Finally, the atomistic model considers that in thin films deposited at high temperatures the thermodynamic parameters control the orientation of the growth, favoring planes with low energy, *i.e.* the (200) planes. This behavior can be understood if adatoms with low diffusion, as in the case of the cation Cl in NaCl, which have three bonds in the (111) surface, and have high diffusion in surface (200), are considered to have only one bond (see Fig. 6), leading to the conclusion that if NaCl thin films are grown at high temperatures, they probably will

On the other hand, the same model establishes that the degree of ion bombardment is influenced by the flow ratio between the density current of the ions and the density current of the atoms, Ji/Ja, and the ion energy, Ei, and both depend on the pressure, substrate target distance, and substrate bias (Ensinger, 1998; Losbichle & Mitterer, 1997). One parameter that combines both the ion energy and the flow has been suggested by Musil and Kadlec (Musil

deformation energy, and (200) planes, with minimum surface energy.

the slope of the straight line represents the deformation per volume.

direction is obtained.

Milman, 2000).

grow along the [200] direction.

$$E\_p = \frac{J\_i}{J\_a} \times E\_i \tag{4}$$

where Ei is incident energy of the ions and is obtained from the plasma potential (Vp) and substrate bias (Vs) and the elementary charge by the expression ( ) *E eV V i s p* .

#### **3.1 Physical mechanisms of thin film growth**

So far, we have studied the influence of the growth parameters on the structure and microstructure of thin films; now we discuss the main physical mechanisms involved in the nucleation and crystallization of the films over the substrate. Studies done through Xray diffraction, optical diffraction and mainly through electron diffraction have allowed establishing that there are three mechanisms of the nucleation and growth of thin films, which depend on the thermodynamic parameters of the deposit and the substrate surface interaction between the adatoms and the substrate material (Green, 1994). The three basic modes are: (i) Volmer–Weber model, (ii) Frank–Van der Merwe model and (iii) Stranski– Krastanov model. A schematic illustration of each of these modes of growth is shown in Fig. 7.

Fig. 7. Modes of growth of film: a- Volmer–Weber island growth b- Frank–vander Merwe layer growth and c- Stranski-Krastanov layer plus island growth (Harsha, 2005).

Thin Film Growth Through Sputtering Technique and Its Applications 409

intersection lines of the crystal side faces and the substrate present a specific structural precondition for the growth of these crystals. The intersection lines can be active or passive in the monolayer nucleation on the side crystal faces. In the first case, the movement of the monolayer growth steps proceeds from the intersection line to the top of the crystal, while in the second case, the movement of the growth steps proceeds in the direction of the intersection line. In the presence of impurities, the direction of the movement of the growth steps will be

Fig. 8. Types of crystal growth in polycrystalline thin films: a- and b- growth of dispersed individual crystals on the substrate surface, the intersection line between the substrate and the side crystal faces are active (a) and passive (b) in the monolayer nucleation; c- and drole of grain boundaries in the growth of crystals as parts of a polycrystalline matrix, pure grain boundary active in the monolayer nucleation (a), contaminated grain boundary

**4. Influence of the normalized temperature (Ts/ Tm) on the structure and**

The structural behavior, as a function of the normalized temperature of Ts/Tm<0.3 is shown in Fig. 9. In general, the patterns recorded from the various films grown present the same reflections as the target (δ-NbN cubic phase), but show important differences in the relative intensities, in particular those corresponding to planes (111) and (200), which increase with an increase of the absolute temperature. Regarding this, it is important to point out that contrary to what occurs in the target and the rest of the film patterns, the XRD pattern of the film grown at 553 K shows a greater intensity for the reflection from plane (111) than from that corresponding to the (200) plane. While the (111) plane in the target pattern has a relative intensity of 100% compared to 86% for the (200) plane, in the 553 K sample the relative intensity of the (200) plane is approximately 6% of the corresponding (111) plane. This result indicates that the supply of additional 22.4 meV to the substrate (which is the energy difference between room temperature and 553 K) does not bring about changes in the polycrystalline character of the materials, but gives place to the growth of films with

passive in the monolayer nucleation (b). (Barna & Adamimik)

**microstructure of thin films**

preferred orientation along the (111) plane.

important in determining the location of the developing second phase, *e.g.*, SCL.

In the Volmer–Weber model, equilibrium exists in a three-dimensional crystal of the film in contact with the substrate, while the rest of the substrate is devoid of any condensed phase. Nucleation of film occurs in the form of discrete three-dimensional nuclei on the surface of the substrate, for example lead on graphite. Both the number of nuclei and the size of a given nucleus increase. Finally, the nuclei grow in size until they intersperse with each other to form a continuous film.

In the Frank and Vander Merwe model, nucleation occurs in the form of a monolayer island of the deposit. Eventually the monolayer's grow together to form a complete continuous monolayer of the deposit. The process repeats itself so that the deposit grows in a layer-bylayer manner, for example rare gases on graphite. In this growth, the interaction between the substrate and the layer atoms is stronger than that between neighboring layer atoms. Layer-by-layer growth is hindered by elastic constraints at the solid–solid interface.

The Stransky and Krastanov (S-K) model combines the features of layer-by-layer growth and discrete three-dimensional nucleation The S-K nucleation is common with metal-onmetal deposition and at low temperatures where the surface mobility is low (Greene, 1987). The conditions for these types of growth are generally described in terms of thermodynamics and surface energy considerations. In this model, nucleation and growth occurs as in the layer-by-layer mode, so that a finite number of monolayer's is produced. Subsequent formation of film occurs by formation of discrete nuclei. The lattice mismatch between the substrate and the deposit cannot be accommodated when the layer thickness increases, so the three-dimensional growth follows the layer-by-layer growth. Alternatively, symmetry or orientation of the overlayers with respect to the substrate might be responsible for the production of this growth mode.

The foregoing models were summarized by Barna and Adamik (Barna & Adamimik, 1998), who established that the growth of the films has the following evolution: nucleation, island growth, coalescence of islands, formation of polycrystalline islands and channels, development of continuous structure, and thickness growth.

According to Barna, the nucleation starting the growth of individual islands takes place on the substrate surface at the very first stage of the condensation (primary nucleation*)* or later on the bare substrate surface area developing upon liquid-like coalescence ( secondary nucleation*).* A peculiar case of nucleation shows up on the surface of a growing crystal when its growth is blocked by a surface covering layer (SCL) of an impurity phase. This is the repeated nucleation*.* The primary nucleation starts the condensation and the film growth on the whole substrate surface simultaneously, while the secondary and the repeated nucleation initiates the start of the growth locally in later stages of film formation. It is important to note that on amorphous substrates the nuclei are randomly oriented.

Crystal growth is the fundamental structure-forming phenomenon which incorporates the depositing material into the condensed phase. Two main cases of crystal growth should be considered in the case of polycrystalline thin films: a- the growth of discrete crystals dispersed on the substrate surface, (Fig. 8a and b), and b- the growth of crystals which are parts of a polycrystalline structure (Fig. 8c and d). Crystals growing from the nuclei are randomly oriented due to the random orientation of the nuclei. The complete coalescence of the crystals touching each other produces a grain coarsening, resulting also in the development of discrete single crystals and is connected to some changes in the orientation controlled mainly by the minimization of the substrate–crystal interface energy. The

In the Volmer–Weber model, equilibrium exists in a three-dimensional crystal of the film in contact with the substrate, while the rest of the substrate is devoid of any condensed phase. Nucleation of film occurs in the form of discrete three-dimensional nuclei on the surface of the substrate, for example lead on graphite. Both the number of nuclei and the size of a given nucleus increase. Finally, the nuclei grow in size until they intersperse with each other

In the Frank and Vander Merwe model, nucleation occurs in the form of a monolayer island of the deposit. Eventually the monolayer's grow together to form a complete continuous monolayer of the deposit. The process repeats itself so that the deposit grows in a layer-bylayer manner, for example rare gases on graphite. In this growth, the interaction between the substrate and the layer atoms is stronger than that between neighboring layer atoms.

