the promising results, however, to date, only limited reports have utilized this vacuum sublimation technique to fabricate perovskite layers [5, 78]. The main reason could be due to the small molecular weight of organic halide and make the monitoring without control of the CH<sup>3</sup>

deposition rate using quartz microbalance sensors [5, 78]. In another way, Zhu et al. reported how to develop this technique deposition to fabricate pinhole-free cesium-substituted perovskite films and enhance the surface coverage as shown in **Figure 4**. The same method used to promise tunable bandgap reduced trap-state density and longer carrier lifetime, with efficiency 20.13%, which is the highest fabrication for planar perovskite solar cells [79].

Chen et al. reported a novel method of perovskite thin-film deposition via a layer-by-layer sequential vacuum sublimation. This method has been easier than the previous technique. The very uniform perovskite thin films can achieve high coverage via incorporating the thin films of perovskite with a poly(3,4-ethylenedioxythiophene):poly(styrene sulfonate) (PEDOT:PSS) hole-transporting layer (HTL) and thermally evaporated C60/bathophenanthroline (Bphen) electron-transporting layers (ETLs). The cells here attain efficiencies as high as 15.4% because

The large-scale production in optoelectronic applications has been achieved by vapor deposition techniques because this technique is widely used in semiconductor industry. The feasibility of organometal halide perovskite materials via vapor deposition techniques

the devices were free of high-temperature-prepared metal oxide layers [80].

NH<sup>3</sup> I 227

By modifying the surface of a planar structure of the TiO<sup>2</sup> compact layer with C60-SAM molecules (Wojciechowski et al.), a PCE of 15.7% has been obtained (**Figure 3**) [43]. Pablo Docampo et al. have demonstrated 10% PCEs for inverted planar PSCs with bilayer of PC60BM and compact TiOx as ETL [44]. Further, hysteresis-less inverted planar hybrid solar cells with 18.1% PCE has been fabricated by Heo et al. Better PCE and stability were attributed to the electron extraction from MAPbI<sup>3</sup> into PCBM, the increased EQE value by the better charge injection/separation efficiency, and the improved FF by the increased diffusion coefficient (Dn) and charge carrier lifetime (*τ*n). In addition, the air and humid stability was improved by the corrosive additive-free device architecture and hydrophobicity of the PCBM top layer [45].

**Figure 3.** Planar structure of the TiO<sup>2</sup> compact layer with C60-SAM molecules [43].

### **3. Film formation of ETL**



(15.8%) [35, 36]. The best-performing small-

) exhibited a high laboratory PCE of 21.4% without hyster-

) cells fabricated on TiO<sup>2</sup>

(Y-TiO<sup>2</sup>

as ETL [44]. Further, hysteresis-less inverted planar hybrid solar



) was fabricated at <150°C

as the ETLs for the conventional

compact layer with C60-SAM

into PCBM, the increased EQE value by the

**2.2. The planar heterojunction**

226 Coatings and Thin-Film Technologies

performance than those on TiO<sup>2</sup>

and Jiang et al., using SnO<sup>2</sup>

PC60BM and compact TiOx

PCBM top layer [45].

**Figure 3.** Planar structure of the TiO<sup>2</sup>

area CsMAFA solar cell (0.049 cm<sup>2</sup>

Planar PSCs have been demonstrated by Tan et al. with smooth and pinhole-free TiO<sup>2</sup>

ible to near-infrared spectrum. Solar cells fabricated on TiO<sup>2</sup>

heterojunction perovskite solar cells. Yttrium-doped TiO<sup>2</sup>

150°C) has been investigated to replace the widely used TiO<sup>2</sup>

[24, 40–42].

By modifying the surface of a planar structure of the TiO<sup>2</sup>

uted to the electron extraction from MAPbI<sup>3</sup>

higher average PCE (19.8%) than the Cl-free TiO<sup>2</sup>

esis in J-V sweeps. Similarly, large-area (1.1 cm<sup>2</sup>

as the ETL. The film also exhibited negligible parasitic absorption loss over the entire vis-

value >20% with negligible hysteresis. Zhou et al. manipulated carrier behavior with planar

annealing temperature as the ETL to enhance electron extraction and transport over reduced work function ITO treated with polyethylenimine ethoxylated (PEIE) solution. These treatments produced a PCE of 19.3% [37, 38]. The use of CdSe nanoparticles (solution processed at

planar heterojunction PSCs. Devices with CdSe nanoparticle ETLs were performed well, with the PCE of 11.7% [39]. A planar device with high PCE was applicable through Hagfeldt et al.

molecules (Wojciechowski et al.), a PCE of 15.7% has been obtained (**Figure 3**) [43]. Pablo Docampo et al. have demonstrated 10% PCEs for inverted planar PSCs with bilayer of

cells with 18.1% PCE has been fabricated by Heo et al. Better PCE and stability were attrib-

better charge injection/separation efficiency, and the improved FF by the increased diffusion coefficient (Dn) and charge carrier lifetime (*τ*n). In addition, the air and humid stability was improved by the corrosive additive-free device architecture and hydrophobicity of the

compact layer with C60-SAM molecules [43].

for all PV parameters. Correspondingly, TiO<sup>2</sup>

There is no doubt that the crystallinity, thickness, material film morphology, and purity have impact on the efficiency of solar cell performance. The film formation has been relied on deposition techniques such as one-step spin-coating [7, 46–50] and two-step deposition techniques [51, 52], solvent-solvent extraction [53], vapor-assisted solution processes [54–57], dual-source vacuum deposition [58–61], hybrid deposition [62, 63], hybrid chemical vapor deposition [62, 64–66], sequential vapor deposition [67, 68], and flash evaporation [69]. The control of the morphology for perovskite thin films has been observed by many efforts. These efforts include the optimization of the annealing time, temperature [70, 71], selection of the underlayer material and thickness [3, 72–75], and the use of alternative deposition methods such as two-step deposition and vacuum sublimation [4, 76, 77]. Herein, some of these methods have been discussed.

### **3.1. Vacuum thermal co-evaporation**

In this method, the authors used vacuum thermal co-evaporation of organic halide and metal halide to resulting perovskite thin films with homogeneous morphology and improved thinfilm coverage. These results achieved a high performance of 12–15% PCE [5, 78]. Despite the promising results, however, to date, only limited reports have utilized this vacuum sublimation technique to fabricate perovskite layers [5, 78]. The main reason could be due to the small molecular weight of organic halide and make the monitoring without control of the CH<sup>3</sup> NH<sup>3</sup> I deposition rate using quartz microbalance sensors [5, 78]. In another way, Zhu et al. reported how to develop this technique deposition to fabricate pinhole-free cesium-substituted perovskite films and enhance the surface coverage as shown in **Figure 4**. The same method used to promise tunable bandgap reduced trap-state density and longer carrier lifetime, with efficiency 20.13%, which is the highest fabrication for planar perovskite solar cells [79].

### **3.2. Layer-by-layer sequential vacuum sublimation**

Chen et al. reported a novel method of perovskite thin-film deposition via a layer-by-layer sequential vacuum sublimation. This method has been easier than the previous technique. The very uniform perovskite thin films can achieve high coverage via incorporating the thin films of perovskite with a poly(3,4-ethylenedioxythiophene):poly(styrene sulfonate) (PEDOT:PSS) hole-transporting layer (HTL) and thermally evaporated C60/bathophenanthroline (Bphen) electron-transporting layers (ETLs). The cells here attain efficiencies as high as 15.4% because the devices were free of high-temperature-prepared metal oxide layers [80].

### **3.3. Vapor deposition by dual source**

The large-scale production in optoelectronic applications has been achieved by vapor deposition techniques because this technique is widely used in semiconductor industry. The feasibility of organometal halide perovskite materials via vapor deposition techniques

**Figure 4.** (a) Illustration of the vacuum co-evaporation, (b) deposition of the Cs-substituted MA1-xCs<sup>x</sup> PbI<sup>3</sup> perovskite thin film and (c) final perovskite thin film [79].

has advantages like the possibility to fabricate films with high purity; these techniques are more proper to prepare multilayered structures of thin films, and suitable optimization of perovskite films also can be deposited by vapor deposition [5, 81]. Liu et al. [5] reported preparation of CH<sup>3</sup> NH<sup>3</sup> PbI3−xCl<sup>x</sup> by dual-source vapor deposition technique in the presence of PbCl<sup>2</sup> and CH<sup>3</sup> NH<sup>3</sup> I. This method leads to high-efficiency photovoltaic devices of 15.4%. Malinkiewicz et al. used the same technique to deposited CH<sup>3</sup> NH<sup>3</sup> PbI<sup>3</sup> in the presence of PbI<sup>2</sup> and CH<sup>3</sup> NH<sup>3</sup> I and give uniform film formation with root-mean-square roughness of 5 nm measured by AFM [81]. In addition, the films showed uniform grainy structures with an average grain size of 150 nm [82]. Schematic illustration of the dual-source vacuum deposition process is shown in **Figure 5** [59].

**Figure 5.** Description of the dual-source vacuum deposition instrument (reproduced with permission from Ref. [59]).

Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite...

http://dx.doi.org/10.5772/intechopen.79678

229

**Figure 6.** Different techniques used for large-area perovskite film deposition: (a) spin-coating technique and (b) spray-

coating technique (reproduced with permission [6, 98]).

### **3.4. Spin coating**

Spin coating is widely used to fabricate a small area from thin films in lab scale. Spin coating used a small amount of solution, which was then dropped on the substrate as shown in **Figure 6a**. Then, the substrate has been covered by a layer of solution and spun to accelerate evaporation of the solvent [83]. This technique controls the thickness of the film by the concentration of the solution and speed [83]. In general, with regular spin coating, a one-step process with PbI<sup>2</sup> /MAI or PbCl<sup>2</sup> /MAI with gamma-butyrolactone (GBL), dimethylformamide (DMF), or dimethyl sulfoxide (DMSO) as solvents prompts poor film quality [71, 84]. Despite the fact that all preparing conditions have been considered, spin-coated perovskite film quality is regularly poor, with a high thickness of pinholes and little grain sizes. These pinholes cause shunt that debase the efficiency. With added substances designing the crystallization Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite... http://dx.doi.org/10.5772/intechopen.79678 229

**Figure 5.** Description of the dual-source vacuum deposition instrument (reproduced with permission from Ref. [59]).

has advantages like the possibility to fabricate films with high purity; these techniques are more proper to prepare multilayered structures of thin films, and suitable optimization of perovskite films also can be deposited by vapor deposition [5, 81]. Liu et al. [5] reported

**Figure 4.** (a) Illustration of the vacuum co-evaporation, (b) deposition of the Cs-substituted MA1-xCs<sup>x</sup>

measured by AFM [81]. In addition, the films showed uniform grainy structures with an average grain size of 150 nm [82]. Schematic illustration of the dual-source vacuum deposition

Spin coating is widely used to fabricate a small area from thin films in lab scale. Spin coating used a small amount of solution, which was then dropped on the substrate as shown in **Figure 6a**. Then, the substrate has been covered by a layer of solution and spun to accelerate evaporation of the solvent [83]. This technique controls the thickness of the film by the concentration of the solution and speed [83]. In general, with regular spin coating, a one-step

(DMF), or dimethyl sulfoxide (DMSO) as solvents prompts poor film quality [71, 84]. Despite the fact that all preparing conditions have been considered, spin-coated perovskite film quality is regularly poor, with a high thickness of pinholes and little grain sizes. These pinholes cause shunt that debase the efficiency. With added substances designing the crystallization

by dual-source vapor deposition technique in the presence

NH<sup>3</sup> PbI<sup>3</sup>

/MAI with gamma-butyrolactone (GBL), dimethylformamide

in the presence of PbI<sup>2</sup>

PbI<sup>3</sup>

perovskite thin

I. This method leads to high-efficiency photovoltaic devices of 15.4%.

I and give uniform film formation with root-mean-square roughness of 5 nm

preparation of CH<sup>3</sup>

NH<sup>3</sup>

**3.4. Spin coating**

process with PbI<sup>2</sup>

and CH<sup>3</sup>

film and (c) final perovskite thin film [79].

228 Coatings and Thin-Film Technologies

of PbCl<sup>2</sup>

and CH<sup>3</sup>

NH<sup>3</sup>

/MAI or PbCl<sup>2</sup>

NH<sup>3</sup>

process is shown in **Figure 5** [59].

PbI3−xCl<sup>x</sup>

Malinkiewicz et al. used the same technique to deposited CH<sup>3</sup>

**Figure 6.** Different techniques used for large-area perovskite film deposition: (a) spin-coating technique and (b) spraycoating technique (reproduced with permission [6, 98]).

of perovskite could be finely tuned, and perovskite films with altogether enhanced quality can be set up for superior power conversion [85, 86]. For instance, by utilizing lead acetic acid as lead source, the crystallization of perovskite is significantly speedier. The free pinhole perovskite films were shown by a basic one-step coating process [87].

ZrO<sup>2</sup>

0.28 cm<sup>2</sup>

trode and the TiO<sup>2</sup>

functions as a porous insulating layer to prevent direct contact between the carbon elec-

[26, 99]. **Figure 7** shows noncontact inkjet printing offering rapid and digital deposi-

technique, infiltration of perovskite precursor solution remains a challenge and is the main reason for lower efficiency compared to those of devices fabricated by other means. The most intriguing properties of this carbon-coated, printed, mesoscopic device are the high stability and outstanding outdoor performance. A certified PCE of 12.8% and stable performance over >1000 h in ambient air under full sunlight has been recorded for a device with an active area of

tion combined with excellent control over the layer formation for printed perovskite solar cells. Mathies et al. [100] reported that inkjet printing is used to deposit triple cation perovskite layers with 10% cesium in a mixed formamidinium/methylammonium lead iodide/bromide composite for solar cells with high temperature and moisture stability. A reliable process control over a wide range of perovskite layer thickness from 175 to 780 nm and corresponding grain sizes is achieved by adjusting the drop spacing of the inkjet printer cartridge. A continuous power output at constant voltage, resulting in a power conversion efficiency of 12.9%, is demonstrated, representing a major improvement from previously reported inkjetprinted methylammonium lead triiodide perovskite solar cells [100]. Compared with solution processes, dry deposition processes may be more environmentally friendly, as they do not require toxic solvents (DMSO, GBL, DMF, chloroform, chlorobenzene, isopropanol, toluene, diethyl ether, etc.), and they are compatible with high-quality, large-area perovskite film

deposition, such as vacuum deposition and chemical vapor deposition.

**Figure 7.** (a) 520 nm-thick inkjet-printed perovskite layer on FTO/TiO<sup>2</sup>

of the solar cell stack, denoting with the different layers: glass/FTO/TiO<sup>2</sup>

Spiro-MeOTAD (HTL)/Au.

printed perovskite solar cells. The substrate contains eight cells with each 3×3 mm<sup>2</sup>

(HTL)/Au [100]. (d) Real solar cell achieved 12% with the configuration: Glass/FTO/TiO<sup>2</sup>

/FTO substrate. Although it is easy to fabricate solar cells with this printing

Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite...

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231


/triple cation perovskite (PVK)/Spiro-MeOTAD

active area. (c) Schematic diagram

/triple cation perovskite (PVK)/

Concerning the two-step process, made to create good morphology perovskite films, convolutes control over the change rate of PbI<sup>2</sup> to perovskite [4]. In the recent date, it has been demonstrated that the order of PbI<sup>2</sup> -DMSO-MAI essentially improves the film morphology and quality utilizing a propelled hostile to dissolved designing strategy [6, 88, 89]. Furthermore, this hot-throwing procedure could be exchanged to a considerably less difficult plunge-covering process for large-area film deposition. There is no doubt that the high performance and scale of large area for PSCs are closely related to perovskite film quality. A 1-cm<sup>2</sup> PSC was fabricated for the first time with a modified interface layer and certified efficiency of 15.6% [90]. Then, enhancing the gradient of heterojunction structure for charge separation/transport, its performance increased to 18.21% [91]. Moreover, a vacuum, flash-assisted process has been produced by a solar cell with a 1-cm<sup>2</sup> area [92]. This technique showed a maximum efficiency of 20.5% and a certified efficiency of 19.6% [93].

### **3.5. Spray coating**

Spray coating has been broadly utilized to deposit perovskite films and compact TiO<sup>2</sup> films and is perfect with large-scale, high-throughput manufacture (**Figure 6b**). The first thinking about how to get perovskite films via spray coating came from polymer solar cell fabrication. An ultrasonic spray-coating technique has been discovered in ambient conditions. In this method a system from DMF or DMSO and perovskite materials was investigated with deposition parameter to achieve higher coverage of perovskite films. The parameters which are related to spray coating require to form high-coverage perovskite film such as drying time, substrate temperature, solvent volatility, and post-annealing conditions. With PCE of 11% and an active area about 0.025 cm<sup>2</sup> , the potential of spray coating in fabricating perovskite solar cells has been indicated [94]. A similar work was performed in TiO<sup>2</sup> and achieved PCE of 13% with 0.065 cm<sup>2</sup> active area on a glass/ITO substrate. In low-temperature PET/ITO substrate in the presence of TiO<sup>2</sup> , an efficiency of 8.1% was attained on a flexible device, which is comparable with rollto-roll processing [95]. This spray-coating process is suitable for various perovskite precursor solutions; for example, spray-coating deposition of large bandgap CsPbIBr<sup>2</sup> thin films has a potential for tandem structure devices [96]. Other mixed cations and halide perovskites, FA1 x Cs<sup>x</sup> PbI<sup>3</sup> mixed cation films, were prepared with a spray-assisted solution process [97]. Solar cell devices based on this mixed cation film showed enhanced stability and performance compared with those based on the other. Efficiency increased from 11.3 to 14.2% for the mixed cations.