The Stransky and Krastanov (S-K) model combines the features of layer-by-layer growth and discrete three-dimensional nucleation The S-K nucleation is common with metal-onmetal deposition and at low temperatures where the surface mobility is low (Greene, 1987). The conditions for these types of growth are generally described in terms of thermodynamics and surface energy considerations. In this model, nucleation and growth occurs as in the layer-by-layer mode, so that a finite number of monolayer's is produced. Subsequent formation of film occurs by formation of discrete nuclei. The lattice mismatch between the substrate and the deposit cannot be accommodated when the layer thickness increases, so the three-dimensional growth follows the layer-by-layer growth. Alternatively, symmetry or orientation of the overlayers with respect to the substrate might be responsible

The foregoing models were summarized by Barna and Adamik (Barna & Adamimik, 1998), who established that the growth of the films has the following evolution: nucleation, island growth, coalescence of islands, formation of polycrystalline islands and channels,

According to Barna, the nucleation starting the growth of individual islands takes place on the substrate surface at the very first stage of the condensation (primary nucleation*)* or later on the bare substrate surface area developing upon liquid-like coalescence ( secondary nucleation*).* A peculiar case of nucleation shows up on the surface of a growing crystal when its growth is blocked by a surface covering layer (SCL) of an impurity phase. This is the repeated nucleation*.* The primary nucleation starts the condensation and the film growth on the whole substrate surface simultaneously, while the secondary and the repeated nucleation initiates the start of the growth locally in later stages of film formation. It is

Crystal growth is the fundamental structure-forming phenomenon which incorporates the depositing material into the condensed phase. Two main cases of crystal growth should be considered in the case of polycrystalline thin films: a- the growth of discrete crystals dispersed on the substrate surface, (Fig. 8a and b), and b- the growth of crystals which are parts of a polycrystalline structure (Fig. 8c and d). Crystals growing from the nuclei are randomly oriented due to the random orientation of the nuclei. The complete coalescence of the crystals touching each other produces a grain coarsening, resulting also in the development of discrete single crystals and is connected to some changes in the orientation controlled mainly by the minimization of the substrate–crystal interface energy. The

important to note that on amorphous substrates the nuclei are randomly oriented.

Layer-by-layer growth is hindered by elastic constraints at the solid–solid interface.

to form a continuous film.

for the production of this growth mode.

development of continuous structure, and thickness growth.

intersection lines of the crystal side faces and the substrate present a specific structural precondition for the growth of these crystals. The intersection lines can be active or passive in the monolayer nucleation on the side crystal faces. In the first case, the movement of the monolayer growth steps proceeds from the intersection line to the top of the crystal, while in the second case, the movement of the growth steps proceeds in the direction of the intersection line. In the presence of impurities, the direction of the movement of the growth steps will be important in determining the location of the developing second phase, *e.g.*, SCL.

Fig. 8. Types of crystal growth in polycrystalline thin films: a- and b- growth of dispersed individual crystals on the substrate surface, the intersection line between the substrate and the side crystal faces are active (a) and passive (b) in the monolayer nucleation; c- and drole of grain boundaries in the growth of crystals as parts of a polycrystalline matrix, pure grain boundary active in the monolayer nucleation (a), contaminated grain boundary passive in the monolayer nucleation (b). (Barna & Adamimik)

## **4. Influence of the normalized temperature (Ts/ Tm) on the structure and microstructure of thin films**

The structural behavior, as a function of the normalized temperature of Ts/Tm<0.3 is shown in Fig. 9. In general, the patterns recorded from the various films grown present the same reflections as the target (δ-NbN cubic phase), but show important differences in the relative intensities, in particular those corresponding to planes (111) and (200), which increase with an increase of the absolute temperature. Regarding this, it is important to point out that contrary to what occurs in the target and the rest of the film patterns, the XRD pattern of the film grown at 553 K shows a greater intensity for the reflection from plane (111) than from that corresponding to the (200) plane. While the (111) plane in the target pattern has a relative intensity of 100% compared to 86% for the (200) plane, in the 553 K sample the relative intensity of the (200) plane is approximately 6% of the corresponding (111) plane. This result indicates that the supply of additional 22.4 meV to the substrate (which is the energy difference between room temperature and 553 K) does not bring about changes in the polycrystalline character of the materials, but gives place to the growth of films with preferred orientation along the (111) plane.

Thin Film Growth Through Sputtering Technique and Its Applications 411

Fig. 10 shows the micro-structural analysis carried out on the NbN films through transmission ion electron microcopy (TEM). Micrographs 10a and 10b correspond to the image obtained in multibeam configuration, and in them one can observe an amorphous matrix with crystalline nanoparticles with spherical morphology that are homogenously distributed along the growth plane of the film. Using Gatan software, it was determined that the average grain size of the nanoparticles was 2.7±0.6 nm with normal distribution (fig. 10c). When doing magnification of nanoparticles of Fig. 10a, we can observe an interferential pattern produced by the atoms belonging to the NbN film (fig. 9d). Higher magnifications of the nanoparticles visualized in Fig. 10a let us observe atomic planes as a product of the interferential pattern produced by the arrangement of atoms belonging to the NbN film (Fig.10d). Image processing of Fig. 10d, using the Fourier transform, lets us visualize the reciprocal space of one NbN nanoparticle (Fig. 10e). Applying a mask (Fig. 10f) over the corresponding diffracting spots and directly measuring the distance between them (0.224 nm), we can confirm that they correspond to distances between (200) planes of δ-NbN. These results confirm the analysis of XRD, which established that NbN grows preferentially

To produce thin films with normalized temperature larger than 0.3 it is necessary to have ultrahigh-vacuum equipment. This condition limits the production the thin films through magnetron sputtering. However, there are studies such as that of Frederick (Frederic & Gall, 2005), who grew CrN thin films on MgO substrate through magnetron sputtering with Ts/tm between 0 .43 to 0.48. These authors found that the films grown at 993 K present complete single-crystal structure with smooth surfaces. The root-mean-square surface roughness for 230-nm-thick layers decreases from 18.8 to 9.3 to 1.1 nm as Ts is raised from

**4.1 Influence of the gas flow on the structure and microstructure of thin films**

In a different set of experiments, nitrogen gas was introduced into the deposition chamber (maintaining the final working pressure constant) in order to study the influence that the addition of this gas during deposition exerts on the structural and micro structural properties of the NbN films. Fig. 11 shows the diffraction patterns recorded from the films grown at 300 W, 553 K and different nitrogen flows. The results obtained make clear that in all cases a preferential growth appears along the (200) plane (Fig.11a). The relative intensity of this diffraction peak is so high that it makes it impossible to distinguish the diffraction peaks corresponding to other planes. To determine the polycrystalline character of the film, we carried out X-ray diffraction experiments at grazing incidence. Fig. 11b shows a representative example. It is clear from Fig. 11 that the films show the same diffraction peaks as the target, confirming their polycrystalline character, although having a preferential orientation along the (200) plane (texture index, 0.65). The grain sizes deduced by Scherer's equation (Cullity, 2001) for the different films grown along the (200) plane vary from 35 nm (Φ= 2 sscm) to 42 nm (Φ = 6 sscm). These results indicate that incorporation of nitrogen during the fabrication process favors the preferential growth of the d-NbN phase along the (200) plane. This effect has also been produced during the growth of TiN films, due to the change of the preferred orientation of the films from the [111] direction to the

The incorporation of nitrogen during the deposition process implies changes in the dynamics of the plasma, since the increase in the number of nitrogen molecules increases the

along the (200) plane.

873 to 973 to 1073 K.

[200] direction when the gas flow is increased.

Fig. 9. XRD pattern of NbN thin films deposited at a normalized temperature (Ts/Tm<0.3) at 300W of power.

Fig. 10.a) and 10.b) TEM micrographs using multibeam configuration of NbN nanoparticles distributed homogeneously; c) average grain size of NbN nanoparticles; d) HRTEM micrographs shows atomic planes; e) reciprocal space of a NbN nanoparticle; and f) distance between (200) planes.

**Ts**

**/Tm<0.3**

**(311)**

**(222)**

**Intensity (a.u)**

300W of power.

between (200) planes.

**(111)**

**(200)**

30 40 50 60 70 80 90

**(220)**

**2**

Fig. 9. XRD pattern of NbN thin films deposited at a normalized temperature (Ts/Tm<0.3) at

a) b) c)

d) e) f)

Fig. 10.a) and 10.b) TEM micrographs using multibeam configuration of NbN nanoparticles

micrographs shows atomic planes; e) reciprocal space of a NbN nanoparticle; and f) distance

distributed homogeneously; c) average grain size of NbN nanoparticles; d) HRTEM

Target

**393K**

**513K**

**553K**

Fig. 10 shows the micro-structural analysis carried out on the NbN films through transmission ion electron microcopy (TEM). Micrographs 10a and 10b correspond to the image obtained in multibeam configuration, and in them one can observe an amorphous matrix with crystalline nanoparticles with spherical morphology that are homogenously distributed along the growth plane of the film. Using Gatan software, it was determined that the average grain size of the nanoparticles was 2.7±0.6 nm with normal distribution (fig. 10c). When doing magnification of nanoparticles of Fig. 10a, we can observe an interferential pattern produced by the atoms belonging to the NbN film (fig. 9d). Higher magnifications of the nanoparticles visualized in Fig. 10a let us observe atomic planes as a product of the interferential pattern produced by the arrangement of atoms belonging to the NbN film (Fig.10d). Image processing of Fig. 10d, using the Fourier transform, lets us visualize the reciprocal space of one NbN nanoparticle (Fig. 10e). Applying a mask (Fig. 10f) over the corresponding diffracting spots and directly measuring the distance between them (0.224 nm), we can confirm that they correspond to distances between (200) planes of δ-NbN. These results confirm the analysis of XRD, which established that NbN grows preferentially along the (200) plane.

To produce thin films with normalized temperature larger than 0.3 it is necessary to have ultrahigh-vacuum equipment. This condition limits the production the thin films through magnetron sputtering. However, there are studies such as that of Frederick (Frederic & Gall, 2005), who grew CrN thin films on MgO substrate through magnetron sputtering with Ts/tm between 0 .43 to 0.48. These authors found that the films grown at 993 K present complete single-crystal structure with smooth surfaces. The root-mean-square surface roughness for 230-nm-thick layers decreases from 18.8 to 9.3 to 1.1 nm as Ts is raised from 873 to 973 to 1073 K.

## **4.1 Influence of the gas flow on the structure and microstructure of thin films**

In a different set of experiments, nitrogen gas was introduced into the deposition chamber (maintaining the final working pressure constant) in order to study the influence that the addition of this gas during deposition exerts on the structural and micro structural properties of the NbN films. Fig. 11 shows the diffraction patterns recorded from the films grown at 300 W, 553 K and different nitrogen flows. The results obtained make clear that in all cases a preferential growth appears along the (200) plane (Fig.11a). The relative intensity of this diffraction peak is so high that it makes it impossible to distinguish the diffraction peaks corresponding to other planes. To determine the polycrystalline character of the film, we carried out X-ray diffraction experiments at grazing incidence. Fig. 11b shows a representative example. It is clear from Fig. 11 that the films show the same diffraction peaks as the target, confirming their polycrystalline character, although having a preferential orientation along the (200) plane (texture index, 0.65). The grain sizes deduced by Scherer's equation (Cullity, 2001) for the different films grown along the (200) plane vary from 35 nm (Φ= 2 sscm) to 42 nm (Φ = 6 sscm). These results indicate that incorporation of nitrogen during the fabrication process favors the preferential growth of the d-NbN phase along the (200) plane. This effect has also been produced during the growth of TiN films, due to the change of the preferred orientation of the films from the [111] direction to the [200] direction when the gas flow is increased.

The incorporation of nitrogen during the deposition process implies changes in the dynamics of the plasma, since the increase in the number of nitrogen molecules increases the

Thin Film Growth Through Sputtering Technique and Its Applications 413

(i = 1–4) admolecule or islands of adatoms (Petrov & Barna, 2003). This can be considered as causing an additional decrease in the (200) surface energy relative to that of the (111) plane. Consequently, the presence of the nitrogen atoms reduces the flow of cations from the (200) to the (111) planes, resulting in the orientation of the growth along the [200] direction. In this example, the growth of the films depends on the final working pressure as well as the

The SEM study (Fig. 12a and 12b) of film growth as a function of the gas flow indicates that the NbN films present a compact granular structure, with a columnar growth of the type described by Movchan and Demchisshim (Movchan & Demchisshim, 1969), having an

average thickness of 0.7 μm, which implies a deposition rate of 20 nm/min.

(a) (b)

(2.0 sccm) and the nitrogen flow was varied.

2001).

Fig. 12. a) Micrograph of the morphology of NbN thin film. b) Micrograph of the crosssection of NbN film, grown at 300 W and 553 K, 20 sccm of Ar and 6 sccm of N2.

An interesting example of the growth of thin films in a reactive phase through the rf sputtering technique is ZrNxOy thin films, which grow at different flow ratios (N2/O2), but with the final working pressure constant (7.4 X10-1 Pa). The results of the XRD studies are shown in figures 13a and 13b, where it is possible to observe the influence that the flows have on the crystallographic of the films. Fig. 13a shows the XRD pattern of the films where the nitrogen flow was maintained constant (2.5 sccm) and the oxygen flow was varied, and Fig. 13b shows the XRD pattern of the films where the oxygen flow was maintained constant

The results obtained allowed establishing that there is an optimum flow ratio of 1.25 in which growth of a film with a high degree of crystallinity is reached. This behavior is very similar to the NbN films discussed above; therefore, the physicochemical mechanisms involved in the growth of thin films are the same for nitrogen molecules as for oxygen molecules, due to the higher reactivity of oxygen as compared to nitrogen (Martin, *et al.*,

The SEM study (Fig. 14a) of ZrNxOy films that were grown at a flow ratio of 1.25 present a highly homogenous and maybe very compact surface in which it is not possible to find contrast, and therefore the growth mechanism is not well defined. The microstructure of the

bombardment energy of the ions.

probability of collisions, promoting a larger number of chemical reactions on the substrate surface. These reactions can be explained using a model of low energy (<20 eV) ion bombardment during film growth at a Ts/Tm ratio ranging between 0.1 and 0.3 (a condition which is satisfied in this case) (Petrov & Barna, 2003). According to this model, 25 eV are sufficient to cause collision dissociation of the N2 ions, providing a continuous source of atomic nitrogen. The nitrogen readily chemisorbs on the (200) planes but not on the Nterminated (111) planes. This in turn reduces the mean free path of the metal cation on the (200) plane due to capture by the nitrogen atoms, and promotes the formation of a NbNi

Fig. 11. a) XRD patterns -2 configuration and b) Grazing angle (2) XRD pattern recorder from the NbN film, grown at 300 W, 553 K, 20 sccm of Ar and 6 sccm of N2.

probability of collisions, promoting a larger number of chemical reactions on the substrate surface. These reactions can be explained using a model of low energy (<20 eV) ion bombardment during film growth at a Ts/Tm ratio ranging between 0.1 and 0.3 (a condition which is satisfied in this case) (Petrov & Barna, 2003). According to this model, 25 eV are sufficient to cause collision dissociation of the N2 ions, providing a continuous source of atomic nitrogen. The nitrogen readily chemisorbs on the (200) planes but not on the Nterminated (111) planes. This in turn reduces the mean free path of the metal cation on the (200) plane due to capture by the nitrogen atoms, and promotes the formation of a NbNi

**a**

=2 sccm

**b**

Target

(222)

(311)

Target

=4 sccm

=6 sccm

30 40 50 60 70 80 90

30 40 50 60 70 80 90

**2** Fig. 11. a) XRD patterns -2 configuration and b) Grazing angle (2) XRD pattern recorder

from the NbN film, grown at 300 W, 553 K, 20 sccm of Ar and 6 sccm of N2.