### **3.6. Screen printing**

Screen printing is a technique used to fabricate PSCs that could be easily fabricated with a printing process. The layer-by-layer printing process starts with screen printing of TiO<sup>2</sup> , followed by printing of ZrO<sup>2</sup> and carbon electrodes. Then, perovskite solution is dropped onto the porous carbon electrode so that it infiltrates into the mesoporous TiO<sup>2</sup> and ZrO<sup>2</sup> . Herein, ZrO<sup>2</sup> functions as a porous insulating layer to prevent direct contact between the carbon electrode and the TiO<sup>2</sup> /FTO substrate. Although it is easy to fabricate solar cells with this printing technique, infiltration of perovskite precursor solution remains a challenge and is the main reason for lower efficiency compared to those of devices fabricated by other means. The most intriguing properties of this carbon-coated, printed, mesoscopic device are the high stability and outstanding outdoor performance. A certified PCE of 12.8% and stable performance over >1000 h in ambient air under full sunlight has been recorded for a device with an active area of 0.28 cm<sup>2</sup> [26, 99]. **Figure 7** shows noncontact inkjet printing offering rapid and digital deposition combined with excellent control over the layer formation for printed perovskite solar cells. Mathies et al. [100] reported that inkjet printing is used to deposit triple cation perovskite layers with 10% cesium in a mixed formamidinium/methylammonium lead iodide/bromide composite for solar cells with high temperature and moisture stability. A reliable process control over a wide range of perovskite layer thickness from 175 to 780 nm and corresponding grain sizes is achieved by adjusting the drop spacing of the inkjet printer cartridge. A continuous power output at constant voltage, resulting in a power conversion efficiency of 12.9%, is demonstrated, representing a major improvement from previously reported inkjetprinted methylammonium lead triiodide perovskite solar cells [100]. Compared with solution processes, dry deposition processes may be more environmentally friendly, as they do not require toxic solvents (DMSO, GBL, DMF, chloroform, chlorobenzene, isopropanol, toluene, diethyl ether, etc.), and they are compatible with high-quality, large-area perovskite film deposition, such as vacuum deposition and chemical vapor deposition.

of perovskite could be finely tuned, and perovskite films with altogether enhanced quality can be set up for superior power conversion [85, 86]. For instance, by utilizing lead acetic acid as lead source, the crystallization of perovskite is significantly speedier. The free pinhole

Concerning the two-step process, made to create good morphology perovskite films, convo-

quality utilizing a propelled hostile to dissolved designing strategy [6, 88, 89]. Furthermore, this hot-throwing procedure could be exchanged to a considerably less difficult plunge-covering process for large-area film deposition. There is no doubt that the high performance and

ricated for the first time with a modified interface layer and certified efficiency of 15.6% [90]. Then, enhancing the gradient of heterojunction structure for charge separation/transport, its performance increased to 18.21% [91]. Moreover, a vacuum, flash-assisted process has been

is perfect with large-scale, high-throughput manufacture (**Figure 6b**). The first thinking about how to get perovskite films via spray coating came from polymer solar cell fabrication. An ultrasonic spray-coating technique has been discovered in ambient conditions. In this method a system from DMF or DMSO and perovskite materials was investigated with deposition parameter to achieve higher coverage of perovskite films. The parameters which are related to spray coating require to form high-coverage perovskite film such as drying time, substrate temperature, solvent volatility, and post-annealing conditions. With PCE of 11% and an active

active area on a glass/ITO substrate. In low-temperature PET/ITO substrate in the presence of

potential for tandem structure devices [96]. Other mixed cations and halide perovskites, FA1-

devices based on this mixed cation film showed enhanced stability and performance compared with those based on the other. Efficiency increased from 11.3 to 14.2% for the mixed cations.

Screen printing is a technique used to fabricate PSCs that could be easily fabricated with a printing process. The layer-by-layer printing process starts with screen printing of TiO<sup>2</sup>

solutions; for example, spray-coating deposition of large bandgap CsPbIBr<sup>2</sup>

the porous carbon electrode so that it infiltrates into the mesoporous TiO<sup>2</sup>

, an efficiency of 8.1% was attained on a flexible device, which is comparable with rollto-roll processing [95]. This spray-coating process is suitable for various perovskite precursor

mixed cation films, were prepared with a spray-assisted solution process [97]. Solar cell

, the potential of spray coating in fabricating perovskite solar cells has been

and carbon electrodes. Then, perovskite solution is dropped onto

scale of large area for PSCs are closely related to perovskite film quality. A 1-cm<sup>2</sup>

Spray coating has been broadly utilized to deposit perovskite films and compact TiO<sup>2</sup>

to perovskite [4]. In the recent date, it has been dem-

area [92]. This technique showed a maximum efficiency

and achieved PCE of 13% with 0.065 cm<sup>2</sup>

PSC was fab-

films and

thin films has a

and ZrO<sup>2</sup>

, fol-

. Herein,


perovskite films were shown by a basic one-step coating process [87].

lutes control over the change rate of PbI<sup>2</sup>

produced by a solar cell with a 1-cm<sup>2</sup>

**3.5. Spray coating**

area about 0.025 cm<sup>2</sup>

**3.6. Screen printing**

lowed by printing of ZrO<sup>2</sup>

TiO<sup>2</sup>

x Cs<sup>x</sup> PbI<sup>3</sup>

of 20.5% and a certified efficiency of 19.6% [93].

indicated [94]. A similar work was performed in TiO<sup>2</sup>

onstrated that the order of PbI<sup>2</sup>

230 Coatings and Thin-Film Technologies

**Figure 7.** (a) 520 nm-thick inkjet-printed perovskite layer on FTO/TiO<sup>2</sup> -coated glass substrate. (b) Photograph of inkjetprinted perovskite solar cells. The substrate contains eight cells with each 3×3 mm<sup>2</sup> active area. (c) Schematic diagram of the solar cell stack, denoting with the different layers: glass/FTO/TiO<sup>2</sup> /triple cation perovskite (PVK)/Spiro-MeOTAD (HTL)/Au [100]. (d) Real solar cell achieved 12% with the configuration: Glass/FTO/TiO<sup>2</sup> /triple cation perovskite (PVK)/ Spiro-MeOTAD (HTL)/Au.

### **4. Summary and future outlook**

The effectivity of photon capture in PSCs has resulted in tangible action and contributed to scientific community [101]. These achievements have an economic impact for future endeavor. Therefore, the innovation in PSC field required a large amount of effort and attention to be reliable and highly efficient at converting sunlight to electricity. Furthermore, the improvement of device engineering methods is urgent. In particular, investigation of photophysical mechanism of the materials also plays an important role. The continuous investigation on current density and voltage characteristics of PSCs would provide a good understanding point for the semiconducting behavior [102]. The improvement in PSC efficiency relies on deposition techniques and material composition [103]. The solution processed in PSCs is more important [104]. We are noticed that the one-step spin coating is broadly the used method because of its simplicity and low cost. The films synthesized by this method have a poor morphology and incomplete coverage, for instance, in the case of planar architecture [84, 85, 105]. On the other hand, in the two-step coating, a layer of lead halide was deposited by spin coating then followed by immerging in organic salt solution and the perovskite films formed by a chemical reaction [51, 52]. The high reaction rates of perovskite materials are important to optimize the coating conditions with sufficient reproducibility. To record high efficiencies by solution processing, it is revealed that the reaction kinetics are required to control and maintain consistent device to minimize batch-to-batch variations. The different vapor-based methods to deposit perovskite films are also discussed, which in many cases show properties different from their counterparts prepared by solution-based methods.

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Organometal Halide Perovskites Thin Film and Their Impact on the Efficiency of Perovskite...

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### **Conflict of interest**

The authors declare no conflict of interest.

### **Author details**

Ahmed Mourtada Elseman

Address all correspondence to: ahmedMourtada5555@yahoo.com

Electronic and Magnetic Materials Department, Advanced Materials Division, Central Metallurgical Research and Development Institute (CMRDI), Helwan, Cairo, Egypt

### **References**

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**4. Summary and future outlook**

232 Coatings and Thin-Film Technologies

**Conflict of interest**

**Author details**

**References**

Ahmed Mourtada Elseman

The authors declare no conflict of interest.

Address all correspondence to: ahmedMourtada5555@yahoo.com

2009;**131**(17):6050-6051. DOI: 10.1021/ja809598r

Electronic and Magnetic Materials Department, Advanced Materials Division, Central Metallurgical Research and Development Institute (CMRDI), Helwan, Cairo, Egypt

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[71] Eperon GE, Burlakov VM, Docampo P, Goriely A, Snaith HJ. Morphological control for high performance, solution-processed planar heterojunction perovskite solar cells.

[72] Bi D, Moon S-J, Häggman L, et al. Using a two-step deposition technique to prepare

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transport layers. Nature Photonics. 2014;**8**(2):128

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perovskite/PCBM planar-heterojunction hybrid solar cells. Advanced Materials. 2014;

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[75] Yella A, Heiniger L-P, Gao P, Nazeeruddin MK, Grätzel M. Nanocrystalline rutile electron extraction layer enables low-temperature solution processed perovskite photovol-

[76] Ma Y, Zheng L, Chung Y-H, et al. A highly efficient mesoscopic solar cell based on

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[80] Chen CW, Kang HW, Hsiao SY, Yang PF, Chiang KM, Lin HW. Efficient and uniform planar-type perovskite solar cells by simple sequential vacuum deposition. Advanced

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PbBr<sup>3</sup>

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238 Coatings and Thin-Film Technologies

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**Chapter 12**

**Provisional chapter**

**Textured BST Thin Film on Silicon Substrate:**

**Textured BST Thin Film on Silicon Substrate:** 

**Tunable Devices**

and Guifu Ding

**Abstract**

**1. Introduction**

**Tunable Devices**

Conchun Zhang, Jianze Huang, Chunsheng Yang and Guifu Ding

http://dx.doi.org/10.5772/intechopen.79270

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

The dielectric properties of Ba0.5Sr0.5TiO3

dielectric properties, tunability

lographic orientation of the films. LaNiO<sup>3</sup>

Conchun Zhang, Jianze Huang, Chunsheng Yang

**Preparation and Its Applications for High Frequency**

**Preparation and Its Applications for High Frequency** 

on Si substrate before BST. The effect of buffer layer such as LNO, MgO and MgO/LNO bilayer on the microstructure and dielectric properties of BST were extensively investigated. The preferred (100) orientation of LNO by radio frequency (RF) magnetron sputtering was dominated by the substrate temperature and highly (100)-oriented LNO thin films were grown on Si substrates at 300°C. The oriented (100) growth of sputtered BST thin films was strongly affected by the orientation of LNO thin films and the tunability of BST thin film was greatly improved with the insertion of (100)-textured LNO buffer layer. In addition, MgO, as a buffer layer, was deposited by RF magnetron sputtering. The results show that the crystallization of BST was also enhanced by the insertion of MgO buffer layer, which enhances the oriental growth of BST along (100). Also, the tunability of the BST thin films was improved and the dielectric loss greatly decreased. Finally, CPW with BST/MgO multilayer was fabricated and the scattering (S) parameters were tested.

Tunable high frequency devices are key components for the next generation of communications and radar systems. Oxides with the perovskite structure, such as barium strontium

**Keywords:** magnetron sputtering, BST, oriented growth, buffer layer,

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

© 2018 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use,

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

(BST) thin films are sensitive to the relative crystal-

(LNO) and MgO were deposited as buffer layer

DOI: 10.5772/intechopen.79270


#### **Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency Tunable Devices Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency Tunable Devices**

DOI: 10.5772/intechopen.79270

Conchun Zhang, Jianze Huang, Chunsheng Yang and Guifu Ding Conchun Zhang, Jianze Huang, Chunsheng Yang and Guifu Ding

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

http://dx.doi.org/10.5772/intechopen.79270

#### **Abstract**

[99] Li X, Tschumi M, Han H, et al. Outdoor performance and stability under elevated temperatures and long-term light soaking of triple-layer mesoporous perovskite photovol-

[100] Mathies F, Eggers H, Richards BS, Hernandez-Sosa G, Lemmer U, Paetzold UW. Inkjetprinted triple cation perovskite solar cells. ACS Applied Energy Materials. 2018;**1**(5):

[101] Park N-G. Perovskite solar cells: an emerging photovoltaic technology. Materials

[102] Meloni S, Moehl T, Tress W, et al. Ionic polarization-induced current-voltage hysteresis

[103] Back H, Kim J, Kim G, et al. Interfacial modification of hole transport layers for efficient large-area perovskite solar cell achieved via blade-coating. Solar Energy Materials and

Solar Cells. 2016;**144**(Supplement C):309-315. DOI: 10.1016/j.solmat.2015.09.018 [104] Kaiyu Y, Fushan L, Jianhua Z, Chandrasekar Perumal V, Tailiang G. All-solution processed semi-transparent perovskite solar cells with silver nanowires electrode. Nano-

[105] Conings B, Baeten L, De Dobbelaere C, D'Haen J, Manca J, Boyen HG. Perovskite-based hybrid solar cells exceeding 10% efficiency with high reproducibility using a thin film

sandwich approach. Advanced Materials. 2014;**26**(13):2041-2046

perovskite solar cells. Nature Communications. Feb 8 2016;**7**:10334.

taics. Energy Technology. 2015;**3**(6):551-555

1834-1839. DOI: 10.1021/acsaem.8b00222

in CH<sup>3</sup>

240 Coatings and Thin-Film Technologies

NH<sup>3</sup> PbX<sup>3</sup>

DOI: 10.1038/ncomms10334

technology. 2016;**27**(9):095202

Today. 2015;**18**(2):65-72. DOI: 10.1016/j.mattod.2014.07.007

The dielectric properties of Ba0.5Sr0.5TiO3 (BST) thin films are sensitive to the relative crystallographic orientation of the films. LaNiO<sup>3</sup> (LNO) and MgO were deposited as buffer layer on Si substrate before BST. The effect of buffer layer such as LNO, MgO and MgO/LNO bilayer on the microstructure and dielectric properties of BST were extensively investigated. The preferred (100) orientation of LNO by radio frequency (RF) magnetron sputtering was dominated by the substrate temperature and highly (100)-oriented LNO thin films were grown on Si substrates at 300°C. The oriented (100) growth of sputtered BST thin films was strongly affected by the orientation of LNO thin films and the tunability of BST thin film was greatly improved with the insertion of (100)-textured LNO buffer layer. In addition, MgO, as a buffer layer, was deposited by RF magnetron sputtering. The results show that the crystallization of BST was also enhanced by the insertion of MgO buffer layer, which enhances the oriental growth of BST along (100). Also, the tunability of the BST thin films was improved and the dielectric loss greatly decreased. Finally, CPW with BST/MgO multilayer was fabricated and the scattering (S) parameters were tested.

**Keywords:** magnetron sputtering, BST, oriented growth, buffer layer, dielectric properties, tunability

### **1. Introduction**

Tunable high frequency devices are key components for the next generation of communications and radar systems. Oxides with the perovskite structure, such as barium strontium

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2018 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

titanate (BST), have ferroelectric, high dielectric, high-Tc superconductive or very large magnetron resistance. Therefore, they are widely investigated for applications in the tunable microwave filters and monolithic microwave integrated circuit decoupling capacitors. For application in frequency agile devices, it is desirable to have as large a capacitance change ratio and as low a dielectric loss as possible. Many researches have been studied to increase the tunability and to reduce the loss of epitaxial BST films grown on single crystal oxide substrates, such as LaAlO<sup>3</sup> and MgO [1–4]. However, in order to integrate BST thin films with Si for frequency agile devices, it is necessary to deposit BST thin films on Si substrates. But it is difficult to prepare high-quality oxide films on Si due to inherent crystallographic incompatibility of the two materials and the lack of an epitaxial template for the growth of well-oriented Ba1−xSr<sup>x</sup> TiO3 [4].

magnetron sputtering. The substrate was silicon or platinized silicon. The sputtering power

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency…

For the LNO deposited at room temperature, the LNO film was post annealed at 700°C for 1 h

MgO buffer layers were deposited at room temperature on Si wafers. The working pres-

is 100 W. The thickness of thin films was measured by a stylus profiler (Veeco Dektak 6M). The thickness of the MgO buffer layer is between 50 and 150 nm. Then, Ba0.5Sr0.5TiO3

films were deposited on MgO buffer layers by RF magnetron sputtering. The deposition pressure was 0.8 Pa. Both the MgO and the BST/MgO composite films were annealed in the O<sup>2</sup>

The compositional analysis was performed by induced coupling plasma emission spectroscopy (ICP) using an Iris Advantage 1000 instrument. The phase and crystallinity of the films were characterized by an X-ray diffractometer (XRD) D/MAX-γB, Rigaku with CuKa radiation. Atomic force microscopy (AFM) surface morphologies were achieved with a Digital Instrument Nanoscope III with AFM tapping mode. Then the resistivity was determined by a four-point probe (China, D41-11A/ZM) at room temperature. The microstructure of the thin films was examined by field emission-scanning electron microscopy (FE-SEM). For the measurements of the dielectric properties, the Cu/Cr top electrodes were deposited by sputtering method and were patterned by liftoff technology. The planar capacitor was prepared based on

BST thin film was deposited under various relative oxygen ratios with total Ar and O<sup>2</sup>

ratio ranged from 10 to 50% with the total pressure of 1.33 Pa.

partial ratio changed was 10%. The RF magnetron sputtering power

/Si thin film, and interdigital capacitor was prepared based on BST/

/Si thin films. The dielectric properties of the BST films were measured using an HP

 **or MgO buffer layer on Si for the textured** 

4194A impedance analyzer at a frequency of 1 MHz. C-V characteristic was carried out with a

BST thin films can be deposited by different methods such as hydrothermal, pulsed laser deposition, metal organic chemical vapor deposition (MOCVD), chemical solution deposition method and sol-gel. Among these methods, we chose radio frequency magnetron sputtering technique due to its industrial process compatibility through ease of implementation, superior compositional reproducibility, medium deposition rate, uniform deposition over large area with additional interest in mass production [22–24]. However, transferring the target stoichiometry to the substrate remains a challenge with on-axis RF-magnetron sputtering due to re-sputtering effects. Thus, the control of the re-sputtering and the etch rate are critical issues related to the sputtering of BST thin films. This etch rate is known to be strongly influenced by

sine wave with a 0.1 V step at a frequency of 1 MHz at room temperature.

pres-

243

thin

atmosphere for crystallization.

http://dx.doi.org/10.5772/intechopen.79270

was 100 W and the relative O<sup>2</sup>

sure is 0.3 Pa, and O<sup>2</sup>

BST/LNO Pt (111)//Ti/SiO<sup>2</sup>

**3. Oriented LaNiO3**

 **thin films**

**3.1. RF sputtered BST thin film on Si substrate**

**Ba0.5Sr0.5TiO3**

atmosphere.