(220)

**2**

= 6 sccm

**(220)**

**(111)**

(111)

(200)

**(200)**

**Intensity (a.u)**

**Intensity(a.u)**

**(222)**

**(311)**

(i = 1–4) admolecule or islands of adatoms (Petrov & Barna, 2003). This can be considered as causing an additional decrease in the (200) surface energy relative to that of the (111) plane. Consequently, the presence of the nitrogen atoms reduces the flow of cations from the (200) to the (111) planes, resulting in the orientation of the growth along the [200] direction. In this example, the growth of the films depends on the final working pressure as well as the bombardment energy of the ions.

The SEM study (Fig. 12a and 12b) of film growth as a function of the gas flow indicates that the NbN films present a compact granular structure, with a columnar growth of the type described by Movchan and Demchisshim (Movchan & Demchisshim, 1969), having an average thickness of 0.7 μm, which implies a deposition rate of 20 nm/min.

Fig. 12. a) Micrograph of the morphology of NbN thin film. b) Micrograph of the crosssection of NbN film, grown at 300 W and 553 K, 20 sccm of Ar and 6 sccm of N2.

An interesting example of the growth of thin films in a reactive phase through the rf sputtering technique is ZrNxOy thin films, which grow at different flow ratios (N2/O2), but with the final working pressure constant (7.4 X10-1 Pa). The results of the XRD studies are shown in figures 13a and 13b, where it is possible to observe the influence that the flows have on the crystallographic of the films. Fig. 13a shows the XRD pattern of the films where the nitrogen flow was maintained constant (2.5 sccm) and the oxygen flow was varied, and Fig. 13b shows the XRD pattern of the films where the oxygen flow was maintained constant (2.0 sccm) and the nitrogen flow was varied.

The results obtained allowed establishing that there is an optimum flow ratio of 1.25 in which growth of a film with a high degree of crystallinity is reached. This behavior is very similar to the NbN films discussed above; therefore, the physicochemical mechanisms involved in the growth of thin films are the same for nitrogen molecules as for oxygen molecules, due to the higher reactivity of oxygen as compared to nitrogen (Martin, *et al.*, 2001).

The SEM study (Fig. 14a) of ZrNxOy films that were grown at a flow ratio of 1.25 present a highly homogenous and maybe very compact surface in which it is not possible to find contrast, and therefore the growth mechanism is not well defined. The microstructure of the

Thin Film Growth Through Sputtering Technique and Its Applications 415

a)

b) Fig. 14. a- Micrograph of the morphology of ZrNxOy thin film. b- AFM micrograph of the

The XRD patterns of MgO films (Fig. 15a and 15b) recorded from films grown at room temperature shows the influence which the power supply applied to the target has on the structural behavior of thin films. The films grown within the range from 150 to 200W showed amorphous behavior (not shown in the XRD pattern). On the other hand, starting at 250W, all films showed the (200) plane corresponding to magnesium oxide in the FCC phase. In the XRD pattern it can also be observed that the intensity of the (200) plane is so high that it makes it impossible to distinguish the diffraction peaks corresponding to other planes. To determine the polycrystalline character of the film, we carried out X-ray diffraction experiments at grazing incidence. From Fig. 15b it is clear that the films present two main diffraction peaks, (200) and (220), which belong to the target, confirming their polycrystalline character, although

**4.2 Influence of the power on the structure and microstructure of thin films**

showing a preferential orientation along the (200) plane (texture index, 0.85).

ZrNxOy film, deposited at 350 W, 623 K and ΦN2/ΦO2= 1.25

films was evaluated through AFM studies (Fig. 14b). The micrograph reveals that the average size grain was 150 nm with an average rugosity of 5.9 nm.

Fig. 13 . XRD Patterns recorded from ZrNxOy thin films deposited on common glass at 623K and different ratios flow (ΦN2/ΦO2 ). a- ΦN2 constant at 2.5 sccm and b- ΦO2 constant at 2.0 sccm.

films was evaluated through AFM studies (Fig. 14b). The micrograph reveals that the

**a**

**b**

10 20 30 40 50 60

10 20 30 40 50 60

Fig. 13 . XRD Patterns recorded from ZrNxOy thin films deposited on common glass at 623K and different ratios flow (ΦN2/ΦO2 ). a- ΦN2 constant at 2.5 sccm and b- ΦO2 constant at 2.0

2

**(111 )**

**(-111)**

**2**

**(111)**

**(-111)**

**N2 /O2**

> **N2 /O2 =1.5**

**N2 /O2**

**N2 /O2 =1.0**

**=1.25**

**N2 /O2 =1.0**

**N2 /O2**

**=0.83**

**=1.25**

average size grain was 150 nm with an average rugosity of 5.9 nm.

**(100)**

**Intensity (a.u.)**

**Intensity (a.u.)**

sccm.

Fig. 14. a- Micrograph of the morphology of ZrNxOy thin film. b- AFM micrograph of the ZrNxOy film, deposited at 350 W, 623 K and ΦN2/ΦO2= 1.25

### **4.2 Influence of the power on the structure and microstructure of thin films**

The XRD patterns of MgO films (Fig. 15a and 15b) recorded from films grown at room temperature shows the influence which the power supply applied to the target has on the structural behavior of thin films. The films grown within the range from 150 to 200W showed amorphous behavior (not shown in the XRD pattern). On the other hand, starting at 250W, all films showed the (200) plane corresponding to magnesium oxide in the FCC phase. In the XRD pattern it can also be observed that the intensity of the (200) plane is so high that it makes it impossible to distinguish the diffraction peaks corresponding to other planes. To determine the polycrystalline character of the film, we carried out X-ray diffraction experiments at grazing incidence. From Fig. 15b it is clear that the films present two main diffraction peaks, (200) and (220), which belong to the target, confirming their polycrystalline character, although showing a preferential orientation along the (200) plane (texture index, 0.85).

Thin Film Growth Through Sputtering Technique and Its Applications 417

**(111)**

**(101)´**

**Intensity (a.u.)**

at different powers applied to the target.

substrate.

**thin films**

10 20 30 40 50

**2** Fig. 16. XRD Patterns recorded from ZrNxOy thin films deposited on common glass at 623K

Fig. 17. Atomic force micrographs the MgO thin films deposited at 400 W.

The study of the microstructure of thin films was carried out using atomic force microscopy (AFM). The micrograph of Fig. 17 shows that in the area swept by the cantilever point (16μm2) the film has an average roughness of 20 nm and an average grain size on the order of 110 nm. These results established that there is a threshold of power for the growth of nano-structured MgO thin films with high texture without intentional heating of the

**4.3 Influence of the substrate bias voltage (Vs) on the structure and microstructure of**

Fig. 18 shows the influence that the substrate bias voltage has on the growth of the Ti alloy thin films deposited on steel and glass substrates. The films grown on steel at –100V showed, in addition to the diffraction peaks of the substrate material (note that in all the XRD

**350 W**

**300 W 250 W**

**200 W**

Another example that shows the influence of the power supply applied to the target on the crystallization of the films is the growth of ZrNxOy thin films; the films were grown from the Zr target in the reactive phase in an atmosphere of nitrogen and oxygen (ΦN= 2.5 sccm, ΦO=2.0 sccm). In the XRD pattern of Fig. 16, it is possible to determine that there is a threshold power (250W) for producing the crystallization of the film on the substrate. At this power, the film grows with a high degree of crystallization along the (111) plane. This behavior can be explained by the energy model, since increasing the power on the target implies increasing the energy of the ions that are bombarding the substrate and therefore improving the mobility of the adatoms, which produce chemical reactions and atomic grouping along the planes with the higher surface energy, which generates films with a high degree of crystallinity.

Fig. 15. a- XRD patterns recorded from MgO films deposited at 293 K and different power supplies applied to the target b- Grazing angle (3) XRD pattern recorded from a MgO film grown at 400 W. The target XRD pattern is included as reference.

Another example that shows the influence of the power supply applied to the target on the crystallization of the films is the growth of ZrNxOy thin films; the films were grown from the Zr target in the reactive phase in an atmosphere of nitrogen and oxygen (ΦN= 2.5 sccm, ΦO=2.0 sccm). In the XRD pattern of Fig. 16, it is possible to determine that there is a threshold power (250W) for producing the crystallization of the film on the substrate. At this power, the film grows with a high degree of crystallization along the (111) plane. This behavior can be explained by the energy model, since increasing the power on the target implies increasing the energy of the ions that are bombarding the substrate and therefore improving the mobility of the adatoms, which produce chemical reactions and atomic grouping along the planes with the

30 40 50 60 70 80

30 40 50 60 70 80

**2** Fig. 15. a- XRD patterns recorded from MgO films deposited at 293 K and different power supplies applied to the target b- Grazing angle (3) XRD pattern recorded from a MgO film

**2**

**(200)**

**(111)**

grown at 400 W. The target XRD pattern is included as reference.

**(220)**

**(220)**

**a**

**400W 350W 300W 250W MgO**

**b**

400W

MgO Target

higher surface energy, which generates films with a high degree of crystallinity.

**(200)**

**Intensity (a.u.)**

**Intensity (u.a.)**

Fig. 16. XRD Patterns recorded from ZrNxOy thin films deposited on common glass at 623K at different powers applied to the target.

Fig. 17. Atomic force micrographs the MgO thin films deposited at 400 W.

The study of the microstructure of thin films was carried out using atomic force microscopy (AFM). The micrograph of Fig. 17 shows that in the area swept by the cantilever point (16μm2) the film has an average roughness of 20 nm and an average grain size on the order of 110 nm. These results established that there is a threshold of power for the growth of nano-structured MgO thin films with high texture without intentional heating of the substrate.

#### **4.3 Influence of the substrate bias voltage (Vs) on the structure and microstructure of thin films**

Fig. 18 shows the influence that the substrate bias voltage has on the growth of the Ti alloy thin films deposited on steel and glass substrates. The films grown on steel at –100V showed, in addition to the diffraction peaks of the substrate material (note that in all the XRD

Thin Film Growth Through Sputtering Technique and Its Applications 419

films. The films grown on glass at 120V and -200V show a similar behavior, although there are two differences: the two peaks appearing in the XRD pattern of the film grown at -120 V show a different intensity ratio from those of the film grown on steel, and the XRD pattern of the film grown at –200 V shows an additional low-intensity peak at 82.3, which we

In summary, it follows from the XRD data that the substrate bias voltage (Vs) has a great influence on the structure of the deposited films: an increase of the bias voltage promotes the growth of different phases of titanium having well-defined different preferential crystallographic orientations. The increase of Vs is also reflected in a better crystallinity of the deposited films. The data also show that the crystalline phases formed at different Vs are the same in both substrates (see fig 18b). These results are in accord with that established in the foregoing section, where it was indicated that the increasing V*s* improves the density of

the film, which in turn allows obtaining films with a high degree of crystallinity.

Fig. 19. a) Micrograph of a Ti6A14V film grown on steel with bias voltage of – 200 V. (b) Atomic force Micrograph recorded from Ti6Al4V grown on steel at bias voltage of –200V.

a) b)

3

μm

2

0

 0 1

1

2 μm

3

Ǻ

607 1174

The analysis of the microstructure of the Ti6Al4V was made through scenic electron microscopy (SEM) and an atomic force microscope (AFM). Fig. 19a shows the SEM results; the recorded micrograph of the film grown at –200V reveals a film with excellent texture, high homogeneity and denser microstructure with grain refinement produced under the enhanced plasma bombardment, which is induced by the substrate bias voltage. Through Scherer's equation (Cullity, 2001) and using the broadening of the (110) and (110) planes of the Ti -phase, it was found that the average grain size was 13 nm for the (100) plane and 16.5 nm for the (110) plane (Alfonso *et al.*, 2005). Fig. 19b shows the results of the AFM, which indicated that the films grown at -200 V possess a roughness of 20 nm for a scan area

**4.4 Influence of the energy parameter (Ep) on the structure and microstructure of thin**

The X-ray diffraction patterns of the CrNx films deposited at different Ep values are shown in Fig. 20. It was necessary to use two different scales, because as the Ep increased, the intensity of the CrN (200) peak increased by two orders of magnitude and became narrower. At Ep values lower than 30 eV/atom, we observed that the films mainly showed a [200] orientation with traces of the CrN[111] and CrN[220] orientations, while at higher Ep values only the CrN[200]

associate with the (220) reflection of the -Ti phase.

of 3μmx3μm.

**films**

patterns recorded from the films deposited onto steel diffraction peaks from the substrate are still visible), a quite broad peak, which corresponds to the (002) plane of the α- phase of Ti. The XRD pattern corresponding to the films deposited at –120 V present two welldefined peaks at 35.1 (100) and 38.4 (110), which we associate with the (α/)-alloy phase. The film deposited at a bias voltage of -160 V shows a peak at 38.4, which can be associated with the (110) plane of the - phase of Ti. When the bias voltage increases to -200 V, the XRD pattern shows only one high-intensity peak at 38.4(110), which can be associated with the - phase of Ti. It can also be observed that the width of the diffraction peaks is narrower for the films deposited at higher bias voltages, especially for that deposited at -200 V, which suggests that increasing the bias voltage increases the grain size of the deposited

Fig. 18. (a). XRD pattern of the Ti6A14V films deposited by rf magnetron on steel substrates. The substrate appears as a reference. (b) XRD pattern of the films growth on glass and XRD pattern of theTi6A14V in bulk.

patterns recorded from the films deposited onto steel diffraction peaks from the substrate are still visible), a quite broad peak, which corresponds to the (002) plane of the α- phase of Ti. The XRD pattern corresponding to the films deposited at –120 V present two well-

The film deposited at a bias voltage of -160 V shows a peak at 38.4, which can be associated with the (110) plane of the - phase of Ti. When the bias voltage increases to -200 V, the XRD pattern shows only one high-intensity peak at 38.4(110), which can be associated with the - phase of Ti. It can also be observed that the width of the diffraction peaks is narrower for the films deposited at higher bias voltages, especially for that deposited at -200 V, which suggests that increasing the bias voltage increases the grain size of the deposited

30 40 50 60 70 80 90

**2**

40 60 80 100

**2** Fig. 18. (a). XRD pattern of the Ti6A14V films deposited by rf magnetron on steel substrates. The substrate appears as a reference. (b) XRD pattern of the films growth on glass and XRD

**Ti6Al4V Target**

**Bias -200 V**

**Bias -120 V**

(110), which we associate with the (α/)-alloy phase.

a

 **AISI 420 Steel**

**b**

**Bias -100 V**

**Bias -120 V**

**Bias -200 V**

**Bias -160 V**

defined peaks at 35.1 (100) and 38.4

**(002)**

**(110)**

**(100)**

**35,21**

**Intensity (a.u.)**

pattern of theTi6A14V in bulk.

**38,42**

**Intensity (a.u.)**

films. The films grown on glass at 120V and -200V show a similar behavior, although there are two differences: the two peaks appearing in the XRD pattern of the film grown at -120 V show a different intensity ratio from those of the film grown on steel, and the XRD pattern of the film grown at –200 V shows an additional low-intensity peak at 82.3, which we associate with the (220) reflection of the -Ti phase.

In summary, it follows from the XRD data that the substrate bias voltage (Vs) has a great influence on the structure of the deposited films: an increase of the bias voltage promotes the growth of different phases of titanium having well-defined different preferential crystallographic orientations. The increase of Vs is also reflected in a better crystallinity of the deposited films. The data also show that the crystalline phases formed at different Vs are the same in both substrates (see fig 18b). These results are in accord with that established in the foregoing section, where it was indicated that the increasing V*s* improves the density of the film, which in turn allows obtaining films with a high degree of crystallinity.