MgO/SiO<sup>2</sup>

atmosphere for crystallization.

sure of 0.8 Pa, and then it was post annealed at 700–750°C in O<sup>2</sup>

in O2

The dielectric properties of BST thin films are sensitive to the relative crystallographic orientation of the films, usually highly oriented BST film leads to higher performance compared to a randomly oriented one. There are various methods for changing the orientation of the ferroelectric films, such as the insertion of a buffer layer and/or different bottom electrodes [4–6]. Kang et al. reported the physical properties of epitaxial Ba0.6Sr0.4TiO3 (BST) films grown on SiO2 /Si using biaxially oriented MgO as template layers [4]. A buffer layer used between the silicon substrate and the perovskite oxides thin film has been shown to be a good way of overcoming the problem [5, 7–12]. Previous study showed that the dielectric constant of BST thin films was improved by LNO buffer layer [8, 11], where the thin films were prepared by pulsed laser deposition. LaNiO<sup>3</sup> (LNO), a perovskite-type metallic oxide, has a lattice parameter of 3.84 Å. The crystal structural and lattice constants of the LNO matches well with ferroelectric thin films such as BST films [5, 10], which offer the benefits of better lattice matching and structural compatibility and the potential for improved dielectric properties. However, some studies showed that the LNO layer hardly affected the texture of BST thin films [13].

On the other hand, MgO shows low microwave loss, good thermal stability and has been increasingly utilized in microwave devices. The crystal parameter of MgO is a = b = c = 0.42 nm, which is close to that of BST (about 0.39 nm). The single crystal MgO has been used as a microwave substrate [14]. However, the high cost of MgO crystal restrains the popularity of massive production [15]. Few works have been reported on MgO thin films deposited by RF magnetron sputtering as a buffer layer for BST [16].

In this chapter, buffer layer such as LaNiO<sup>3</sup> thin film or MgO thin film were deposited on silicon substrate by RF magnetron sputtering. Here, both the effect of the sputtering parameters on the orientation of LNO thin film and the effect of the LNO/MgO buffer layer on the microstructure and dielectric properties of sputtered BST were studied [17–21].

### **2. Experimental details**

BST and LNO targets for our experiment were synthesized by conventional sintering process using a conventional mixing oxides method. The LNO films with varied thickness (600– 2400 Å), mainly used as a buffer layer, were deposited at room temperature or at 300°C by RF magnetron sputtering. The substrate was silicon or platinized silicon. The sputtering power was 100 W and the relative O<sup>2</sup> ratio ranged from 10 to 50% with the total pressure of 1.33 Pa. For the LNO deposited at room temperature, the LNO film was post annealed at 700°C for 1 h in O2 atmosphere for crystallization.

BST thin film was deposited under various relative oxygen ratios with total Ar and O<sup>2</sup> pressure of 0.8 Pa, and then it was post annealed at 700–750°C in O<sup>2</sup> atmosphere for crystallization.

MgO buffer layers were deposited at room temperature on Si wafers. The working pressure is 0.3 Pa, and O<sup>2</sup> partial ratio changed was 10%. The RF magnetron sputtering power is 100 W. The thickness of thin films was measured by a stylus profiler (Veeco Dektak 6M). The thickness of the MgO buffer layer is between 50 and 150 nm. Then, Ba0.5Sr0.5TiO3 thin films were deposited on MgO buffer layers by RF magnetron sputtering. The deposition pressure was 0.8 Pa. Both the MgO and the BST/MgO composite films were annealed in the O<sup>2</sup> atmosphere.

The compositional analysis was performed by induced coupling plasma emission spectroscopy (ICP) using an Iris Advantage 1000 instrument. The phase and crystallinity of the films were characterized by an X-ray diffractometer (XRD) D/MAX-γB, Rigaku with CuKa radiation. Atomic force microscopy (AFM) surface morphologies were achieved with a Digital Instrument Nanoscope III with AFM tapping mode. Then the resistivity was determined by a four-point probe (China, D41-11A/ZM) at room temperature. The microstructure of the thin films was examined by field emission-scanning electron microscopy (FE-SEM). For the measurements of the dielectric properties, the Cu/Cr top electrodes were deposited by sputtering method and were patterned by liftoff technology. The planar capacitor was prepared based on BST/LNO Pt (111)//Ti/SiO<sup>2</sup> /Si thin film, and interdigital capacitor was prepared based on BST/ MgO/SiO<sup>2</sup> /Si thin films. The dielectric properties of the BST films were measured using an HP 4194A impedance analyzer at a frequency of 1 MHz. C-V characteristic was carried out with a sine wave with a 0.1 V step at a frequency of 1 MHz at room temperature.

#### **3. Oriented LaNiO3 or MgO buffer layer on Si for the textured Ba0.5Sr0.5TiO3 thin films**

### **3.1. RF sputtered BST thin film on Si substrate**

titanate (BST), have ferroelectric, high dielectric, high-Tc superconductive or very large magnetron resistance. Therefore, they are widely investigated for applications in the tunable microwave filters and monolithic microwave integrated circuit decoupling capacitors. For application in frequency agile devices, it is desirable to have as large a capacitance change ratio and as low a dielectric loss as possible. Many researches have been studied to increase the tunability and to reduce the loss of epitaxial BST films grown on single crystal oxide

with Si for frequency agile devices, it is necessary to deposit BST thin films on Si substrates. But it is difficult to prepare high-quality oxide films on Si due to inherent crystallographic incompatibility of the two materials and the lack of an epitaxial template for the growth of

The dielectric properties of BST thin films are sensitive to the relative crystallographic orientation of the films, usually highly oriented BST film leads to higher performance compared to a randomly oriented one. There are various methods for changing the orientation of the ferroelectric films, such as the insertion of a buffer layer and/or different bottom electrodes [4–6].

/Si using biaxially oriented MgO as template layers [4]. A buffer layer used between the silicon substrate and the perovskite oxides thin film has been shown to be a good way of overcoming the problem [5, 7–12]. Previous study showed that the dielectric constant of BST thin films was improved by LNO buffer layer [8, 11], where the thin films were prepared by pulsed

3.84 Å. The crystal structural and lattice constants of the LNO matches well with ferroelectric thin films such as BST films [5, 10], which offer the benefits of better lattice matching and structural compatibility and the potential for improved dielectric properties. However, some

On the other hand, MgO shows low microwave loss, good thermal stability and has been increasingly utilized in microwave devices. The crystal parameter of MgO is a = b = c = 0.42 nm, which is close to that of BST (about 0.39 nm). The single crystal MgO has been used as a microwave substrate [14]. However, the high cost of MgO crystal restrains the popularity of massive production [15]. Few works have been reported on MgO thin films deposited by RF

silicon substrate by RF magnetron sputtering. Here, both the effect of the sputtering parameters on the orientation of LNO thin film and the effect of the LNO/MgO buffer layer on the

BST and LNO targets for our experiment were synthesized by conventional sintering process using a conventional mixing oxides method. The LNO films with varied thickness (600– 2400 Å), mainly used as a buffer layer, were deposited at room temperature or at 300°C by RF

microstructure and dielectric properties of sputtered BST were studied [17–21].

studies showed that the LNO layer hardly affected the texture of BST thin films [13].

Kang et al. reported the physical properties of epitaxial Ba0.6Sr0.4TiO3

and MgO [1–4]. However, in order to integrate BST thin films

(LNO), a perovskite-type metallic oxide, has a lattice parameter of

thin film or MgO thin film were deposited on

(BST) films grown on

substrates, such as LaAlO<sup>3</sup>

242 Coatings and Thin-Film Technologies

TiO3 [4].

magnetron sputtering as a buffer layer for BST [16].

In this chapter, buffer layer such as LaNiO<sup>3</sup>

**2. Experimental details**

well-oriented Ba1−xSr<sup>x</sup>

laser deposition. LaNiO<sup>3</sup>

SiO2

BST thin films can be deposited by different methods such as hydrothermal, pulsed laser deposition, metal organic chemical vapor deposition (MOCVD), chemical solution deposition method and sol-gel. Among these methods, we chose radio frequency magnetron sputtering technique due to its industrial process compatibility through ease of implementation, superior compositional reproducibility, medium deposition rate, uniform deposition over large area with additional interest in mass production [22–24]. However, transferring the target stoichiometry to the substrate remains a challenge with on-axis RF-magnetron sputtering due to re-sputtering effects. Thus, the control of the re-sputtering and the etch rate are critical issues related to the sputtering of BST thin films. This etch rate is known to be strongly influenced by the deposition parameters. There have been several reported methods to reduce the impact of this phenomenon. The common method is to use an off-axis deposition [25], where the surface is not parallel to the target surface. Another method is to increase the sputtering gas pressure, providing higher anion-gas collisions to reduce the energy of negative particles or deflecting them away from the film. In the present study, the key to achieve stoichiometric BST thin film deposition was by controlling both target composition and oxygen concentration to avoid the re-sputtering and the substrate etching. BST thin films were deposited by on-axis radiofrequency magnetron sputtering, then the films were annealed at 750°C in O<sup>2</sup> atmosphere.

**Figure 1** shows the composition of the BST thin films deposited by on-axis RF-magnetron sputtering in Ar/O<sup>2</sup> mix gas is deviated from the target, which is due to re-sputtering effects, that is, oxygen negative ions bombardment of the growing film. The BST thin films close to stoichiometric composition, Ba0.5Sr0.5TiO3 , were sputtered with a (Ba0.8Sr0.8)TiO3 target in Ar gas. The re-sputtering effect can be decreased when no oxygen was introduced during sputtering since the re-sputtering phenomenon is attributed to the presence of energetic particles (negative oxygen ions).

As can be seen in **Figure 2**, the BST thin film deposited at room temperature was amorphous. When annealed at 750°C in O<sup>2</sup> atmosphere for 30 min, thin film crystallized. With increasing annealing time, the grain size increased and some large grains appeared. Annealing at 750°C for no more than 30 min is beneficial for the homogenous grain growth.

**Figure 3** shows the XRD patterns of the BST thin films on SiO<sup>2</sup> /Si deposited in different oxygen ratio (PO2 ) and annealed at 750°C for 30 min. BST sputtered in Ar/O<sup>2</sup> mix gas cannot form complete perovskite phase and some non-perovskite phase appears because of oxygen negative ions bombardment of the growing film. But Perovskite phase formed when the BST thin film were sputtered in Ar gas.

**Figure 4** shows the C-V characteristic of the BST thin films directly deposited on Pt/SiO<sup>2</sup> /Si. The curve was asymmetric because the bottom electrode (Pt) and top electrode (Cu) were different. When the bias voltage switched, the two curves were symmetric with x = 0 V as axis of symmetry.

**3.2. RF sputtered LaNiO<sup>3</sup>**

**Figure 2.** SEM of the BST thin films on SiO<sup>2</sup>

atmosphere for 60 min.

(c) annealed at 750 in O2

of O2

the O2

ratio when the O<sup>2</sup>

 **on Si substrate**

**Figure 3.** XRD spectra of BST film sputtered under different relative oxygen ratio (PO2

As shown in **Figure 5**, the chemical composition analysis by ICP, that is, the La/Ni ratio (mole ratio) in the LNO thin films varied with relative oxygen ratios. The La/Ni ratio keeps stoichiometric ratio of 1:1 while the relative oxygen ratio is between 10 and 25%; however, as the relative oxygen ratio exceeds 25%, the content of La decreases. In the sputtering process, oxygen participated in the sputtering as well as a reaction gas. Because of the energy difference between oxygen ion and argon ion, the change of the relative oxygen ratio will lead to the change of sputtering yield of La and Ni, but the change should be continuous and monotonous. The structure with La/Ni ratio of 1:1 is relatively stable, La and Ni are selectively adsorbed in the film in the process of sputtering, so the component of the film keeps stable when the relative oxygen ratio maintains between 10 and 25%. This indicates that the films tend to grow according to stoichiometric ratio during the sputtering process [26]. **Figure 6** shows the XRD patterns of the LNO films. According to XRD results of **Figure 6a**, the LNO films deposited at room temperature with different relative oxygen ratio and post annealed at 700°C for 1 h show a perovskite-type cubic structure with a random orientation irrespective

).

/Si: (a) as-deposited (b) annealed at 750 in O<sup>2</sup>

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency…

atmosphere for 30 min and

http://dx.doi.org/10.5772/intechopen.79270

245

ratio. The full width at half maximum (FWHM) of the (110) peak increases slightly when

 ratio exceeds 25%, that is, the LNO films have better crystalline when the relative oxygen ratio is no more than 25%. The mole ratio of La/Ni for the film deviates from stoichiometric

ratio exceeds 25%, therefore, the LNO films deposited above 25% O<sup>2</sup>

ratio

**Figure 1.** (a) The effect of target composition and (b) PO2 on (Ba + Sr)/Ti ratio of the BST thin films SiO<sup>2</sup> /Si (total pressure = 0.8 Pa).

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency… http://dx.doi.org/10.5772/intechopen.79270 245

**Figure 2.** SEM of the BST thin films on SiO<sup>2</sup> /Si: (a) as-deposited (b) annealed at 750 in O<sup>2</sup> atmosphere for 30 min and (c) annealed at 750 in O2 atmosphere for 60 min.

**Figure 3.** XRD spectra of BST film sputtered under different relative oxygen ratio (PO2 ).

#### **3.2. RF sputtered LaNiO<sup>3</sup> on Si substrate**

the deposition parameters. There have been several reported methods to reduce the impact of this phenomenon. The common method is to use an off-axis deposition [25], where the surface is not parallel to the target surface. Another method is to increase the sputtering gas pressure, providing higher anion-gas collisions to reduce the energy of negative particles or deflecting them away from the film. In the present study, the key to achieve stoichiometric BST thin film deposition was by controlling both target composition and oxygen concentration to avoid the re-sputtering and the substrate etching. BST thin films were deposited by on-axis radio-

**Figure 1** shows the composition of the BST thin films deposited by on-axis RF-magnetron

that is, oxygen negative ions bombardment of the growing film. The BST thin films close to

gas. The re-sputtering effect can be decreased when no oxygen was introduced during sputtering since the re-sputtering phenomenon is attributed to the presence of energetic particles

As can be seen in **Figure 2**, the BST thin film deposited at room temperature was amorphous.

annealing time, the grain size increased and some large grains appeared. Annealing at 750°C

complete perovskite phase and some non-perovskite phase appears because of oxygen negative ions bombardment of the growing film. But Perovskite phase formed when the BST thin

**Figure 4** shows the C-V characteristic of the BST thin films directly deposited on Pt/SiO<sup>2</sup>

The curve was asymmetric because the bottom electrode (Pt) and top electrode (Cu) were different. When the bias voltage switched, the two curves were symmetric with x = 0 V as axis

) and annealed at 750°C for 30 min. BST sputtered in Ar/O<sup>2</sup>

mix gas is deviated from the target, which is due to re-sputtering effects,

, were sputtered with a (Ba0.8Sr0.8)TiO3

atmosphere for 30 min, thin film crystallized. With increasing

on (Ba + Sr)/Ti ratio of the BST thin films SiO<sup>2</sup>

atmosphere.

/Si deposited in different oxy-

mix gas cannot form

/Si.

/Si (total

target in Ar

frequency magnetron sputtering, then the films were annealed at 750°C in O<sup>2</sup>

for no more than 30 min is beneficial for the homogenous grain growth.

**Figure 3** shows the XRD patterns of the BST thin films on SiO<sup>2</sup>

sputtering in Ar/O<sup>2</sup>

244 Coatings and Thin-Film Technologies

(negative oxygen ions).

gen ratio (PO2

of symmetry.

pressure = 0.8 Pa).

When annealed at 750°C in O<sup>2</sup>

film were sputtered in Ar gas.

**Figure 1.** (a) The effect of target composition and (b) PO2

stoichiometric composition, Ba0.5Sr0.5TiO3

As shown in **Figure 5**, the chemical composition analysis by ICP, that is, the La/Ni ratio (mole ratio) in the LNO thin films varied with relative oxygen ratios. The La/Ni ratio keeps stoichiometric ratio of 1:1 while the relative oxygen ratio is between 10 and 25%; however, as the relative oxygen ratio exceeds 25%, the content of La decreases. In the sputtering process, oxygen participated in the sputtering as well as a reaction gas. Because of the energy difference between oxygen ion and argon ion, the change of the relative oxygen ratio will lead to the change of sputtering yield of La and Ni, but the change should be continuous and monotonous. The structure with La/Ni ratio of 1:1 is relatively stable, La and Ni are selectively adsorbed in the film in the process of sputtering, so the component of the film keeps stable when the relative oxygen ratio maintains between 10 and 25%. This indicates that the films tend to grow according to stoichiometric ratio during the sputtering process [26]. **Figure 6** shows the XRD patterns of the LNO films. According to XRD results of **Figure 6a**, the LNO films deposited at room temperature with different relative oxygen ratio and post annealed at 700°C for 1 h show a perovskite-type cubic structure with a random orientation irrespective of O2 ratio. The full width at half maximum (FWHM) of the (110) peak increases slightly when the O2 ratio exceeds 25%, that is, the LNO films have better crystalline when the relative oxygen ratio is no more than 25%. The mole ratio of La/Ni for the film deviates from stoichiometric ratio when the O<sup>2</sup> ratio exceeds 25%, therefore, the LNO films deposited above 25% O<sup>2</sup> ratio

**Figure 4.** C-V diagram before and after changing the DC bias direction.

oriented growth of the LNO thin films was dominated by deposition temperature, which was due to the collision between the energetic particles and the thin film during sputtering [12],

/Si with 10% oxygen ratio at 300°C.

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency…

 **(nm) Surface morphology**

/Si at room temperature with different oxygen ratios

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247

/Si at room temperature and annealed at 700°C for 1 h with

/Si at different oxygen ratios and postannealed at 700°C.

As we know, the preferred orientations of thin films are affected by many parameters. Surface free energies of two-dimensional planes affect the preferred orientation of films in the initial stage of film growth [28]. The film will be oriented to the plane which has the smallest surface free energy if the effect of the substrate is negligible. In ceramic systems with unit cell containing cations and anions, both the electrostatic charges of two-dimensional planes and surface packing densities should be considered in the calculation of surface free energies. It is difficult to find out the surface free energies of ceramics. Furthermore, in RF magnetron sputtering process ionic species ejected from the target surface will also influence the orientations of films. The film orientation can be adjusted by changing the process conditions such as substrate temperature, ambient pressure. In the film deposited at low pressures, the size of the plume was large and the substrate was located at the middle of plume. The kinetic energies of ionic species were higher than those of atomic and molecular species. While the absolute number of ionic species was small, they were absorbed on the substrate preferentially and the influence of electrostatic charges became dominant at the initial stage of film growth. The

the detailed discussion was given in Ref [27].