Fig. 19. a) Micrograph of a Ti6A14V film grown on steel with bias voltage of – 200 V. (b) Atomic force Micrograph recorded from Ti6Al4V grown on steel at bias voltage of –200V.

The analysis of the microstructure of the Ti6Al4V was made through scenic electron microscopy (SEM) and an atomic force microscope (AFM). Fig. 19a shows the SEM results; the recorded micrograph of the film grown at –200V reveals a film with excellent texture, high homogeneity and denser microstructure with grain refinement produced under the enhanced plasma bombardment, which is induced by the substrate bias voltage. Through Scherer's equation (Cullity, 2001) and using the broadening of the (110) and (110) planes of the Ti -phase, it was found that the average grain size was 13 nm for the (100) plane and 16.5 nm for the (110) plane (Alfonso *et al.*, 2005). Fig. 19b shows the results of the AFM, which indicated that the films grown at -200 V possess a roughness of 20 nm for a scan area of 3μmx3μm.

#### **4.4 Influence of the energy parameter (Ep) on the structure and microstructure of thin films**

The X-ray diffraction patterns of the CrNx films deposited at different Ep values are shown in Fig. 20. It was necessary to use two different scales, because as the Ep increased, the intensity of the CrN (200) peak increased by two orders of magnitude and became narrower. At Ep values lower than 30 eV/atom, we observed that the films mainly showed a [200] orientation with traces of the CrN[111] and CrN[220] orientations, while at higher Ep values only the CrN[200]

Thin Film Growth Through Sputtering Technique and Its Applications 421

the residual stress remained low, but the texture of the films changed abruptly and completely from the [111] to the [002] direction. According to the Petrov Model, the key factor was that the energy of the N2 ions should be near 20 eV to promote the dissociation of the ions through collision with the film surface, in this way providing a supply of atomic nitrogen which can chemisorb on [002] oriented grains and later capture metal atoms, resulting in the development of the 〈002〉 texture. In CrN films, it was observed that at the lowest Ep values (Ep<50 eV/atom) with f(200) ~80%, the film texture was equivalent to the transition regime in Petrov *et al.*'s research. The ion energy and ion–atom flows were not sufficiently low to produce pure 〈111〉 texture, as explained above, but were also not sufficiently high to induce 100% 〈002〉. In this transition regime, a competitive growth between the high-trapping (111) plane and the low surface energy (002) plane was established, since as the ds–t increased, the ion energy of the N2 ions approached 20 eV (see Fig. 20). As the Ep value increased, the conditions favored the formation of the 〈002〉 texture. Thus we considered Petrov's theory to be a rather good model to explain the microstructure evolution of the texture in metal nitride films. The competition of the different orientations during growth also affected the microstructure, as observed in Fig. 21a–d. As the ion bombardment increased, the microstructure changed from a non-ordered columnar growth to well-oriented grains that look like fibers coming out from the substrate, *i.e*. equiaxed grains.

Fig. 21. Cross-section SEM images of samples deposited at increasing Ep values (a) 13, (b)

**1 m 1 m**

**1 m**

**1 m**

46, (c) 143 (d) 316 eV/atom, respectively.

**a b**

**c d**

orientation was detected. An estimation of the 〈200〉 texture was obtained comparing the (200) peak intensity (I200) relative to the intensity of all the orientations (Ii) appearing in the XRD pattern. Table 3 summarizes some experimental conditions of the deposition of the CrN films.

Fig. 20. X-ray diffraction patterns of CrNx coatings: (a) low Ep values, (b) high Ep values.

Table 3. Summary of experimental growth conditions of CrN thin films and some of the results: deposition distance ds–t, r.f. bias voltage Vs, deposition rate R, ion energy Ei, ion– atom flow Ji/Ja, energy parameter Ep and texture factor f<200>.

The results of the microstructure behavior of the CrN thin films are in good agreement with that of the theory of Petrov *et al.,* which was explained above. In the case of CrN thin films, increasing the ion energy by applying a higher substrate bias and keeping low ion–atom flows resulted in a film densification and a change in the preferred orientation from [111] to [002] with a subsequent increase in the levels of stress and the incorporation of argon ions. Moreover, by increasing the energy to around 20 eV and working at higher ion–atom ratios,

orientation was detected. An estimation of the 〈200〉 texture was obtained comparing the (200) peak intensity (I200) relative to the intensity of all the orientations (Ii) appearing in the XRD pattern. Table 3 summarizes some experimental conditions of the deposition of the CrN films.

s

CrN (200)

40 60 65

Fig. 20. X-ray diffraction patterns of CrNx coatings: (a) low Ep values, (b) high Ep values.

CrN-3 3 0 0.10 10 1.3 13 0.80 CrN-4 4 0 0.064 16 2 32 0.89 CrN-5 5 0 0.051 12 2.2 26 0.90 CrN-6 6 0 0.035 13.5 2.6 35 0.92 CrN-7 7 0 0.026 14.5 2.7 39 0.93 CrN-8 8 0 0.019 17 2.7 46 0.92 CrN-3-100 3 -100 0.10 110 1.3 143 0.99 CrN-4-100 4 -100 0.064 115.6 2.0 231 0.99 CrN-5-100 5 -100 0.05 112.3 2.2 247 1.0 CrN-6-100 6 -100 0.034 113.5 2.6 295 1.0 CrN-7-100 7 -100 0.026 114.5 2.7 310 1.0 CrN-8-100 8 -100 0.019 117 2.7 316 1.0 Table 3. Summary of experimental growth conditions of CrN thin films and some of the results: deposition distance ds–t, r.f. bias voltage Vs, deposition rate R, ion energy Ei, ion–

The results of the microstructure behavior of the CrN thin films are in good agreement with that of the theory of Petrov *et al.,* which was explained above. In the case of CrN thin films, increasing the ion energy by applying a higher substrate bias and keeping low ion–atom flows resulted in a film densification and a change in the preferred orientation from [111] to [002] with a subsequent increase in the levels of stress and the incorporation of argon ions. Moreover, by increasing the energy to around 20 eV and working at higher ion–atom ratios,

40 60 65

CrN (220)

Ei (eV) Ji/Ja Ep

2

**b**

(eV/atom)

F<200>

**a**

 Ep (eV/atom) 13 39 27

 Ep(eV/atom) 146 243 308

0

Vs (V) R

atom flow Ji/Ja, energy parameter Ep and texture factor f<200>.

(μm/min)

1x104

4x104

5,0x102 1,0x103 1,5x103

0,0

CrN (111)

Intensity (a.u.)

Sample d s-t

(cm)

the residual stress remained low, but the texture of the films changed abruptly and completely from the [111] to the [002] direction. According to the Petrov Model, the key factor was that the energy of the N2 ions should be near 20 eV to promote the dissociation of the ions through collision with the film surface, in this way providing a supply of atomic nitrogen which can chemisorb on [002] oriented grains and later capture metal atoms, resulting in the development of the 〈002〉 texture. In CrN films, it was observed that at the lowest Ep values (Ep<50 eV/atom) with f(200) ~80%, the film texture was equivalent to the transition regime in Petrov *et al.*'s research. The ion energy and ion–atom flows were not sufficiently low to produce pure 〈111〉 texture, as explained above, but were also not sufficiently high to induce 100% 〈002〉. In this transition regime, a competitive growth between the high-trapping (111) plane and the low surface energy (002) plane was established, since as the ds–t increased, the ion energy of the N2 ions approached 20 eV (see Fig. 20). As the Ep value increased, the conditions favored the formation of the 〈002〉 texture. Thus we considered Petrov's theory to be a rather good model to explain the microstructure evolution of the texture in metal nitride films. The competition of the different orientations during growth also affected the microstructure, as observed in Fig. 21a–d. As the ion bombardment increased, the microstructure changed from a non-ordered columnar growth to well-oriented grains that look like fibers coming out from the substrate, *i.e*. equiaxed grains.

Fig. 21. Cross-section SEM images of samples deposited at increasing Ep values (a) 13, (b) 46, (c) 143 (d) 316 eV/atom, respectively.