**Table 1.** The morphology of the LNO films deposited on SiO<sup>2</sup>

**Figure 6.** XRD patterns of the LNO thin films. (a) Deposited on SiO<sup>2</sup>

10 1.5 Smooth 25 2.4 Smooth

33 4.6 Rough, some voids

50 5.0 Many voids

and post annealed at 700°C. (b) Deposited on Pt(111)/SiO<sup>2</sup>

10% oxygen ratio. (c) Deposited on Pt (111)/SiO<sup>2</sup>

**Oxygen ratios (%) Ra**

**Figure 5.** The dependence of La/Ni ratio on the relative oxygen ratio.

show bad crystalline. The surface morphology of the LNO thin films by AFM are summarized in **Table 1**. The mean roughness value of the film deposited with 10% O<sup>2</sup> ratio was 1.5 nm, while the surface roughness value was 5.0 nm and many voids were observed when the O<sup>2</sup> ratio exceeded 25%. **Figure 6(b)** and **(c)** shows the thin films deposited on Pt (111)/SiO<sup>2</sup> /Si with 10% oxygen at room temperature and 300°C, respectively. It can be seen that the films show an entirely perovskite phase with a (100) preferential orientation when the LNO thin film was deposited at 300°C (**Figure 6(c)**). However, when the LNO thin film was deposited at room temperature and was post-annealed for crystallization (**Figure 6(b)**), the thin film showed random orientation. Combined the XRD results in **Figure 6(a)–(c)**, it revealed that the Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency… http://dx.doi.org/10.5772/intechopen.79270 247

**Figure 6.** XRD patterns of the LNO thin films. (a) Deposited on SiO<sup>2</sup> /Si at room temperature with different oxygen ratios and post annealed at 700°C. (b) Deposited on Pt(111)/SiO<sup>2</sup> /Si at room temperature and annealed at 700°C for 1 h with 10% oxygen ratio. (c) Deposited on Pt (111)/SiO<sup>2</sup> /Si with 10% oxygen ratio at 300°C.


**Table 1.** The morphology of the LNO films deposited on SiO<sup>2</sup> /Si at different oxygen ratios and postannealed at 700°C.

oriented growth of the LNO thin films was dominated by deposition temperature, which was due to the collision between the energetic particles and the thin film during sputtering [12], the detailed discussion was given in Ref [27].

As we know, the preferred orientations of thin films are affected by many parameters. Surface free energies of two-dimensional planes affect the preferred orientation of films in the initial stage of film growth [28]. The film will be oriented to the plane which has the smallest surface free energy if the effect of the substrate is negligible. In ceramic systems with unit cell containing cations and anions, both the electrostatic charges of two-dimensional planes and surface packing densities should be considered in the calculation of surface free energies. It is difficult to find out the surface free energies of ceramics. Furthermore, in RF magnetron sputtering process ionic species ejected from the target surface will also influence the orientations of films. The film orientation can be adjusted by changing the process conditions such as substrate temperature, ambient pressure. In the film deposited at low pressures, the size of the plume was large and the substrate was located at the middle of plume. The kinetic energies of ionic species were higher than those of atomic and molecular species. While the absolute number of ionic species was small, they were absorbed on the substrate preferentially and the influence of electrostatic charges became dominant at the initial stage of film growth. The

show bad crystalline. The surface morphology of the LNO thin films by AFM are summarized

while the surface roughness value was 5.0 nm and many voids were observed when the O<sup>2</sup> ratio exceeded 25%. **Figure 6(b)** and **(c)** shows the thin films deposited on Pt (111)/SiO<sup>2</sup>

with 10% oxygen at room temperature and 300°C, respectively. It can be seen that the films show an entirely perovskite phase with a (100) preferential orientation when the LNO thin film was deposited at 300°C (**Figure 6(c)**). However, when the LNO thin film was deposited at room temperature and was post-annealed for crystallization (**Figure 6(b)**), the thin film showed random orientation. Combined the XRD results in **Figure 6(a)–(c)**, it revealed that the

ratio was 1.5 nm,

/Si

in **Table 1**. The mean roughness value of the film deposited with 10% O<sup>2</sup>

**Figure 5.** The dependence of La/Ni ratio on the relative oxygen ratio.

**Figure 4.** C-V diagram before and after changing the DC bias direction.

246 Coatings and Thin-Film Technologies

kinetic energies of absorbed species were determined by the substrate temperature. When the substrate temperature was high enough, the absorbed species had sufficient energies to rearrange along the plane with the electrically neutral planes, that is, the (100) equivalent planes, therefore textured LNO films were obtained.

The resistivity of LNO measured by four points probe method is shown in **Table 2**. It shows that the resistivity increases rapidly when the O<sup>2</sup> ratio is more than 33%, which can be explained as following: one is the decreased La resulting in the composition deviated from the stoichiometric ratio, another one is that the film contains many voids. This indicates that the films can be deposited at 300°C with O<sup>2</sup> ratio less than 33%. The effects of the relative O<sup>2</sup> ratio, substrate temperatures on the microstructure and electrical properties of the LaNiO<sup>3</sup> thin films by RF sputtering have been investigated [26]. It revealed the films deposited at 300°C show (100) preferred orientation. This indicates that the substrate temperature plays an important role in the determination of the films orientation. The LNO films deposited with 10% O<sup>2</sup> ratio had the lowest resistivity and as such be suitable as a buffer layer or electrode of perovskite oxide thin films.

#### **3.3. Textured BST thin film on LNO/Pt/SiO<sup>2</sup> /Si (100) substrate**

BST thin films were deposited on different electrodes and post annealed in O<sup>2</sup> at 700–750 °C atmosphere for 30 min. **Figure 7(a)–(c)** shows the XRD patterns of BST thin films deposited on different electrodes. The films in **Figure 7b** and **c** show a well-developed perovskite phase without other crystalline phases, but the BST thin films directly deposited on Pt electrode showed relatively weak crystallization compared with those deposited on LNO buffer layer. Mainly oriented (100) peak with small extra (110) peaks, was observed in the BST film deposited on LNO (100)/Pt (111) in **Figure 8b** and **c** [10]. It can be seen that the (100)-textured LNO films due to a good match of lattice parameters between LNO and BST, as well as a similar perovskite structures, further facilitated the crystallization and growth of the BST films. Also, the BST thin film deposited on 120 nm LNO shows better crystallinity than that of the BST film on 60 nm LNO. The AFM images in **Figure 7(d)–(f)** revealed that the surface of BST on Pt shows some white hillocks because Ti can migrate into the surface via Pt grain boundaries and results in the formation of hillocks, and larger grain size is observed in the BST thin films deposited on LNO/Pt (111) compared to that deposited on Pt (111). These results are in agreement with XRD, which means that LNO buffer layers enhance the growth of BST grains. On

the other hand, increasing LNO thickness leads to the rougher surface when the thickness of LNO is more than 60 nm. The root-mean-square (RMS) values of surface roughness of the BST

**Figure 7.** Characterization of BST thin films [30]. (a–c) X-ray diffraction patterns of BST films deposited on various substrates (a) Pt (111), (b) LNO (100) 120 nm/Pt (111), (c) LNO (100) 60 nm/Pt (111); (d–f) AFM surface morphology of the BST thin films deposited on various substrates (d) Pt (111), (e) LNO (100) 120 nm/Pt (111), (f) LNO (100) 60 nm/Pt (111).

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249

For the measurements of the dielectric properties, parallel capacitor was prepared. **Figure 8(1)** shows the structure of BST capacitor and **Figure 8(2)** shows the dielectric properties of the BST thin films varied with the electric field. The asymmetric C-V curves may arise from the difference between bottom electrode and top electrode [29]. The tunability of the ferroelectric

According to Eq. (1), the tunability of BST thin films on the LNO (100)/Pt (111) and on the LNO (110)/Pt (111) were about 63 and 50% at 500 kV/cm, which were higher than that (about 30%) of BST thin film directly deposited on the Pt (111). The value is comparable with that

*Cmax*

(1)

film in **Figure 7(d)–f)** are 7, 5, 8 nm, respectively.

**3.4. The dielectric properties of BST multilayer thin films**

film can be expressed with the equation as following:

Tunability <sup>=</sup> *Cmax* <sup>−</sup> *<sup>C</sup>* \_\_\_\_\_\_\_\_*min*


**Table 2.** The resistivity of the LNO thin film sputtered with different O<sup>2</sup> ratios.

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency… http://dx.doi.org/10.5772/intechopen.79270 249

**Figure 7.** Characterization of BST thin films [30]. (a–c) X-ray diffraction patterns of BST films deposited on various substrates (a) Pt (111), (b) LNO (100) 120 nm/Pt (111), (c) LNO (100) 60 nm/Pt (111); (d–f) AFM surface morphology of the BST thin films deposited on various substrates (d) Pt (111), (e) LNO (100) 120 nm/Pt (111), (f) LNO (100) 60 nm/Pt (111).

the other hand, increasing LNO thickness leads to the rougher surface when the thickness of LNO is more than 60 nm. The root-mean-square (RMS) values of surface roughness of the BST film in **Figure 7(d)–f)** are 7, 5, 8 nm, respectively.

#### **3.4. The dielectric properties of BST multilayer thin films**

kinetic energies of absorbed species were determined by the substrate temperature. When the substrate temperature was high enough, the absorbed species had sufficient energies to rearrange along the plane with the electrically neutral planes, that is, the (100) equivalent planes,

The resistivity of LNO measured by four points probe method is shown in **Table 2**. It shows

explained as following: one is the decreased La resulting in the composition deviated from the stoichiometric ratio, another one is that the film contains many voids. This indicates that

ratio, substrate temperatures on the microstructure and electrical properties of the LaNiO<sup>3</sup> thin films by RF sputtering have been investigated [26]. It revealed the films deposited at 300°C show (100) preferred orientation. This indicates that the substrate temperature plays an important role in the determination of the films orientation. The LNO films deposited with

atmosphere for 30 min. **Figure 7(a)–(c)** shows the XRD patterns of BST thin films deposited on different electrodes. The films in **Figure 7b** and **c** show a well-developed perovskite phase without other crystalline phases, but the BST thin films directly deposited on Pt electrode showed relatively weak crystallization compared with those deposited on LNO buffer layer. Mainly oriented (100) peak with small extra (110) peaks, was observed in the BST film deposited on LNO (100)/Pt (111) in **Figure 8b** and **c** [10]. It can be seen that the (100)-textured LNO films due to a good match of lattice parameters between LNO and BST, as well as a similar perovskite structures, further facilitated the crystallization and growth of the BST films. Also, the BST thin film deposited on 120 nm LNO shows better crystallinity than that of the BST film on 60 nm LNO. The AFM images in **Figure 7(d)–(f)** revealed that the surface of BST on Pt shows some white hillocks because Ti can migrate into the surface via Pt grain boundaries and results in the formation of hillocks, and larger grain size is observed in the BST thin films deposited on LNO/Pt (111) compared to that deposited on Pt (111). These results are in agreement with XRD, which means that LNO buffer layers enhance the growth of BST grains. On

BST thin films were deposited on different electrodes and post annealed in O<sup>2</sup>

**Oxygen ratio (%) Resistivity(MΩ cm) (Td = 300°C)**

10 2.45 20 2.51 33 2.49 50 23.1

**Table 2.** The resistivity of the LNO thin film sputtered with different O<sup>2</sup>

ratio had the lowest resistivity and as such be suitable as a buffer layer or electrode of

**/Si (100) substrate**

ratios.

ratio is more than 33%, which can be

at 700–750 °C

ratio less than 33%. The effects of the relative O<sup>2</sup>

therefore textured LNO films were obtained.

the films can be deposited at 300°C with O<sup>2</sup>

**3.3. Textured BST thin film on LNO/Pt/SiO<sup>2</sup>**

10% O<sup>2</sup>

perovskite oxide thin films.

248 Coatings and Thin-Film Technologies

that the resistivity increases rapidly when the O<sup>2</sup>

For the measurements of the dielectric properties, parallel capacitor was prepared. **Figure 8(1)** shows the structure of BST capacitor and **Figure 8(2)** shows the dielectric properties of the BST thin films varied with the electric field. The asymmetric C-V curves may arise from the difference between bottom electrode and top electrode [29]. The tunability of the ferroelectric film can be expressed with the equation as following:

$$\text{Tunability} = \frac{C\_{\text{max}} - C\_{\text{min}}}{C\_{\text{max}}} \tag{1}$$

According to Eq. (1), the tunability of BST thin films on the LNO (100)/Pt (111) and on the LNO (110)/Pt (111) were about 63 and 50% at 500 kV/cm, which were higher than that (about 30%) of BST thin film directly deposited on the Pt (111). The value is comparable with that

layer kept 120 nm, the tunability of BST with 100–125 nm is better than that with thicker BST. **Figure 9(b)** shows the DC bias voltage dependence of dielectric loss for the thin film

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency…

has the lowest dielectric loss, which demonstrated the LNO buffer layer help in reducing the dielectric loss. However, when the BST/LNO film was annealed at 750°C, the dielectric loss increased, and it increased rapidly with increasing bias voltage, which may be due to two reasons: one is that the Ti diffusion toward BST and reduced the effective dielectric layer, on the other hand, the lattice parameter of LNO is 3.84 and that of BST is 3.94, mismatch between BST and LNO is 2.6% the dislocations increased with increasing annealing temperature.

The effect of buffer layer thickness on the dielectric properties of the BST thin films was

**Figure 10(a)** shows the DC bias voltage dependence of dielectric constant for the BST deposited on LNO buffer layer with varied thickness δ. The commonly used figure of merit (FOM) for electrically tuned device applications is the ratio of the tunability to dielectric loss (tgδ),

**Table 3** shows the tunability and FOM of BST with different thickness of LNO buffer layer at applied electric field of 400 kV/cm. The tunability of BST thin film on the LNO (100)/Pt (111) was about 63% at a vertical applied field of 400 kV/cm, and it was higher than that (about 32%) of BST thin film on the Pt (111). **Figure 10(b)** reveals that the dielectric loss for the BST on 120 nm LNO was the smallest when the DC bias is no more than 1 V. The best dielectric properties were obtained with 120 nm LNO buffer layer, which was confirmed by microstructure examination. Many factors, such as orientation, compositions, crystallinity, strain and stress have been found to affect the dielectric properties. The crystalline quality of the LNO template layer affects the quality of BST. When the LNO was too thin, such as 60 nm, the microstructure was inhomogeneous and the crystalline quality was poor. On the other hand, too thick LNO was detrimental to the dielectric properties of BST, probably due to the interdiffusion between LNO layer and BST during long time deposition and post-annealing. The higher tunability

**Figure 10.** The DC bias voltage dependence of dielectric property for the Cu/BST/LNO/Pt capacitor of varied LNO

thickness δ. (a) Dielectric constant [30] and (b) dielectric loss.

/Si annealed at 700°C

251

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annealed at different temperature. It reveals that the BST/LNO/Pt/SiO<sup>2</sup>

examined.

the so-called K factor.

**Figure 8.** (1) Schematic of the capacitor, (2) the electric field dependence of dielectric property for: (a) Cu/BST/LNO (100)/Pt capacitor, (b) LNO (110)/Pt (111) and (c) Cu/BST/Pt capacitor, respectively [10].

of pulsed laser deposited BST on LNO (110)/Pt (111) (62% tunability at 262.5 kV/cm [11] and that of the BST thin film on the LNO/Pt (111) (about 51% tunability at 400 kV/cm) deposited by metal organic deposition process [12]. The results of higher capacitance are consistent with the previous structural observation, that is, the crystallization of the BST films was enhanced by LNO buffer layer. Furthermore, the dielectric loss of BST thin films was reduced for BST thin films with LNO buffer layer. It has been reported that conducting oxide like LNO aid in migration of oxygen vacancy [11], which is a possible reason for the decrease in dielectric loss of the BST thin films with LNO buffer layer. The best tunability of BST films on (100)-LNO seems to be attributed to (100) texturing of the BST films.

In addition, the effect of BST thin films thickness and post annealing on the dielectric properties of the parallel capacitor was examined. **Figure 9(a)** shows the electric field dependence of the dielectric properties of the BST thin films of different thickness with 120 nm LNO buffer layer. It should be noted that the dielectric constant did not increase with increasing thickness of BST, which is related with the interface between LNO and BST. When the LNO buffer

**Figure 9.** (a) Dielectric constant versus DC voltage with different BST thicknesses (LNO is 120 nm) and (b) the DC bias voltage dependence of dielectric loss for the thin film annealed at different temperature.

layer kept 120 nm, the tunability of BST with 100–125 nm is better than that with thicker BST. **Figure 9(b)** shows the DC bias voltage dependence of dielectric loss for the thin film annealed at different temperature. It reveals that the BST/LNO/Pt/SiO<sup>2</sup> /Si annealed at 700°C has the lowest dielectric loss, which demonstrated the LNO buffer layer help in reducing the dielectric loss. However, when the BST/LNO film was annealed at 750°C, the dielectric loss increased, and it increased rapidly with increasing bias voltage, which may be due to two reasons: one is that the Ti diffusion toward BST and reduced the effective dielectric layer, on the other hand, the lattice parameter of LNO is 3.84 and that of BST is 3.94, mismatch between BST and LNO is 2.6% the dislocations increased with increasing annealing temperature.

The effect of buffer layer thickness on the dielectric properties of the BST thin films was examined.

**Figure 10(a)** shows the DC bias voltage dependence of dielectric constant for the BST deposited on LNO buffer layer with varied thickness δ. The commonly used figure of merit (FOM) for electrically tuned device applications is the ratio of the tunability to dielectric loss (tgδ), the so-called K factor.

of pulsed laser deposited BST on LNO (110)/Pt (111) (62% tunability at 262.5 kV/cm [11] and that of the BST thin film on the LNO/Pt (111) (about 51% tunability at 400 kV/cm) deposited by metal organic deposition process [12]. The results of higher capacitance are consistent with the previous structural observation, that is, the crystallization of the BST films was enhanced by LNO buffer layer. Furthermore, the dielectric loss of BST thin films was reduced for BST thin films with LNO buffer layer. It has been reported that conducting oxide like LNO aid in migration of oxygen vacancy [11], which is a possible reason for the decrease in dielectric loss of the BST thin films with LNO buffer layer. The best tunability of BST films on (100)-LNO

**Figure 8.** (1) Schematic of the capacitor, (2) the electric field dependence of dielectric property for: (a) Cu/BST/LNO (100)/Pt

In addition, the effect of BST thin films thickness and post annealing on the dielectric properties of the parallel capacitor was examined. **Figure 9(a)** shows the electric field dependence of the dielectric properties of the BST thin films of different thickness with 120 nm LNO buffer layer. It should be noted that the dielectric constant did not increase with increasing thickness of BST, which is related with the interface between LNO and BST. When the LNO buffer

**Figure 9.** (a) Dielectric constant versus DC voltage with different BST thicknesses (LNO is 120 nm) and (b) the DC bias

voltage dependence of dielectric loss for the thin film annealed at different temperature.

seems to be attributed to (100) texturing of the BST films.

capacitor, (b) LNO (110)/Pt (111) and (c) Cu/BST/Pt capacitor, respectively [10].

250 Coatings and Thin-Film Technologies

**Table 3** shows the tunability and FOM of BST with different thickness of LNO buffer layer at applied electric field of 400 kV/cm. The tunability of BST thin film on the LNO (100)/Pt (111) was about 63% at a vertical applied field of 400 kV/cm, and it was higher than that (about 32%) of BST thin film on the Pt (111). **Figure 10(b)** reveals that the dielectric loss for the BST on 120 nm LNO was the smallest when the DC bias is no more than 1 V. The best dielectric properties were obtained with 120 nm LNO buffer layer, which was confirmed by microstructure examination. Many factors, such as orientation, compositions, crystallinity, strain and stress have been found to affect the dielectric properties. The crystalline quality of the LNO template layer affects the quality of BST. When the LNO was too thin, such as 60 nm, the microstructure was inhomogeneous and the crystalline quality was poor. On the other hand, too thick LNO was detrimental to the dielectric properties of BST, probably due to the interdiffusion between LNO layer and BST during long time deposition and post-annealing. The higher tunability

**Figure 10.** The DC bias voltage dependence of dielectric property for the Cu/BST/LNO/Pt capacitor of varied LNO thickness δ. (a) Dielectric constant [30] and (b) dielectric loss.


**Table 3.** The tunability and FOM of BST thin film on different thickness of LNO buffer layer.

is related to (100) texture of the BST films [30]. As a BST film subjected to tensile stress, a contraction occurred along the C axis, which leads to an enhancement of the in-plane oriented polar axis [10]. By a converse electrostrictive effect, the in-plane tensile stress reduces the capacitance in the thickness direction of the film [11]. When the (100)-oriented BST thin films were applied under higher electric fields, the in-plane orientation of the polar axis resulted in higher tunability. Hence, the tunability of (100)-oriented BST films was higher compared to the randomly oriented BST films.

**Figure 12(a)** show that the surface of BST/MgO/LNO is smooth and the BST grain is about 15–30 nm, the XRD results in **Figure 12(b)** reveal that the BST show (100) preferential orientation and no peak of MgO was observed because MgO annealed at 750°C is still amorphous.

**Figure 11.** XRD patterns of (a) MgO films deposited on (100)Si, (b) MgO films deposited on (111) Si, (c) BST/MgO films

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**Figure 13** shows the tunable dielectric properties of the BST thin film with and without MgO buffer layer. The frequency is fixed at 1 MHz and the additional AC voltage is 0.05 V. The thickness of interdigital Cr/Au electrodes is about 0.5 μm, and the width of the finger is 10 μm and the gap between the fingers is 5 μm. The tunability of BST thin film without MgO buffer

V. However, the insertion of MgO buffer layer increases the tunability to 62.4% for BST/MgO (150 nm)/Si and 61.5% for BST/MgO (50 nm)/Si respectively, which is more than four times than that of BST films directly deposited on Si [31]. The dielectric loss of BST without the MgO buffer layer is about 0.1–0.15, and it decreases to about 0.05 after the insertion of the MgO buffer layer. The insertion of MgO buffer layer notably improves the tunability and at the mean time reduces the dielectric loss. The higher capacitance of BST/MgO (150 nm)/Si film compared with that of BST/MgO (50 nm)/Si film resulted from that the electrical field distribution [28] in interface (Dead layer [29]). The dead layer can reduce the dielectric capacitance of the composite films because the thinner MgO buffer layer means the more electrical distribution

The C-V characteristic of the BST on MgO/LNO bilayer was shown in **Figure 14(a)**. For the interdigital thin film capacitor, the gap is 5 μm and the width of the finger is 5, 10, 15 μm,

*3.5.2. The tunable dielectrical properties of the BST thin film with and without LNO/MgO* 

*buffer layer*

and (d) BST/SiO<sup>2</sup>

/Si.

layer is 14.7% at 10.

in the interface between Si and MgO thin film.

*3.5.3. The tunable dielectrical properties of the BST/MgO/LNO multilayer*

#### **3.5. BST thin film on MgO/SiO<sup>2</sup> /Si (100) substrate and its dielectric properties**

Interdigital capacitor electrodes were patterned by lift-off process [31]. The thickness of electroplated Cu is about 5 μm which is larger than the skin depth of copper at 10 GHz.

### *3.5.1. XRD and SEM studies*

MgO thin film was deposited on Si as buffer layer and its microstructure was investigated in our previous study [31]. The sputtering was carried out under room temperature, with a partial pressure of oxygen kept at 10% and the total argon and oxygen pressure of 0.3 Pa. **Figure 11a–d** shows the XRD patterns of the MgO films and BST/MgO. MgO thin film was annealed at 1000°C for the crystallization after sputtering. BST/MgO composite thin films undergo a two-step annealing process, in which the BST film was deposited on annealed MgO buffer layer (1000°C) then BST/MgO composite film was annealed at 750°C.

**Figure 11(a)** and **(b)** reveals that the MgO thin film deposited on (100) Si prefer to (200) orientation, but the MgO thin film on (111) Si show (111) orientation. The orientation of the BST films with and without MgO buffer layer is much different in **Figure 11c** and **d**. The MgO buffer layers enhanced the BST (100) orientation and weakened the BST (111) orientations. However, the BST/Si films did not show desirable crystal orientation when annealed at 750°C.

Furthermore, considering LNO, a kind of conductive oxide, may increase the leakage of BST on it, a thin layer of MgO (<50 nm) was inserted between LNO (100 oriented) and BST The BST was deposited on MgO/LNO bilayer and then annealed at 750°C. The microstructure of the multilayers was examined.

**Figure 11.** XRD patterns of (a) MgO films deposited on (100)Si, (b) MgO films deposited on (111) Si, (c) BST/MgO films and (d) BST/SiO<sup>2</sup> /Si.

**Figure 12(a)** show that the surface of BST/MgO/LNO is smooth and the BST grain is about 15–30 nm, the XRD results in **Figure 12(b)** reveal that the BST show (100) preferential orientation and no peak of MgO was observed because MgO annealed at 750°C is still amorphous.

### *3.5.2. The tunable dielectrical properties of the BST thin film with and without LNO/MgO buffer layer*

is related to (100) texture of the BST films [30]. As a BST film subjected to tensile stress, a contraction occurred along the C axis, which leads to an enhancement of the in-plane oriented polar axis [10]. By a converse electrostrictive effect, the in-plane tensile stress reduces the capacitance in the thickness direction of the film [11]. When the (100)-oriented BST thin films were applied under higher electric fields, the in-plane orientation of the polar axis resulted in higher tunability. Hence, the tunability of (100)-oriented BST films was higher compared to

**LNO under layer thickness (nm) Tunability (%) Figure of merit (FOM)**

0 35 1.2 60 26 3.1 120 63 6.3 180 31 4.1 240 32 1.5

**Table 3.** The tunability and FOM of BST thin film on different thickness of LNO buffer layer.

Interdigital capacitor electrodes were patterned by lift-off process [31]. The thickness of elec-

MgO thin film was deposited on Si as buffer layer and its microstructure was investigated in our previous study [31]. The sputtering was carried out under room temperature, with a partial pressure of oxygen kept at 10% and the total argon and oxygen pressure of 0.3 Pa. **Figure 11a–d** shows the XRD patterns of the MgO films and BST/MgO. MgO thin film was annealed at 1000°C for the crystallization after sputtering. BST/MgO composite thin films undergo a two-step annealing process, in which the BST film was deposited on annealed MgO

**Figure 11(a)** and **(b)** reveals that the MgO thin film deposited on (100) Si prefer to (200) orientation, but the MgO thin film on (111) Si show (111) orientation. The orientation of the BST films with and without MgO buffer layer is much different in **Figure 11c** and **d**. The MgO buffer layers enhanced the BST (100) orientation and weakened the BST (111) orientations. However, the BST/Si films did not show desirable crystal orientation when annealed at 750°C. Furthermore, considering LNO, a kind of conductive oxide, may increase the leakage of BST on it, a thin layer of MgO (<50 nm) was inserted between LNO (100 oriented) and BST The BST was deposited on MgO/LNO bilayer and then annealed at 750°C. The microstructure of

troplated Cu is about 5 μm which is larger than the skin depth of copper at 10 GHz.

buffer layer (1000°C) then BST/MgO composite film was annealed at 750°C.

**/Si (100) substrate and its dielectric properties**

the randomly oriented BST films.

252 Coatings and Thin-Film Technologies

**3.5. BST thin film on MgO/SiO<sup>2</sup>**

*3.5.1. XRD and SEM studies*

the multilayers was examined.

**Figure 13** shows the tunable dielectric properties of the BST thin film with and without MgO buffer layer. The frequency is fixed at 1 MHz and the additional AC voltage is 0.05 V. The thickness of interdigital Cr/Au electrodes is about 0.5 μm, and the width of the finger is 10 μm and the gap between the fingers is 5 μm. The tunability of BST thin film without MgO buffer layer is 14.7% at 10.

V. However, the insertion of MgO buffer layer increases the tunability to 62.4% for BST/MgO (150 nm)/Si and 61.5% for BST/MgO (50 nm)/Si respectively, which is more than four times than that of BST films directly deposited on Si [31]. The dielectric loss of BST without the MgO buffer layer is about 0.1–0.15, and it decreases to about 0.05 after the insertion of the MgO buffer layer. The insertion of MgO buffer layer notably improves the tunability and at the mean time reduces the dielectric loss. The higher capacitance of BST/MgO (150 nm)/Si film compared with that of BST/MgO (50 nm)/Si film resulted from that the electrical field distribution [28] in interface (Dead layer [29]). The dead layer can reduce the dielectric capacitance of the composite films because the thinner MgO buffer layer means the more electrical distribution in the interface between Si and MgO thin film.

### *3.5.3. The tunable dielectrical properties of the BST/MgO/LNO multilayer*

The C-V characteristic of the BST on MgO/LNO bilayer was shown in **Figure 14(a)**. For the interdigital thin film capacitor, the gap is 5 μm and the width of the finger is 5, 10, 15 μm,

**Figure 12.** Characterization of the BST/MgO/LNO multilayer thin films: (a) SEM and (b) XRD patterns.

**Figure 13.** The dielectric properties of the BST thin film (a) without MgO; (b) on different thickness of MgO buffer layer [31].

respectively. **Figure 14(b)** reveal that the tunability of BST/MgO/LNO multilayer is higher with increasing width of the finger, but the dielectric loss also increases rapidly when the bias is larger than 15 V. Probably the interface between the multilayers resulted in the large dielectric loss at high bias voltage. Further intensive investigation is needed in order to obtain high tunability with small dielectric loss.

#### *3.5.4. Coplanar waveguide with BST/MgO on Si substrate*

Finally, coplanar waveguide (CPW) on Si substrate with BST/MgO multilayer were fabricated. The schematic of a CPW was shown in **Figure 15**, where g is the gap between the ground and signal line, w is the width of the signal line and the total of w and g is kept as constant (90 μm). To obtain the transmission and loss characteristics of the coplanar waveguide, S parameters were measured using a vector network analyzer (Agilent 8722ES). The frequency was swept from 5 to 15 GHz. The measured S-parameters of the waveguide are shown in **Figure 16**. The measured return loss S11 are lower than −40 dB at the central frequency of 10 GHz. Meanwhile, the measured S12 were close to −2.8 to −3.8 dB in measuring frequency

band. These results indicated that the CPW based on Si substrate with BST/MgO thin films showed good transmission property near 10 GHz. It has promising application in tunable high frequency devices although the insertion loss was little higher, which need to be further

**Figure 14.** Dielectric property for BST/MgO/LNO interdigital capacitor with different widths of the finger (gap = 5 μm)

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency…

http://dx.doi.org/10.5772/intechopen.79270

255

(a) C-V curves and (b) the tunability.

**Figure 15.** Schematic of a CPW with a double-layer dielectric on Si substrate.

optimized by design optimization of CPW and improvement of BST multilayers.

**Figure 16.** S parameters of the CPW with BST/MgO(w + g = 90 μm). (a) S12 and (b) S11.

Textured BST Thin Film on Silicon Substrate: Preparation and Its Applications for High Frequency… http://dx.doi.org/10.5772/intechopen.79270 255

**Figure 14.** Dielectric property for BST/MgO/LNO interdigital capacitor with different widths of the finger (gap = 5 μm) (a) C-V curves and (b) the tunability.

**Figure 15.** Schematic of a CPW with a double-layer dielectric on Si substrate.

respectively. **Figure 14(b)** reveal that the tunability of BST/MgO/LNO multilayer is higher with increasing width of the finger, but the dielectric loss also increases rapidly when the bias is larger than 15 V. Probably the interface between the multilayers resulted in the large dielectric loss at high bias voltage. Further intensive investigation is needed in order to obtain

**Figure 13.** The dielectric properties of the BST thin film (a) without MgO; (b) on different thickness of MgO buffer layer

**Figure 12.** Characterization of the BST/MgO/LNO multilayer thin films: (a) SEM and (b) XRD patterns.

Finally, coplanar waveguide (CPW) on Si substrate with BST/MgO multilayer were fabricated. The schematic of a CPW was shown in **Figure 15**, where g is the gap between the ground and signal line, w is the width of the signal line and the total of w and g is kept as constant (90 μm). To obtain the transmission and loss characteristics of the coplanar waveguide, S parameters were measured using a vector network analyzer (Agilent 8722ES). The frequency was swept from 5 to 15 GHz. The measured S-parameters of the waveguide are shown in **Figure 16**. The measured return loss S11 are lower than −40 dB at the central frequency of 10 GHz. Meanwhile, the measured S12 were close to −2.8 to −3.8 dB in measuring frequency

high tunability with small dielectric loss.

[31].

254 Coatings and Thin-Film Technologies

*3.5.4. Coplanar waveguide with BST/MgO on Si substrate*

**Figure 16.** S parameters of the CPW with BST/MgO(w + g = 90 μm). (a) S12 and (b) S11.

band. These results indicated that the CPW based on Si substrate with BST/MgO thin films showed good transmission property near 10 GHz. It has promising application in tunable high frequency devices although the insertion loss was little higher, which need to be further optimized by design optimization of CPW and improvement of BST multilayers.

### **4. Conclusions**

The effect of LNO/MgO buffer layer on the microstructure and dielectric properties of BST thin films was investigated. The MgO, BST and LNO films were deposited on the Si substrate by RF magnetron sputtering. It reveals that the orientation of the LNO thin films were dominated by the substrate temperature, highly (100)-oriented LNO was obtained when the substrate temperature was 300°C, and the orientation of BST thin film was tailored by the introduction of LNO buffer layer. Highly (100)-oriented (Ba0.5Sr0.5)TiO3 /LaNiO<sup>3</sup> heterostructures were obtained on Pt(111) by RF sputtering, and LNO buffer layers enhance the growth of BST grains. The tunability were greatly improved to 63% by the introduction of (100)-textured LNO buffer layer with proper thickness, and FOM of BST thin films was also greatly improved. Also, the BST thin film interdigital capacitors were fabricated on silicon with MgO or LNO/MgO as buffer layer. The results show that insertion of the MgO buffer layer can enhance the tunability of the BST film and simultaneously reduce the dielectric loss. The MgO buffer layer can also enhance the crystallization of the BST thin films. However, LNO/MgO bilayer can greatly increase the tunability, at the same time the dielectric loss is large when the applied voltage is more than 10 V. Finally, the CPW with BST/ MgO were fabricated and their S were tested, the CPW showed good transmission property near 10 GHz and showed promising application in tunable high frequency devices.

[3] Zhu X, Peng HW, Miao J, Zheng DN. Fabrication and characterization of tunable dielec-

[4] Kang BS, Lee J, Stan L, Lee JK, DePaula RF, Arendt PN. Dielectric properties of epitaxial

[5] Chu CM, Lin P. Electrical properties and crystal structure of (Ba,Sr)TiO films prepared at low temperatures on a LaNiO electrode by radio-frequency magnetron sputtering.

[8] Chen XY, Wong KH, Mak CL, Yin XB, Liu JM, Wang M, Liu ZG. The orientation-selec-

[9] Wang C, Cheng BL, Wang SY, Lu HB, Zhou YL, Chen ZH, Yang GZ.Effects of oxygen pressure on lattice parameter, orientation, surface morphology and deposition rate of (Ba0.02Sr0.98) TiO3 thin films grown on MgO substrate by pulsed laser deposition. Thin Solid Films. 2005;

[11] Tang XG, Xiong HF, Jiang LL, Chan LHW. Dielectric properties and high tunability of (1

[12] Gao YH, Sun JL, Ma JH, et al. Improved dielectric and electrical properties of (Ba,Sr)TiO<sup>3</sup>

[14] Chun YH, Hong JS, Peng B, Jackson TJ, Lancaster MJ. An Electronically Tuned Bandstop Filter Using BST Varactors. Microwave Conference, EuMC, 38th European. 2008.

MgO heterostructured thin films by thin MgO layers insertion. Journal of Sol-Gel Science and

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∕Si using biaxially oriented ion-beam-assisted-deposited

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Ba0.6Sr0.4TiO3

### **Acknowledgements**

The authors would like to express their gratitude for the support from the National Natural Science Foundation of China (no. 60701012).

### **Author details**

Conchun Zhang\*, Jianze Huang, Chunsheng Yang and Guifu Ding

\*Address all correspondence to: zhcc@sjtu.edu.cn

National Key Laboratory of Nano/Micro Fabrication Technology, Department of Micro/ Nano Electronics, Shanghai Jiao Tong University, Shanghai, PR China

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256 Coatings and Thin-Film Technologies

**Acknowledgements**

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Science Foundation of China (no. 60701012).

\*Address all correspondence to: zhcc@sjtu.edu.cn

films with conductive SrRuO<sup>3</sup>

Conchun Zhang\*, Jianze Huang, Chunsheng Yang and Guifu Ding

Nano Electronics, Shanghai Jiao Tong University, Shanghai, PR China

properties of tetragonally distorted (Ba0.5Sr0.5)TiO3

The effect of LNO/MgO buffer layer on the microstructure and dielectric properties of BST thin films was investigated. The MgO, BST and LNO films were deposited on the Si substrate by RF magnetron sputtering. It reveals that the orientation of the LNO thin films were dominated by the substrate temperature, highly (100)-oriented LNO was obtained when the substrate temperature was 300°C, and the orientation of BST thin film was tailored by

structures were obtained on Pt(111) by RF sputtering, and LNO buffer layers enhance the growth of BST grains. The tunability were greatly improved to 63% by the introduction of (100)-textured LNO buffer layer with proper thickness, and FOM of BST thin films was also greatly improved. Also, the BST thin film interdigital capacitors were fabricated on silicon with MgO or LNO/MgO as buffer layer. The results show that insertion of the MgO buffer layer can enhance the tunability of the BST film and simultaneously reduce the dielectric loss. The MgO buffer layer can also enhance the crystallization of the BST thin films. However, LNO/MgO bilayer can greatly increase the tunability, at the same time the dielectric loss is large when the applied voltage is more than 10 V. Finally, the CPW with BST/ MgO were fabricated and their S were tested, the CPW showed good transmission property

near 10 GHz and showed promising application in tunable high frequency devices.