Thin Film Growth Through Sputtering Technique and Its Applications 423

observed for the passive oxide layer on a Ti6Al4V alloy. The O 1s peak shows three different contributions: a main one at 530.2 eV that can be associated with metal-oxygen bonds, and two much less intense contributions at 531.7 eV and 532.8 eV, which correspond to the presence of OH- groups and chemisorbed water, respectively. The fit of the V 2p spectrum is complicated by the fact that the V peaks overlap strongly with the O 1s K3,4 X-ray satellite peaks. Therefore, the whole spectrum was refined considering the presence of these satellites and the presence of vanadium peaks. The results show the presence of an intense peak at 515.4 eV, which we associate with the presence of VO2, and a less intense vanadium contribution at 516.7 eV, which corresponds to V2O5. This finding contrasts with those of Milosev (Milosev *et al.*, 2000), who did not find oxidized vanadium in the passive layer

These results show that the films that have been grown through rf magnetron sputtering reproduced the stiochoimetry of the target, although on the surface of the thin films

The elemental composition of thin films is determined through energy dispersive X-ray analysis (EDX), although is important to indicate that elements such as oxygen and nitrogen are not possible to assess because the binding energies are very near to the binding energy of the carbon. As an example, the EDX results for NbN nanoparticle are shown. Fig. 23 shows the elements that constitute the nanoparticles of NbN (N and Nb) that formed the NbN studied in section 4, and the elements of the substrate (common glass) on which the

chemical reactions will be produced, forming hydroxides and passivation layers.

Fig. 23. EDX spectrum of NbN nanoparticle that integrates the NbN thin films.

formed on bulk Ti6Al4V (Alfonso *et al.*, 2006).

**5.2 Energy dispersive X-ray (EDX) analysis**

thin films were grown. The Cu belongs to the sampling.

## **5. Chemical characterization of thin films**

#### **5.1 X-ray photoelectron spectroscopy (XPS) analysis**

The last 10 nm of the surface of thin films can be analyzed using X-ray photoelectron spectroscopy (XPS). As an example, we show the results of XPS on Ti6Al4V thin films. The spontaneously passive film formed on the deposited Ti6Al4V films upon exposure to the air was studied through XPS. Fig. 4 shows high resolution narrow-scan spectra recorded from the Ti 2p, Al 2p, V 2p and O 1s spectra recorded from one of the samples. All the narrow-scan spectra recorded from the rest of the samples are almost identical to those presented in Fig. 22.

Fig. 22. 2p, Al 2p, V 2p and O 1s XPS narrow-scan spectra recorded from a film deposited at a bias voltage of -140 V.

The Ti 2p spectrum shows several contributions, and is similar to that reported previously by other authors for the spontaneously-formed passive oxide layer on bulk Ti6Al4V alloy. The spectrum is dominated by a major TiO2 contribution (BE Ti 2p3/2=458.4 eV, 82%), and shows smaller Ti (BE Ti 2p3/2=453.3 eV, 5%), TiO (BE Ti 2p3/2=454.8 eV, 9%) and Ti2O3 (BE Ti 2p3/2=456.5 eV, 4%) contributions. The Al 2p spectrum shows a major Al2O3 contribution at 73.6 eV and a minor Al0 contribution at 71.0 eV. Again this spectrum is similar to that

The last 10 nm of the surface of thin films can be analyzed using X-ray photoelectron spectroscopy (XPS). As an example, we show the results of XPS on Ti6Al4V thin films. The spontaneously passive film formed on the deposited Ti6Al4V films upon exposure to the air was studied through XPS. Fig. 4 shows high resolution narrow-scan spectra recorded from the Ti 2p, Al 2p, V 2p and O 1s spectra recorded from one of the samples. All the narrow-scan spectra recorded from the rest of the samples are almost identical to

Fig. 22. 2p, Al 2p, V 2p and O 1s XPS narrow-scan spectra recorded from a film deposited at

The Ti 2p spectrum shows several contributions, and is similar to that reported previously by other authors for the spontaneously-formed passive oxide layer on bulk Ti6Al4V alloy. The spectrum is dominated by a major TiO2 contribution (BE Ti 2p3/2=458.4 eV, 82%), and shows smaller Ti (BE Ti 2p3/2=453.3 eV, 5%), TiO (BE Ti 2p3/2=454.8 eV, 9%) and Ti2O3 (BE Ti 2p3/2=456.5 eV, 4%) contributions. The Al 2p spectrum shows a major Al2O3 contribution at 73.6 eV and a minor Al0 contribution at 71.0 eV. Again this spectrum is similar to that

**5. Chemical characterization of thin films**

those presented in Fig. 22.

a bias voltage of -140 V.

**5.1 X-ray photoelectron spectroscopy (XPS) analysis**

observed for the passive oxide layer on a Ti6Al4V alloy. The O 1s peak shows three different contributions: a main one at 530.2 eV that can be associated with metal-oxygen bonds, and two much less intense contributions at 531.7 eV and 532.8 eV, which correspond to the presence of OH- groups and chemisorbed water, respectively. The fit of the V 2p spectrum is complicated by the fact that the V peaks overlap strongly with the O 1s K3,4 X-ray satellite peaks. Therefore, the whole spectrum was refined considering the presence of these satellites and the presence of vanadium peaks. The results show the presence of an intense peak at 515.4 eV, which we associate with the presence of VO2, and a less intense vanadium contribution at 516.7 eV, which corresponds to V2O5. This finding contrasts with those of Milosev (Milosev *et al.*, 2000), who did not find oxidized vanadium in the passive layer formed on bulk Ti6Al4V (Alfonso *et al.*, 2006).

These results show that the films that have been grown through rf magnetron sputtering reproduced the stiochoimetry of the target, although on the surface of the thin films chemical reactions will be produced, forming hydroxides and passivation layers.

#### **5.2 Energy dispersive X-ray (EDX) analysis**

The elemental composition of thin films is determined through energy dispersive X-ray analysis (EDX), although is important to indicate that elements such as oxygen and nitrogen are not possible to assess because the binding energies are very near to the binding energy of the carbon. As an example, the EDX results for NbN nanoparticle are shown. Fig. 23 shows the elements that constitute the nanoparticles of NbN (N and Nb) that formed the NbN studied in section 4, and the elements of the substrate (common glass) on which the thin films were grown. The Cu belongs to the sampling.

Fig. 23. EDX spectrum of NbN nanoparticle that integrates the NbN thin films.

Thin Film Growth Through Sputtering Technique and Its Applications 425

Films of CrN, TiN, ZrN, TaN and NbN were deposited using an unbalanced magnetron sputtering system with different energy parameters to investigate its effect on some film properties. Fig. 23 shows the main results obtained in these films. In general, it may be observed that the residual compressive stresses, determined using the curvature method and Stoney's equation (Stoney, 1909), and increased with the energy parameter, although for

Fig. 25a shows the microhardness of the substrate–film system. Hardness measurements were made on samples deposited on AISI M2 tool steel. The hardness of the substrate was 800 kg mm−2. The hardness of the films was 2–3 times higher than the substrate hardness. Group IV nitrides generally are harder than those of group V (Hofmann, 1990). Moreover, the data showed that the hardness increased as the energy parameter increased. This could be attributed to the increase in ion bombardment on the substrate surface, increasing adatom mobility and producing denser films, and the effect on hardness of compressive

Fig. 25b presents the wear coefficient, Kwear, which was evaluated using a ball cratering system, which is a micro scale abrasion test. Two-body grooving abrasion was the wear mechanism observed for all coatings (not shown), probably due to the micro-cutting action of abrasive particles that were dragged across the ball, basically remaining fixed to the ball surface during the test (Adachi, K. & Hutchings, 2003). The wear coefficients were all in the 10−6 mm3 N−1m−1 range, and the lowest values were obtained for NbN and TaN films. It may be seen that there was a slight decrease in k wear as the energy parameter increased, but the variation was too small to be considered important. In addition, no clear trend was observed as the unbalance coefficient was increased. This might be a consequence of variations in other film properties, such as the coefficient of friction or the roughness, parameters known