The authors would like to express their gratitude for the support from the National Natural

National Key Laboratory of Nano/Micro Fabrication Technology, Department of Micro/

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**Chapter 13**

**Provisional chapter**

O4

) thin films

**Spin-Coating Technique for Fabricating Nickel Zinc**

**Spin-Coating Technique for Fabricating Nickel Zinc** 

Functional nanoferrite thin films are used in various fields of our life. There are many different methods used to fabricate thin films including sputter deposition, flash laser evaporation pulsed laser deposition (PLD), chemical vapor deposition (PVD) and spin-coating process. In each of these methods, it produces an amorphous phase of the deposited film. To produce a crystalline film, an additional high-temperature processing is required. The high-temperature process can lead to considerable constraints in combining the desirable characteristics of a crystalline nanoferrite thin film with those of thermally unstable substrates and other device components. High-temperature thin-film processing is also a considerable cost to manufacturing. This chapter will report a simple procedure of the

and spin-coating method in coating a chemical solution. This method generally provides for both low-temperature deposition and crystallization of NiZn nanoferrite thin films. **Keywords:** NiZn ferrite thin films, spin coating, structural, magnetic, optical properties

NiZn nanocrystalline ferrite thin films have a spinel crystal structure which have been a subject of extensive attempt because of their potential applications in high-density magneto-optic recording devices, magnetic refrigeration and microwave materials due to its high electrical resistivity, low magnetic coercivity and low eddy current losses. NiZn nanoferrite thin films are structure sensitive, and it is not easy to produce a stoichiometric and point-defectfree NiZn ferrite, for high-resistivity applications**.** In NiZn ferrite thin-film fabrication, the

sol-gel precursor method for fabrication of NiZn nanoferrite (Ni0.3Zn0.7Fe2

**O4) Thin Films**

DOI: 10.5772/intechopen.80461

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

© 2018 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use,

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

**Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films**

Yusnita Yusuf, Raba'ah Syahidah Azis and

Yusnita Yusuf, Raba'ah Syahidah Azis and

Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

**1. Introduction of NiZn ferrite thin films**

Muhammad Syazwan Mustaffa

Muhammad Syazwan Mustaffa

**Nanoferrite (Ni0.3Zn0.7Fe<sup>2</sup>**

http://dx.doi.org/10.5772/intechopen.80461

**Abstract**


#### **Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe<sup>2</sup> O4) Thin Films**

DOI: 10.5772/intechopen.80461

Yusnita Yusuf, Raba'ah Syahidah Azis and Muhammad Syazwan Mustaffa Yusnita Yusuf, Raba'ah Syahidah Azis and Muhammad Syazwan Mustaffa

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

http://dx.doi.org/10.5772/intechopen.80461

#### **Abstract**

[17] Liu SS, Ma BH, Narayanan M, Chao S, Koritala R, Balachandran U. Improved properties of barium strontium titanate thin films grown on copper foils by pulsed laser deposition using a self-buffered layer. Journal of Physics D: Applied Physics. 2012;**45**(6pp):175304

[18] Yan DX, Xu ZP, Chen XL, Xiao DQ, Yu P, Zhu JG. Microstructure and electrical proper-

[19] Kwon SR, Huang WB, Zhang SJ, Yuan FG, Jiang XN. Flexoelectric sensing using a multilayered barium strontium titanate structure. Smart Materials and Structures. 2013;

[20] Iskandar J, Syafutra H, Juansah JJ, Irzaman. Characterizations of electrical and optical properties on ferroelectric photodiode of barium strontium titanate (Ba0.5Sr0.5TiO3

[21] Hu SD, Li H, Tzou HS. Distributed flexoelectric structural sensing: Theory and experi-

[22] Kwon SR, Huang W, Shu L, Yuan FG, Maria JP, Jiang X. Flexoelectricity in barium stron-

[23] Nam SH, Lee WJ, Kim HJ. Oriented growth of SrTiO, thin films on Si substrate by radio frequency I magnetron sputtering. Journal of Physics D: Applied Physics. 1994;**27**:866-870

[24] Song TK, Ahn JS, Choi HS, Noh TW, Kwun SI. Infrared properties of epitaxial SrTiO<sup>3</sup> thin films on MgO(001) substrates. Journal- Korean Physical Society. 1997;**30**(3):623-627

[25] Yang YF, Nordman JE, Lee JU. Effects of deposition conditions on stoichiometry of off-axis RF sputtered BiSrCaCuO thin films. IEEE Transactions on Applied Super-

[26] Zhang C, Hou J, Rao R, Yang C, Ding G. Effects of process parameters on the LaNiO<sup>3</sup> thin films deposited by radio-frequency magnetron sputtering. Thin Solid Films.

[28] Kim DY, Lee SG, Park YK, Park SJ. Effect of ambient gas pressure on the preferred orientation of barium titanate thin films prepared by pulsed laser deposition. Japanese

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con substrate with MgO buffer layer. International Journal of Nanomanufacturing. 2011;

Thin Film on Si Substrate. Journal of Shanghai Jiaotong University.

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satellite technology. Procedia Environmental Sciences. 2015;**24**:324-328

ment. Journal of Sound and Vibration. 2015;**348**:126-136

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films based on the annealing time differences and its development as light sensor on

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**7**(5/6):567-574

2009;**14**(2):133-136

**22**:115017

258 Coatings and Thin-Film Technologies

Functional nanoferrite thin films are used in various fields of our life. There are many different methods used to fabricate thin films including sputter deposition, flash laser evaporation pulsed laser deposition (PLD), chemical vapor deposition (PVD) and spin-coating process. In each of these methods, it produces an amorphous phase of the deposited film. To produce a crystalline film, an additional high-temperature processing is required. The high-temperature process can lead to considerable constraints in combining the desirable characteristics of a crystalline nanoferrite thin film with those of thermally unstable substrates and other device components. High-temperature thin-film processing is also a considerable cost to manufacturing. This chapter will report a simple procedure of the sol-gel precursor method for fabrication of NiZn nanoferrite (Ni0.3Zn0.7Fe2 O4 ) thin films and spin-coating method in coating a chemical solution. This method generally provides for both low-temperature deposition and crystallization of NiZn nanoferrite thin films.

**Keywords:** NiZn ferrite thin films, spin coating, structural, magnetic, optical properties

### **1. Introduction of NiZn ferrite thin films**

NiZn nanocrystalline ferrite thin films have a spinel crystal structure which have been a subject of extensive attempt because of their potential applications in high-density magneto-optic recording devices, magnetic refrigeration and microwave materials due to its high electrical resistivity, low magnetic coercivity and low eddy current losses. NiZn nanoferrite thin films are structure sensitive, and it is not easy to produce a stoichiometric and point-defectfree NiZn ferrite, for high-resistivity applications**.** In NiZn ferrite thin-film fabrication, the

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2018 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

accurate composition control and the uniformity throughout the film deposition are important. It is well known that properties of ferrite materials strongly depend on the preparation conditions. This issue is important in NiZn ferrite thin films because the effect of temperatures will lead to the change in the chemical composition of the ferrite films. These also will result in non-uniformity of film composition and the magnetic hysteresis parameter of ferrites. Hightemperature synthesis of NiZn thin-film ferrite results in the evaporation of some constituents that lead to the nonstoichiometry, and zinc volatilization at higher temperature can result in the formation of Fe2+ ions that lead to increase the electron hopping and reducing the electrical resistivity [1]. Therefore, a low-temperature synthesis is required for the synthesis of NiZn ferrite film. Properties of ferrite film depend on the preparation route, due to its strong influence on type of the film (polycrystalline and epitaxial), particle size, chemical homogeneity, microstructure and cationic distribution between tetrahedral and octahedral sublattice sites [2, 3]. Synthesis of ferrite thin film is of great interest among researchers in this field of study. Widely used techniques are utilized to produce desirable final product of nickel zinc ferrite. These techniques can be classified into two major techniques which are the conventional technique and the nonconventional technique. The starting materials are conventionally oxides or precursors of oxide of the cations. This process involves the interdiffusion of the various metal ions of preselected compositions to form a mixed crystal. The nonconventional powder processing in a liquid medium may produce intermediate, finely divided mixed hydroxides or mixed organic salts to assist the subsequent diffusion process [4]. Most of ferrite films have been prepared using sputtering and pulsed laser deposition. Somehow, sol-gel method is a kind of potential film preparation process, which possesses advantages of chemical homogeneity, easy component adjustment, low calcination temperature and low cost. Spin coating gives an advantage to liquid film that leads to uniformity in thickness during spin-off [5]. Once uniform, it tends to remain provided the viscosity is not shear-dependent and does not vary over the substrate. Other than that, sol-gel and hydrothermal routes of ferrite synthesis have shown increasing importance. Recent years are marked by growing interest in sol-gel processed films in new areas, particularly in microelectronics. This is mainly due to intensively developing applications of silicate or siloxane sol-gel films in the VLSI multilevel interconnection process, the preparation of ferroelectric films for nonvolatile memory [6].

hydrothermal grow. Nickel zinc ferrite thin films are successfully prepared using spin-deposited citrate-precursor route [12]. The formation of crystalline film at low temperature even though films were found to be X-ray amorphous revealed the formation of uniform grains in nanometer size range. Besides that, NiZn ferrite film was successfully fabricated by using photosensitive sol-gel method. The photosensitive gel film can be the photoresist of itself during the

rising and heat treatment [13]. The great potential of combining the microwave technique with nonaqueous sol-gel chemistry was successful [14]. Many transition metal ferrite nanoparticles with high crystallinity are uniformly morphological besides homogenous metal ferrite thin films on flat and curved substrates. The thickness of film can easily be adjusted in the range of 20–80 nm using precursor concentration. Other synthesis techniques of nickel zinc ferrite thin films are chemical vapor deposition, spray pyrolysis, sputtering, pulsed laser deposition and spin spray. Of these methods the earliest used was vapor deposition of metals followed by oxidation [15]. The films were porous and polycrystalline and approximately 1000 Å thick. Many common ferrites produced by this method were single-phase spinel in crystal structure. Spray pyrolysis is complicated and expensive and required special equipment and sometimes high processing temperature above 500°C. By using spray pyrolysis technique [3, 16], very homogenous ferrite thin films were obtained with good reproducibility. Nevertheless, for spinel ferrite thin-film growth, this method is used rarely. Also only a few works can be found on NiZn ferrites, obtained by spray pyrolysis where a focus on the investigation of microstructural, optical and magnetic properties was held [17]. Apart from that, the effect of oxygen plasma treatment on magnetic and NiZn ferrite films using the spin-spray plating method has been employed [18]. The oxygen plasma treatment increased the number of nucleation sites of ferrite and enhanced adhesion of the films to the substrates. It has been reported that spinel Zn ferrite can be synthesized without substrate heating by pulsed laser deposition [19]. However, this technique needs post-deposition and requires sputtered film at a high temperature to grow the spinel ferrite structure. The results were optimized and obtained 4000 Å NiZn ferrite films with low in-plane coercivity of *H*<sup>c</sup> = 15.2 Oe and relatively high saturation magnetization

O4

Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films

. Some other research works was working on preparing NiZn ferrite films

by magnetron sputtering method. Most of the sputtered ferrite films must be deposited at a high substrate temperature and need high heat treatment to obtain ordered spinel structure. Sputtering method is prepared at room temperature without any post-annealing treatment [20]. By controlling the relative oxygen flow, grain size in a range 10–20 nm was developed.

measurement results are affected by the crystallinity, grain dimension and cation distribution.

NiZn ferrite thin film is prepared by a sol-gel process and spin-coating technique. The start-

)2 ·6H<sup>2</sup>

It revealed a maximum saturation magnetization of about 151 emu/cm3

**3. Brief overview of preparation methods**

ing materials nickel nitrate hexahydrate (Ni(NO3

) 3 ·9H<sup>2</sup>

**3.1. Preparation of NiZn ferrite thin films**

nitrate nonahydrate (Fe(NO3

film is obtained through UV radiation,

http://dx.doi.org/10.5772/intechopen.80461

261

. The static magnetic

O) (Sigma Aldrich, 99.999%), iron

O) (Alfa Aesar, 99.999%) and zinc nitrate hexahydrate

preparation process. The fine pattern of Ni0.5Zn0.5Fe2

*M*<sup>s</sup> = 318 emu/cm3

### **2. Literature review on synthesis techniques of NiZn ferrite thin films**

Spin coating is widely used in modern optical and microelectronic industries [7]. The understanding of its underlying physics remains limited, a fact attributed to the lack of experimental data for the evolution of various parameters during the process [8, 9], leading to the need for new evolution tools. A relative study of nickel zinc ferrite by sol-gel route and conventional solid-state reaction was carried out [10]. It was claimed that the homogeneity and high purity in the sol-gel samples and small grains confirmed the finer particles. Ni0.36Zn0.64Fe2 O4 (NZF) thick films have been synthesized using sol-gel dip-coating method [11]. Combination with dispersion of ceramic NZF particles in starting sols has been proved to be useful for producing thick nickel zinc ferrite films. The best NZF powders are formed from dispersing at 300°C by hydrothermal grow. Nickel zinc ferrite thin films are successfully prepared using spin-deposited citrate-precursor route [12]. The formation of crystalline film at low temperature even though films were found to be X-ray amorphous revealed the formation of uniform grains in nanometer size range. Besides that, NiZn ferrite film was successfully fabricated by using photosensitive sol-gel method. The photosensitive gel film can be the photoresist of itself during the preparation process. The fine pattern of Ni0.5Zn0.5Fe2 O4 film is obtained through UV radiation, rising and heat treatment [13]. The great potential of combining the microwave technique with nonaqueous sol-gel chemistry was successful [14]. Many transition metal ferrite nanoparticles with high crystallinity are uniformly morphological besides homogenous metal ferrite thin films on flat and curved substrates. The thickness of film can easily be adjusted in the range of 20–80 nm using precursor concentration. Other synthesis techniques of nickel zinc ferrite thin films are chemical vapor deposition, spray pyrolysis, sputtering, pulsed laser deposition and spin spray. Of these methods the earliest used was vapor deposition of metals followed by oxidation [15]. The films were porous and polycrystalline and approximately 1000 Å thick. Many common ferrites produced by this method were single-phase spinel in crystal structure. Spray pyrolysis is complicated and expensive and required special equipment and sometimes high processing temperature above 500°C. By using spray pyrolysis technique [3, 16], very homogenous ferrite thin films were obtained with good reproducibility. Nevertheless, for spinel ferrite thin-film growth, this method is used rarely. Also only a few works can be found on NiZn ferrites, obtained by spray pyrolysis where a focus on the investigation of microstructural, optical and magnetic properties was held [17]. Apart from that, the effect of oxygen plasma treatment on magnetic and NiZn ferrite films using the spin-spray plating method has been employed [18]. The oxygen plasma treatment increased the number of nucleation sites of ferrite and enhanced adhesion of the films to the substrates. It has been reported that spinel Zn ferrite can be synthesized without substrate heating by pulsed laser deposition [19]. However, this technique needs post-deposition and requires sputtered film at a high temperature to grow the spinel ferrite structure. The results were optimized and obtained 4000 Å NiZn ferrite films with low in-plane coercivity of *H*<sup>c</sup> = 15.2 Oe and relatively high saturation magnetization *M*<sup>s</sup> = 318 emu/cm3 . Some other research works was working on preparing NiZn ferrite films by magnetron sputtering method. Most of the sputtered ferrite films must be deposited at a high substrate temperature and need high heat treatment to obtain ordered spinel structure. Sputtering method is prepared at room temperature without any post-annealing treatment [20]. By controlling the relative oxygen flow, grain size in a range 10–20 nm was developed. It revealed a maximum saturation magnetization of about 151 emu/cm3 . The static magnetic measurement results are affected by the crystallinity, grain dimension and cation distribution.

### **3. Brief overview of preparation methods**

### **3.1. Preparation of NiZn ferrite thin films**

accurate composition control and the uniformity throughout the film deposition are important. It is well known that properties of ferrite materials strongly depend on the preparation conditions. This issue is important in NiZn ferrite thin films because the effect of temperatures will lead to the change in the chemical composition of the ferrite films. These also will result in non-uniformity of film composition and the magnetic hysteresis parameter of ferrites. Hightemperature synthesis of NiZn thin-film ferrite results in the evaporation of some constituents that lead to the nonstoichiometry, and zinc volatilization at higher temperature can result in the formation of Fe2+ ions that lead to increase the electron hopping and reducing the electrical resistivity [1]. Therefore, a low-temperature synthesis is required for the synthesis of NiZn ferrite film. Properties of ferrite film depend on the preparation route, due to its strong influence on type of the film (polycrystalline and epitaxial), particle size, chemical homogeneity, microstructure and cationic distribution between tetrahedral and octahedral sublattice sites [2, 3]. Synthesis of ferrite thin film is of great interest among researchers in this field of study. Widely used techniques are utilized to produce desirable final product of nickel zinc ferrite. These techniques can be classified into two major techniques which are the conventional technique and the nonconventional technique. The starting materials are conventionally oxides or precursors of oxide of the cations. This process involves the interdiffusion of the various metal ions of preselected compositions to form a mixed crystal. The nonconventional powder processing in a liquid medium may produce intermediate, finely divided mixed hydroxides or mixed organic salts to assist the subsequent diffusion process [4]. Most of ferrite films have been prepared using sputtering and pulsed laser deposition. Somehow, sol-gel method is a kind of potential film preparation process, which possesses advantages of chemical homogeneity, easy component adjustment, low calcination temperature and low cost. Spin coating gives an advantage to liquid film that leads to uniformity in thickness during spin-off [5]. Once uniform, it tends to remain provided the viscosity is not shear-dependent and does not vary over the substrate. Other than that, sol-gel and hydrothermal routes of ferrite synthesis have shown increasing importance. Recent years are marked by growing interest in sol-gel processed films in new areas, particularly in microelectronics. This is mainly due to intensively developing applications of silicate or siloxane sol-gel films in the VLSI multilevel interconnection process, the preparation of ferroelectric films for nonvolatile memory [6].

**2. Literature review on synthesis techniques of NiZn ferrite thin** 

in the sol-gel samples and small grains confirmed the finer particles. Ni0.36Zn0.64Fe2

Spin coating is widely used in modern optical and microelectronic industries [7]. The understanding of its underlying physics remains limited, a fact attributed to the lack of experimental data for the evolution of various parameters during the process [8, 9], leading to the need for new evolution tools. A relative study of nickel zinc ferrite by sol-gel route and conventional solid-state reaction was carried out [10]. It was claimed that the homogeneity and high purity

thick films have been synthesized using sol-gel dip-coating method [11]. Combination with dispersion of ceramic NZF particles in starting sols has been proved to be useful for producing thick nickel zinc ferrite films. The best NZF powders are formed from dispersing at 300°C by

O4

(NZF)

**films**

260 Coatings and Thin-Film Technologies

NiZn ferrite thin film is prepared by a sol-gel process and spin-coating technique. The starting materials nickel nitrate hexahydrate (Ni(NO3 )2 ·6H<sup>2</sup> O) (Sigma Aldrich, 99.999%), iron nitrate nonahydrate (Fe(NO3 ) 3 ·9H<sup>2</sup> O) (Alfa Aesar, 99.999%) and zinc nitrate hexahydrate (Zn(NO3 ) 2 ·6H<sup>2</sup> O) (Alfa Aesar, 99.999%) with high purity were used as a precursor for the starting sol preparation. The materials are in metal nitrate hydrates which are soluble in alcohol solvents. Acetone and deionized water were used as a medium for sol-gel reaction. Acetic acid (C<sup>6</sup> H8 O7 H2 O) (Alfa Aesar, 99.99%) acts as the chelating agent. The precursors were dissolved in deionized water and stirred for 15 min with a molar ratio of Ni:Zn:Fe = 1:1:2 using hot plate. The former salt solution was dissolved into acetic acid solution with a molar ratio of 1:1 and stirred for 3 h at 80°C. A sol-gel formed was left 24 h for age.

### **3.2. Thin film deposition and spin-coating technique**

The thin film was deposited on indium tin oxide (ITO) glass. ITO has higher melting point around 1926°C. There is no phase change of substrate during deposition of the film. The typical properties of ITO glass substrate are listed in **Table 1**. The film deposition consists of substrate wash and spin coating. The steps are to wash the substrates firstly with distilled water in ultrasonic bath for 15 min. The substrate was then washed in ultrasonic bath using acetone liquid for 15 min. Coating was carried out in a clean room by using a spin coater. The setting parameters were listed in **Table 2**. The aged sol of 1.0 ml (Section 3.1) was dropped on ITO glass substrate and spin coated for 25 s at 3500 rpm (revolutions per min). The deposition was repeated several times to obtain the required thickness (300 nm). The film thickness can be controlled by the number of coating. The film with desired thickness can be obtained by repeating the deposition cycle. Then drying films were performed in a room temperature for a few minutes and annealed in air at temperature 400, 500, 600 and 700°C, respectively, with an increment of 100°C for 1 h. Annealing process was performed in a box furnace with rate of 5°C/min.

observed using a FEI Nova NanoSEM 230 field emission scanning electron microscope. The distributions of grain sizes were obtained by taking more than 200 different grain images for the sample using J-image software. Hysteresis parameters of the loop of the

Speed rate (rpm) 3500 rpm (constant speed)

Temperature Room temperature

Spin time (s) 25 s Volume of solution 1.0 ml Acceleration and deceleration 140 rpm/s

model 7404 Lake Shore. UV-Vis SHIMADZU model UV-3600 spectrophotometer has been used to analyse the optical transmission of the NiZn ferrite thin film in the wavelength

**Figure 1** shows the XRD pattern of spin coating and air-annealed ferrite thin films on the ITO glass substrate. The XRD patterns show single-cubic spinel-phase structures of (220),

code 74-2081 and 82-1049, respectively. Plane (311) is most intense in each annealing temperature, whereas others are at relatively low intensity [22]. These plane formed nickel zinc ferrite phases. The small peak intensities in XRD pattern revealed the existence of fine grain nanocrystalline with the most part as amorphous. The height of the highest XRD intensity is more intense at high temperature and improves the crystallinity of the films. As the annealing temperature increases, the grain size also increases, as indicated in the narrowing of the XRD spectrum lines. Increasing annealing temperature will enhance the crystallinity besides releasing the internal strains within the samples which results in better optical and magnetic properties. The intensity of the (311) peak increases as a function of the substrate temperature showing an improvement of the film crystallinity. Moreover the peak intensity increases, while its full width at half maximum (FWHM) decreases. Further increasing of substrate temperature leads to a slight decrease of the peak intensity for films prepared at 700°C. It indicates a saturation of film crystallinity. The crystallite sizes of all ferrite thin films are found to be between 16 and 18 nm. The average crystallite size, *D* (**Table 4**), was determined using

O4

thin films were measured by using vibrating sample magnetometer (VSM)

Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films

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263

ferrite thin films according to JCPDS reference

*<sup>β</sup>* cos*<sup>θ</sup>* (1)

Ni0.3Zn0.7Fe2

range 200–800 nm.

O4

Reproduced with permission from [21].

**Table 2.** Spin-coating setting parameters.

**4. Research findings**

**4.1. Structural analysis**

(311), (400), (511) and (440) in Ni0.3Zn0.7Fe2

the Scherrer's formula [23] as given by Eq. (1):

*D* = \_\_\_\_\_\_ 0.9*<sup>λ</sup>*

#### **3.3. Ni0.3Zn0.7Fe2 O4 thin-film characterizations**

The X-ray diffraction (XRD) pattern of Ni0.3Zn0.7Fe2 O4 thin films was obtained by using a Philips X'pert diffractometer model 7602 EA Almelo operating at 40 kV/30 mA in the 2θ range (20–80°) with CuKα radiation, λ = 1.5418 Å. The microstructural properties were


**Table 1.** Experimental parameters of Ni0.3Zn0.7Fe2 O4 thin-film sol-gel spin-coating deposition process.


**Table 2.** Spin-coating setting parameters.

(Zn(NO3 ) 2 ·6H<sup>2</sup>

acid (C<sup>6</sup>

of 5°C/min.

**3.3. Ni0.3Zn0.7Fe2**

**O4**

**Table 1.** Experimental parameters of Ni0.3Zn0.7Fe2

H8 O7 H2

262 Coatings and Thin-Film Technologies

O) (Alfa Aesar, 99.999%) with high purity were used as a precursor for the

O) (Alfa Aesar, 99.99%) acts as the chelating agent. The precursors were dis-

O4

thin-film sol-gel spin-coating deposition process.

Philips X'pert diffractometer model 7602 EA Almelo operating at 40 kV/30 mA in the 2θ range (20–80°) with CuKα radiation, λ = 1.5418 Å. The microstructural properties were

thin films was obtained by using a

starting sol preparation. The materials are in metal nitrate hydrates which are soluble in alcohol solvents. Acetone and deionized water were used as a medium for sol-gel reaction. Acetic

solved in deionized water and stirred for 15 min with a molar ratio of Ni:Zn:Fe = 1:1:2 using hot plate. The former salt solution was dissolved into acetic acid solution with a molar ratio of

The thin film was deposited on indium tin oxide (ITO) glass. ITO has higher melting point around 1926°C. There is no phase change of substrate during deposition of the film. The typical properties of ITO glass substrate are listed in **Table 1**. The film deposition consists of substrate wash and spin coating. The steps are to wash the substrates firstly with distilled water in ultrasonic bath for 15 min. The substrate was then washed in ultrasonic bath using acetone liquid for 15 min. Coating was carried out in a clean room by using a spin coater. The setting parameters were listed in **Table 2**. The aged sol of 1.0 ml (Section 3.1) was dropped on ITO glass substrate and spin coated for 25 s at 3500 rpm (revolutions per min). The deposition was repeated several times to obtain the required thickness (300 nm). The film thickness can be controlled by the number of coating. The film with desired thickness can be obtained by repeating the deposition cycle. Then drying films were performed in a room temperature for a few minutes and annealed in air at temperature 400, 500, 600 and 700°C, respectively, with an increment of 100°C for 1 h. Annealing process was performed in a box furnace with rate

1:1 and stirred for 3 h at 80°C. A sol-gel formed was left 24 h for age.

 **thin-film characterizations**

Glass substrate ITO coating glass Size of substrate 25 × 12 × 1.1 mm

Gel drop 1 ml Wash ultrasonic bath 15 min Spin rotations per min (rpm) 3500 Duration of cycle 25 s Number of cycle 5

Annealing temperature 400–700°C Cooling rate after annealing 5°C/min

O4

The X-ray diffraction (XRD) pattern of Ni0.3Zn0.7Fe2

**3.2. Thin film deposition and spin-coating technique**

observed using a FEI Nova NanoSEM 230 field emission scanning electron microscope. The distributions of grain sizes were obtained by taking more than 200 different grain images for the sample using J-image software. Hysteresis parameters of the loop of the Ni0.3Zn0.7Fe2 O4 thin films were measured by using vibrating sample magnetometer (VSM) model 7404 Lake Shore. UV-Vis SHIMADZU model UV-3600 spectrophotometer has been used to analyse the optical transmission of the NiZn ferrite thin film in the wavelength range 200–800 nm.

### **4. Research findings**

#### **4.1. Structural analysis**

**Figure 1** shows the XRD pattern of spin coating and air-annealed ferrite thin films on the ITO glass substrate. The XRD patterns show single-cubic spinel-phase structures of (220), (311), (400), (511) and (440) in Ni0.3Zn0.7Fe2 O4 ferrite thin films according to JCPDS reference code 74-2081 and 82-1049, respectively. Plane (311) is most intense in each annealing temperature, whereas others are at relatively low intensity [22]. These plane formed nickel zinc ferrite phases. The small peak intensities in XRD pattern revealed the existence of fine grain nanocrystalline with the most part as amorphous. The height of the highest XRD intensity is more intense at high temperature and improves the crystallinity of the films. As the annealing temperature increases, the grain size also increases, as indicated in the narrowing of the XRD spectrum lines. Increasing annealing temperature will enhance the crystallinity besides releasing the internal strains within the samples which results in better optical and magnetic properties. The intensity of the (311) peak increases as a function of the substrate temperature showing an improvement of the film crystallinity. Moreover the peak intensity increases, while its full width at half maximum (FWHM) decreases. Further increasing of substrate temperature leads to a slight decrease of the peak intensity for films prepared at 700°C. It indicates a saturation of film crystallinity. The crystallite sizes of all ferrite thin films are found to be between 16 and 18 nm. The average crystallite size, *D* (**Table 4**), was determined using the Scherrer's formula [23] as given by Eq. (1):

$$D = \frac{0.9\lambda}{\beta \cos \theta} \tag{1}$$

**Figure 1.** XRD pattern of air-annealed Ni0.3Zn0.7Fe2 O4 ferrite thin films. Reproduced with permission from [21].

where *D* is the crystallite size, *β* is full width at half maximum of the diffraction peak, *λ* is the wavelength of 1.54 Å, and *θ* is scanning angle.

The lattice constants of these films were calculated using indexing method [24] given by Eq. (2):

$$\frac{\lambda}{4\,a^2} = \frac{\sin^2\theta}{d} = \frac{\sin^2\theta}{(h^2 + k^2 + l^2)}\tag{2}$$

28.12 nm (600°C) and 41.32 nm (700°C). The grains of the films are spherical and uniform, and cohesion of grains is due to the magnetic attraction. The average grain size of the films is presented in **Figure 2**. The histogram of grain size distribution shifted to the larger grain size as the

thin films from XRD spectra.

**Figure 3** illustrates the cross section of the samples annealed at 400, 500, 600 and 700°C, respectively. The deposited films were uniform with two cycles of number deposition cycle. It was found that the thin films have thickness in the range of 145.7–285.6 nm which was confirmed by cross-sectional FESEM images. The grain size over 26 nm was further increased with a higher annealing temperature. Accordingly, the number of grain sizes beyond the single domain to multidomain critical size also increased. Therefore, the number of domain wall increased as the movement of domain wall contribution to make ease of magnetization

ship with the grain size. Annealing is a process related to secondary grain growth in the film. Thompson discussed the secondary grain growth mechanism and came to the conclusion that the secondary grain growth is driven by the reduction of the total grain boundary energy. Since the grain boundary energy is film-thickness-dependent, the secondary grain growth

The plots of magnetization, *M*, against magnetic field strength, *H* (*M−H* hysteresis loop), for

teresis shape is narrow and has linear loops which have a low saturation magnetization, *M*<sup>s</sup>

could be attributed to the varied grain size and crystallinity. The lower value of saturation

, and coercivity, *H*<sup>c</sup>

these curves and have been listed for various annealing temperatures in **Table 4**. The *M*<sup>s</sup>

films annealed at 400, 500, 600 and 700°C were shown in **Figure 5**. The hys-

results (Section 4.3),

O ferrite thin films and shows its relation-

, values have been directly extracted from

value.

.

and

annealing temperature increased. However, based on the coercivity, *Hc*

O4

**Figure 4** presented the thickness of the Ni0.7Zn0.3Fe2

**Table 3.** Structural parameters of Ni0.3Zn0.7Fe2

rate increases when the film thickness is reduced [29].

increased [28].

**4.3. Magnetic properties**

O4

The saturation magnetization, *M*<sup>s</sup>

Ni0.3Zn0.7Fe2

*H*c

the transition from 600 to 700°C of annealing temperature exhibits a fall of the *Hc*

**Annealing temperature (°C) 400 500 600 700** Rel. intensity counts (%) 100 100 100 100 Position (2θ) 35.3661 35.3717 35.3991 35.3247 FWHM (2θ) 0.5215 0.5371 0.5116 0.4723 *d*-spacing (nm) 2.53595 2.53557 2.53207 2.53993 Crystallite size (nm) 16.71 16.22 17.03 18.45 Lattice strain (%) 0.71 0.73 0.70 0.65 Space group *F d 3 m F d 3 m F d 3 m F d 3 m* Lattice parameter *a* = *b* = *c* (Å) 8.4030 8.4030 8.4030 8.4030 Volume/Å3 593.34 593.34 593.34 593.34

Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films

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265

where *d* is interplanar spacing, *λ* is X-ray wavelength, and *θ* is reflection angle. The *λ*/4*a2* is a constant and *d* = *h*<sup>2</sup> + *k*<sup>2</sup> + *l* 2 , which is determined by sin<sup>2</sup> *θ* value.

These lattice constants are tabulated in **Table 3**. The lattice parameters of all the ferrite films do not match precisely with the standard JCPDS bulk values which could be attributed to the strains present on the surface of the films during the synthesis [25, 26]. Annealing temperature has a pronounced effect on grain size. The lattice parameter calculated for nickel ferrite thin film is 8.338 Å [27]. This is in accordance with the variation in lattice parameter with Zn content reported for the bulk ferrites.

#### **4.2. Microstructure analysis**

The FESEM images revealed that the Ni0.3Zn0.7Fe2 O4 films have dense and homogenous grains with an average grain size. Film annealed at 400°C was homogenous with dense microstructure, and they have high adhesion to the substrate. Film annealed at 500°C shows a welldeveloped grain. The grains slowly appeared with increasing annealed temperature. This is because the grain tends to combine with closer grain to form larger grain size. The structure formed in the thin films is a normal characteristic of film derived from sol-gel. The average grain sizes of the Ni0.3Zn0.7Fe2 O4 nanoferrite thin films are 18.61 nm (400°C), 26.25 nm (500°C),


**Table 3.** Structural parameters of Ni0.3Zn0.7Fe2 O4 thin films from XRD spectra.

28.12 nm (600°C) and 41.32 nm (700°C). The grains of the films are spherical and uniform, and cohesion of grains is due to the magnetic attraction. The average grain size of the films is presented in **Figure 2**. The histogram of grain size distribution shifted to the larger grain size as the annealing temperature increased. However, based on the coercivity, *Hc* results (Section 4.3), the transition from 600 to 700°C of annealing temperature exhibits a fall of the *Hc* value.

**Figure 3** illustrates the cross section of the samples annealed at 400, 500, 600 and 700°C, respectively. The deposited films were uniform with two cycles of number deposition cycle. It was found that the thin films have thickness in the range of 145.7–285.6 nm which was confirmed by cross-sectional FESEM images. The grain size over 26 nm was further increased with a higher annealing temperature. Accordingly, the number of grain sizes beyond the single domain to multidomain critical size also increased. Therefore, the number of domain wall increased as the movement of domain wall contribution to make ease of magnetization increased [28].

**Figure 4** presented the thickness of the Ni0.7Zn0.3Fe2 O ferrite thin films and shows its relationship with the grain size. Annealing is a process related to secondary grain growth in the film. Thompson discussed the secondary grain growth mechanism and came to the conclusion that the secondary grain growth is driven by the reduction of the total grain boundary energy. Since the grain boundary energy is film-thickness-dependent, the secondary grain growth rate increases when the film thickness is reduced [29].

#### **4.3. Magnetic properties**

where *D* is the crystallite size, *β* is full width at half maximum of the diffraction peak, *λ* is the

The lattice constants of these films were calculated using indexing method [24] given by

where *d* is interplanar spacing, *λ* is X-ray wavelength, and *θ* is reflection angle. The *λ*/4*a2*

These lattice constants are tabulated in **Table 3**. The lattice parameters of all the ferrite films do not match precisely with the standard JCPDS bulk values which could be attributed to the strains present on the surface of the films during the synthesis [25, 26]. Annealing temperature has a pronounced effect on grain size. The lattice parameter calculated for nickel ferrite thin film is 8.338 Å [27]. This is in accordance with the variation in lattice parameter with Zn

O4

with an average grain size. Film annealed at 400°C was homogenous with dense microstructure, and they have high adhesion to the substrate. Film annealed at 500°C shows a welldeveloped grain. The grains slowly appeared with increasing annealed temperature. This is because the grain tends to combine with closer grain to form larger grain size. The structure formed in the thin films is a normal characteristic of film derived from sol-gel. The average

*<sup>d</sup>* <sup>=</sup> sin<sup>2</sup> \_\_\_\_\_\_\_\_ *<sup>θ</sup>* (*h*<sup>2</sup> + *k*<sup>2</sup> + *l*

*θ* value.

ferrite thin films. Reproduced with permission from [21].

nanoferrite thin films are 18.61 nm (400°C), 26.25 nm (500°C),

<sup>2</sup>) (2)

films have dense and homogenous grains

is a

<sup>4</sup> *<sup>a</sup>* <sup>2</sup> <sup>=</sup> sin<sup>2</sup> \_\_\_\_\_*<sup>θ</sup>*

O4

, which is determined by sin<sup>2</sup>

wavelength of 1.54 Å, and *θ* is scanning angle.

2

The FESEM images revealed that the Ni0.3Zn0.7Fe2

O4

\_\_\_*<sup>λ</sup>*

**Figure 1.** XRD pattern of air-annealed Ni0.3Zn0.7Fe2

content reported for the bulk ferrites.

**4.2. Microstructure analysis**

grain sizes of the Ni0.3Zn0.7Fe2

constant and *d* = *h*<sup>2</sup> + *k*<sup>2</sup> + *l*

264 Coatings and Thin-Film Technologies

Eq. (2):

The plots of magnetization, *M*, against magnetic field strength, *H* (*M−H* hysteresis loop), for Ni0.3Zn0.7Fe2 O4 films annealed at 400, 500, 600 and 700°C were shown in **Figure 5**. The hysteresis shape is narrow and has linear loops which have a low saturation magnetization, *M*<sup>s</sup> . The saturation magnetization, *M*<sup>s</sup> , and coercivity, *H*<sup>c</sup> , values have been directly extracted from these curves and have been listed for various annealing temperatures in **Table 4**. The *M*<sup>s</sup> and *H*c could be attributed to the varied grain size and crystallinity. The lower value of saturation

**Figure 3.** Cross-sectional of Ni0.3Zn0.7Fe2

from [21].

O nanoferrite films annealed at different temperatures.

Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films

http://dx.doi.org/10.5772/intechopen.80461

267

**Figure 4.** The average grain size and film thickness as a function of annealed temperatures. Reproduced with permission

**Figure 2.** FESEM images of the Ni0.3Zn0.7Fe2 O nanoferrite films annealed at (a) 400, (b) 500, (c) 600 and (d) 700°C.Reproduced with permission from [21].

magnetization, *M*<sup>s</sup> , Ni0.3Zn0.7Fe2 O4 films (**Figure 6(a)**) could be caused by several reasons. A large grain boundary volume presented in thin films would result in the increase of the *M*<sup>s</sup> [30]. Other reasons for the increase *M*<sup>s</sup> are due to complex spinel structure; it was difficult to gain Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films http://dx.doi.org/10.5772/intechopen.80461 267

**Figure 3.** Cross-sectional of Ni0.3Zn0.7Fe2 O nanoferrite films annealed at different temperatures.

magnetization, *M*<sup>s</sup>

with permission from [21].

266 Coatings and Thin-Film Technologies

, Ni0.3Zn0.7Fe2

Other reasons for the increase *M*<sup>s</sup>

**Figure 2.** FESEM images of the Ni0.3Zn0.7Fe2

O4

grain boundary volume presented in thin films would result in the increase of the *M*<sup>s</sup>

films (**Figure 6(a)**) could be caused by several reasons. A large

O nanoferrite films annealed at (a) 400, (b) 500, (c) 600 and (d) 700°C.Reproduced

are due to complex spinel structure; it was difficult to gain

[30].

**Figure 4.** The average grain size and film thickness as a function of annealed temperatures. Reproduced with permission from [21].

**Figure 5.** (a) Hysteresis loop of Ni0.7Zn0.3Fe2 O ferrite thin films and (b) first quadrant of the magnetic hysteresis loops of the samples. Magnetization at any given field increased with heat treatment temperature. Reproduced with permission from [21].


Ni0.3Zn0.7Fe2

O nanoferrite films with perfect crystallization. The metal cations can occupy either

observed were closed to the reported value of *Hc*

Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films

http://dx.doi.org/10.5772/intechopen.80461

269

) increases with the grain

was 16.54 Oe

which is

O

, was decreased as the annealing tem-

as a function of annealed temperatures of the Ni0.3Zn0.7Fe2

*A* sites (tetrahedral) or *B* sites (octahedral), which will result in a partially disordered cation

size, and the observations on larger decrease are interpreted mostly by oxygen absorption, char-

distribution in the crystal lattice [31]. The saturation magnetization (*M*<sup>s</sup>

perature and average grain size increased (**Figure 6(b)**). The maximum value of *Hc*

and (b) *H*<sup>c</sup>

acteristic to the preparation technique. The coercivity, *H*<sup>c</sup>

for the grain size 26.25 nm. The *H*<sup>c</sup>

**Figure 6.** Comparison of the variation in (a) *M*<sup>s</sup>

nanoferrite thin films. Reproduced with permission from [21].

**Table 4.** Saturation magnetization, *M*<sup>s</sup> , coercivity, *H*<sup>c</sup> , grain size, *D*, and calculated crystallite size from XRD *d*xrd of Ni0.3Zn0.7Fe2 O4 nanoferrite thin films.

**Figure 6.** Comparison of the variation in (a) *M*<sup>s</sup> and (b) *H*<sup>c</sup> as a function of annealed temperatures of the Ni0.3Zn0.7Fe2 O nanoferrite thin films. Reproduced with permission from [21].

**Figure 5.** (a) Hysteresis loop of Ni0.7Zn0.3Fe2

Reproduced with permission from [21].

**Table 4.** Saturation magnetization, *M*<sup>s</sup>

nanoferrite thin films.

**Temperature (°C)** *M***<sup>s</sup>**

268 Coatings and Thin-Film Technologies

from [21].

Ni0.3Zn0.7Fe2

O4

O ferrite thin films and (b) first quadrant of the magnetic hysteresis loops of

 **(±0.001 Oe)** *D* **(±0.1 nm)** *d***xrd (±0.01 nm)**

, grain size, *D*, and calculated crystallite size from XRD *d*xrd of

the samples. Magnetization at any given field increased with heat treatment temperature. Reproduced with permission

 **(±0.01 emu/g)** *H***<sup>c</sup>**

 1.287 16.184 18.6 16.71 2.395 16.536 26.3 16.22 2.653 12.288 28.1 17.03 3.421 8.297 42.32 18.45

, coercivity, *H*<sup>c</sup>

Ni0.3Zn0.7Fe2 O nanoferrite films with perfect crystallization. The metal cations can occupy either *A* sites (tetrahedral) or *B* sites (octahedral), which will result in a partially disordered cation distribution in the crystal lattice [31]. The saturation magnetization (*M*<sup>s</sup> ) increases with the grain size, and the observations on larger decrease are interpreted mostly by oxygen absorption, characteristic to the preparation technique. The coercivity, *H*<sup>c</sup> , was decreased as the annealing temperature and average grain size increased (**Figure 6(b)**). The maximum value of *Hc* was 16.54 Oe for the grain size 26.25 nm. The *H*<sup>c</sup> observed were closed to the reported value of *Hc* which is within the range of 20–210 Oe [20]. The decreases of *Hc* were contributed from the transition of the single domain to the multidomain [32]. The coercivity (*H*<sup>c</sup> ) has a maximum grain size of about 26 nm and a steep decrease at larger grain sizes (41.3 nm). The smaller grain sizes and the decrease of *H*<sup>c</sup> are due to the randomizing effects of thermal energy. Thermal energy has an important role in magnetic instability of single-domain magnetic particles. Due to the smaller grain sizes, the thermal agitation becomes small and will not be able to cause fluctuations in the magnetic spin orientations of the nanoparticles where they freeze in random orientations. The latter is probably due to the decreased anisotropy constant, which leads to a sharp decrease in coercivity according to the random anisotropy model. The relation between the decrease *H*<sup>c</sup> and increase grain size shows the linear inverse proportionality between coercivity, *H*<sup>c</sup>

**Figure 7** demonstrates the curves of absorbance and transmittance, respectively. The absorption spectrum exhibits that NiZn ferrite thin films have low absorbance in visible region and it is close to infrared region (**Figure 7(a)**). However, absorbance in UV region is high. This result of the optical behaviour is analogous to those claimed by [34] or cobalt ferrite thin film using microwave-assisted nonaqueous sol-gel process. Optical transmittance is plotted in a wavelength range of 200–800 nm as shown in **Figure 7(b)**. The films are highly transparent in the visible range below 90%. The average transmittance is calculated and tabulated in **Table 5**. The optical transmittance spectra of annealed thin films show a good transmission in the visible region and a sharp fall in the UV region which corresponds to the band gap. The decrease of the transmittance is due to the interaction of the incident long-wavelength radiation with

**Figure 8(a)** and (**b)** demonstrates the optical band-gap energy direct and indirect of the films

UV–Vis absorbance spectra. Direct and indirect band-gap energies can be obtained from the

line on *x*-axis gives the value of optical band gap. The values of band gap as listed in **Table 5** do change with thickness. A dependence of band-gap energy shift on the grain size is attributed to electron confinement effect related with the grain size in the films. As a result, the observed

film thickness. It has been studied that the band gap does not change significantly with the thickness after the film grows completely [37]. The band gap becomes saturated for a particular value of thickness [38]. The presented values of optical band-gap energy are larger than reported value for NiZn ferrite film 1.66 eV [39] and bulk NiZn ferrite, 1.55–1.66 eV [40]. The direct and indirect band energy increases could be the effect of strain present in the films during heat treatment [41].

> **Direct band-gap energy (α***hv***)**

**(±0.01 eV)**

**2**

 145.7 3.76 3.03 85.0 18.61 180.7 3.66 3.08 78.0 26.25 221.5 3.58 3.16 70.0 28.12 285.6 3.04 3.30 70.0 41.32

**Table 5.** Thickness, band gap, transmittance and grain size of sample at various annealing temperatures.

on *hv*, where *α* is the absorption coefficient, whereas *hv* is the photon energy

Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films

with increasing grain size is due to the decrease of resistivity and the increase of

**Indirect band-gap energy (α***hv***)**

**(±0.01 eV)**

**−1/2**

**Average** 

**(±0.1%)**

**transmittance (T%)** 

**Grain size (±0.01 nm)**

size, *D,* by *H*<sup>c</sup> α 1/*D* [22, 33].

the free electron in the films [35].

2

**Thickness (nm) (±0.1 nm)**

dependencies (*αhv*)

decrease in *Eg*

**Annealing temperature (°C)**

annealed at various temperatures. The band-gap energy (*Eg*

in eV [36]. A linear line was obtained by plotting (*αhv*)

**4.4. Optical properties**

, and grain

271

) of the thin films was calculated from

http://dx.doi.org/10.5772/intechopen.80461

1/n against *hv*. The intersection of this straight

**Figure 7.** (a) Absorbance spectra of thin film at various annealed temperatures and (b) transmittance spectra of NiZn ferrite thin films.

increase grain size shows the linear inverse proportionality between coercivity, *H*<sup>c</sup> , and grain size, *D,* by *H*<sup>c</sup> α 1/*D* [22, 33].

### **4.4. Optical properties**

within the range of 20–210 Oe [20]. The decreases of *Hc*

the decrease of *H*<sup>c</sup>

270 Coatings and Thin-Film Technologies

ferrite thin films.

of the single domain to the multidomain [32]. The coercivity (*H*<sup>c</sup>

about 26 nm and a steep decrease at larger grain sizes (41.3 nm). The smaller grain sizes and

important role in magnetic instability of single-domain magnetic particles. Due to the smaller grain sizes, the thermal agitation becomes small and will not be able to cause fluctuations in the magnetic spin orientations of the nanoparticles where they freeze in random orientations. The latter is probably due to the decreased anisotropy constant, which leads to a sharp decrease in coercivity according to the random anisotropy model. The relation between the decrease *H*<sup>c</sup>

**Figure 7.** (a) Absorbance spectra of thin film at various annealed temperatures and (b) transmittance spectra of NiZn

are due to the randomizing effects of thermal energy. Thermal energy has an

were contributed from the transition

) has a maximum grain size of

and

**Figure 7** demonstrates the curves of absorbance and transmittance, respectively. The absorption spectrum exhibits that NiZn ferrite thin films have low absorbance in visible region and it is close to infrared region (**Figure 7(a)**). However, absorbance in UV region is high. This result of the optical behaviour is analogous to those claimed by [34] or cobalt ferrite thin film using microwave-assisted nonaqueous sol-gel process. Optical transmittance is plotted in a wavelength range of 200–800 nm as shown in **Figure 7(b)**. The films are highly transparent in the visible range below 90%. The average transmittance is calculated and tabulated in **Table 5**. The optical transmittance spectra of annealed thin films show a good transmission in the visible region and a sharp fall in the UV region which corresponds to the band gap. The decrease of the transmittance is due to the interaction of the incident long-wavelength radiation with the free electron in the films [35].

**Figure 8(a)** and (**b)** demonstrates the optical band-gap energy direct and indirect of the films annealed at various temperatures. The band-gap energy (*Eg* ) of the thin films was calculated from UV–Vis absorbance spectra. Direct and indirect band-gap energies can be obtained from the dependencies (*αhv*) 2 on *hv*, where *α* is the absorption coefficient, whereas *hv* is the photon energy in eV [36]. A linear line was obtained by plotting (*αhv*) 1/n against *hv*. The intersection of this straight line on *x*-axis gives the value of optical band gap. The values of band gap as listed in **Table 5** do change with thickness. A dependence of band-gap energy shift on the grain size is attributed to electron confinement effect related with the grain size in the films. As a result, the observed decrease in *Eg* with increasing grain size is due to the decrease of resistivity and the increase of film thickness. It has been studied that the band gap does not change significantly with the thickness after the film grows completely [37]. The band gap becomes saturated for a particular value of thickness [38]. The presented values of optical band-gap energy are larger than reported value for NiZn ferrite film 1.66 eV [39] and bulk NiZn ferrite, 1.55–1.66 eV [40]. The direct and indirect band energy increases could be the effect of strain present in the films during heat treatment [41].


**Table 5.** Thickness, band gap, transmittance and grain size of sample at various annealing temperatures.

Further annealing temperature demonstrated the improvement in the degree of crystallin-

been obtained at room temperature from the hysteresis loops which increases with annealing temperatures. The hysteresis shape shown is narrow and has linear loops which have a

• The micrograph revealed the increasing average grain size with the annealing temperature. The grains of the films are spherical and uniform, and cohesion of grains is due to the

• The absorption spectrum exhibits that NiZn ferrite thin films have low absorbance in visible region and close to infrared region. The films are highly transparent in the visible range below 90%. The optical transmittance spectra of annealed thin films show a good transmission in the visible region and a sharp fall in the UV region which corresponds to the band

The authors are grateful to the Ministry of Higher Education Malaysia (MOHE) and Universiti Putra Malaysia for Research University Grant (vote number 5526200). Reprinted by permission from Springer Nature and Copyright Clearance Center, Springer, Journal of the Australian Ceramic Society, Microstructure and magnetic properties of Ni-Zn ferrite thin film synthesized using sol-gel and spin-coating technique, Yusnita, Y.; Azis, R. S.; Kanagesan,

, Raba'ah Syahidah Azis1,2\* and Muhammad Syazwan Mustaffa<sup>2</sup>

1 Materials Synthesis and Characterization Laboratory, Institute of Advanced Technology,

2 Department of Physics, Faculty of Science, Universiti Putra Malaysia, Selangor, Malaysia

S.; and Bahmanrokh, G., licence number: 4393541427573 (2017).

The authors declare that they have no competing interest.

\*Address all correspondence to: rabaah@upm.edu.my

Universiti Putra Malaysia, Selangor, Malaysia

of the magnetic moments inside the domain is fully controlled by thermal energy.

, of the synthesized Ni0.3Zn0.7Fe2

, decreases as the average grain size increases since the alignment

Spin-Coating Technique for Fabricating Nickel Zinc Nanoferrite (Ni0.3Zn0.7Fe2O4) Thin Films

O4

http://dx.doi.org/10.5772/intechopen.80461

ferrite thin films has

273

ity of the annealed films.

magnetic attraction.

**Acknowledgements**

**Conflict of interest**

**Author details**

Yusnita Yusuf<sup>1</sup>

low *M*<sup>s</sup>

gap.

• The saturation magnetization, *M*<sup>s</sup>

. The coercivity, *Hc*

**Figure 8.** (a) Plot of (*αhv*) 2 as a function of photon energy (eV) and (b) Plot of (*αhv*) 1/2 as a function of photon energy (hv).

### **5. Summary**

NiZn ferrite thin film with composition Ni0.3Zn0.7Fe2 O4 was successfully prepared using solgel spin-coating technique. The structure formed in the thin films is a normal characteristic of film derived from sol-gel. Sol-gel spin-coating method was able to produce the similar trend and behaviour, among others, of ferrite thin film. The results are summarized as follows:

• The phase analysis of films produced the complete phase with the formation of spinel structures of Ni0.3Zn0.7Fe2 O4 ferrite which were observable at annealed 400°C and upwards. Further annealing temperature demonstrated the improvement in the degree of crystallinity of the annealed films.


### **Acknowledgements**

The authors are grateful to the Ministry of Higher Education Malaysia (MOHE) and Universiti Putra Malaysia for Research University Grant (vote number 5526200). Reprinted by permission from Springer Nature and Copyright Clearance Center, Springer, Journal of the Australian Ceramic Society, Microstructure and magnetic properties of Ni-Zn ferrite thin film synthesized using sol-gel and spin-coating technique, Yusnita, Y.; Azis, R. S.; Kanagesan, S.; and Bahmanrokh, G., licence number: 4393541427573 (2017).

### **Conflict of interest**

The authors declare that they have no competing interest.

### **Author details**

**5. Summary**

**Figure 8.** (a) Plot of (*αhv*)

272 Coatings and Thin-Film Technologies

structures of Ni0.3Zn0.7Fe2

NiZn ferrite thin film with composition Ni0.3Zn0.7Fe2

2

O4

O4

ferrite which were observable at annealed 400°C and upwards.

gel spin-coating technique. The structure formed in the thin films is a normal characteristic of film derived from sol-gel. Sol-gel spin-coating method was able to produce the similar trend and behaviour, among others, of ferrite thin film. The results are summarized as follows:

as a function of photon energy (eV) and (b) Plot of (*αhv*)

• The phase analysis of films produced the complete phase with the formation of spinel

was successfully prepared using sol-

1/2 as a function of photon energy (hv).

Yusnita Yusuf<sup>1</sup> , Raba'ah Syahidah Azis1,2\* and Muhammad Syazwan Mustaffa<sup>2</sup>

\*Address all correspondence to: rabaah@upm.edu.my

1 Materials Synthesis and Characterization Laboratory, Institute of Advanced Technology, Universiti Putra Malaysia, Selangor, Malaysia

2 Department of Physics, Faculty of Science, Universiti Putra Malaysia, Selangor, Malaysia

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