Ceramic films like metallic nitrides on a metal substrate are commonly believed to be immune to corrosion. Fig. 25 c-d shows the potentiodynamic polarization curves for the films deposited on AISI 304 substrate at the two energy parameters and the AISI 304 substrate in an electrolyte of 0.5 M H2SO4 + 0.05 M KSCN. The corrosion resistance of a material in the polarization curve is determined by its ability to retain low current densities as the electric potential increases. Quantitatively, Tafel analysis was used to determine the corrosion potential, Ecorr, while the critical passivation current density, Icrit, was estimated from the maximum anodic dissolution current before passivation. The corrosion behavior of PVD ceramic-coated steels in aqueous solutions has been increasingly investigated in recent years (Vyas *et al*., 2010; Yang *et al*., 2008; Wesley, 2001; Li *et al.*, 2009; Kumar, A. & Kaur, 2009). One of the major drawbacks is the presence of defects, which are associated with the growth process in PVD ceramic coatings. These coating defects (*e.g.* pores) are particularly deleterious for corrosion protection, since they provide direct paths through which the electrolyte can reach the coating/substrate interface, where localized galvanic corrosion occurs due to the difference in the corrosion potential between the coating and the steel. A lot of research has been undertaken to deal with this problem of coating defects (Stansbury& Buchnan, 2000; Kaciulis *et al.*, 2000; Lang & Yu, 2001; Lee *et al.*, 2009; Chou & Huang, 2003) particularly for transition metal nitride coatings, which in general have excellent wear and

some films the variation was very small.

stresses present in these films (Olaya *et al*., 2009).

to exert a strong influence on the wear performance of the surfaces.

corrosion resistance and therefore are widely used in industry.

## **6. Optical characterization of thin films**

An important characteristic of thin films is the optical behavior, since this determines possible industrial applications, which range from transparent coatings to optical filters; optical studies are carried out through measurements of transmittance, absorbance and reflectance. Fig. 24 shows the transmittance percent as a function of the wavelength on MgO thin films that have been grown through rf sputtering with different power supplies applied to the target.

Fig. 24. Transmission spectrum vs wavelength the films the MgO with various strengths of power supply to the target.

The results show (see Fig. 24) that the films have percentages of transparency that range from 84% to 95% in films grown at 150 and 400 W, respectively. These results indicate that films that were grown at a higher power have a high value of transmittance. This behavior is due to the fact that the films are denser and possibly present lower diffraction.

## **7. Applications of thin films**

Surface modification by means of thin film deposition is an important industrial process used to protect basic materials against wear, fatigue, corrosion and many other surfacerelated damage phenomena (Vyas *et al.*, 2010; Yang *et al*., 2008; Wesley, 2001; Li *et al*., 2009; Kumar, A. & Kaur, 2009). The modern methods of plasma-assisted physical vapor deposition techniques provide great flexibility for designing films with specific chemistry and microstructure, leading to coatings with unique properties. Among these, ceramic coatings deposited on metallic substrates have shown excellent improvement of the surface properties, such as a low friction coefficient and a high degree of hardness with associated good wear resistance and also corrosion resistance to aggressive environments; (Yang, *et al*., 2008; Wesley, 2012).

An important characteristic of thin films is the optical behavior, since this determines possible industrial applications, which range from transparent coatings to optical filters; optical studies are carried out through measurements of transmittance, absorbance and reflectance. Fig. 24 shows the transmittance percent as a function of the wavelength on MgO thin films that have been grown through rf sputtering with different power supplies applied

400 500 600 700 800

**Wavelength (nm)**

Fig. 24. Transmission spectrum vs wavelength the films the MgO with various strengths of

The results show (see Fig. 24) that the films have percentages of transparency that range from 84% to 95% in films grown at 150 and 400 W, respectively. These results indicate that films that were grown at a higher power have a high value of transmittance. This behavior is

Surface modification by means of thin film deposition is an important industrial process used to protect basic materials against wear, fatigue, corrosion and many other surfacerelated damage phenomena (Vyas *et al.*, 2010; Yang *et al*., 2008; Wesley, 2001; Li *et al*., 2009; Kumar, A. & Kaur, 2009). The modern methods of plasma-assisted physical vapor deposition techniques provide great flexibility for designing films with specific chemistry and microstructure, leading to coatings with unique properties. Among these, ceramic coatings deposited on metallic substrates have shown excellent improvement of the surface properties, such as a low friction coefficient and a high degree of hardness with associated good wear resistance and also corrosion resistance to aggressive environments; (Yang, *et al*.,

due to the fact that the films are denser and possibly present lower diffraction.

 150 W 200 W 250 W 300 W 350 W 400 W

**6. Optical characterization of thin films**

20

power supply to the target.

**7. Applications of thin films**

2008; Wesley, 2012).

40

60

**Tranmittance (%)**

80

100

to the target.

Films of CrN, TiN, ZrN, TaN and NbN were deposited using an unbalanced magnetron sputtering system with different energy parameters to investigate its effect on some film properties. Fig. 23 shows the main results obtained in these films. In general, it may be observed that the residual compressive stresses, determined using the curvature method and Stoney's equation (Stoney, 1909), and increased with the energy parameter, although for some films the variation was very small.

Fig. 25a shows the microhardness of the substrate–film system. Hardness measurements were made on samples deposited on AISI M2 tool steel. The hardness of the substrate was 800 kg mm−2. The hardness of the films was 2–3 times higher than the substrate hardness. Group IV nitrides generally are harder than those of group V (Hofmann, 1990). Moreover, the data showed that the hardness increased as the energy parameter increased. This could be attributed to the increase in ion bombardment on the substrate surface, increasing adatom mobility and producing denser films, and the effect on hardness of compressive stresses present in these films (Olaya *et al*., 2009).

Fig. 25b presents the wear coefficient, Kwear, which was evaluated using a ball cratering system, which is a micro scale abrasion test. Two-body grooving abrasion was the wear mechanism observed for all coatings (not shown), probably due to the micro-cutting action of abrasive particles that were dragged across the ball, basically remaining fixed to the ball surface during the test (Adachi, K. & Hutchings, 2003). The wear coefficients were all in the 10−6 mm3 N−1m−1 range, and the lowest values were obtained for NbN and TaN films. It may be seen that there was a slight decrease in k wear as the energy parameter increased, but the variation was too small to be considered important. In addition, no clear trend was observed as the unbalance coefficient was increased. This might be a consequence of variations in other film properties, such as the coefficient of friction or the roughness, parameters known to exert a strong influence on the wear performance of the surfaces.

Ceramic films like metallic nitrides on a metal substrate are commonly believed to be immune to corrosion. Fig. 25 c-d shows the potentiodynamic polarization curves for the films deposited on AISI 304 substrate at the two energy parameters and the AISI 304 substrate in an electrolyte of 0.5 M H2SO4 + 0.05 M KSCN. The corrosion resistance of a material in the polarization curve is determined by its ability to retain low current densities as the electric potential increases. Quantitatively, Tafel analysis was used to determine the corrosion potential, Ecorr, while the critical passivation current density, Icrit, was estimated from the maximum anodic dissolution current before passivation. The corrosion behavior of PVD ceramic-coated steels in aqueous solutions has been increasingly investigated in recent years (Vyas *et al*., 2010; Yang *et al*., 2008; Wesley, 2001; Li *et al.*, 2009; Kumar, A. & Kaur, 2009). One of the major drawbacks is the presence of defects, which are associated with the growth process in PVD ceramic coatings. These coating defects (*e.g.* pores) are particularly deleterious for corrosion protection, since they provide direct paths through which the electrolyte can reach the coating/substrate interface, where localized galvanic corrosion occurs due to the difference in the corrosion potential between the coating and the steel. A lot of research has been undertaken to deal with this problem of coating defects (Stansbury& Buchnan, 2000; Kaciulis *et al.*, 2000; Lang & Yu, 2001; Lee *et al.*, 2009; Chou & Huang, 2003) particularly for transition metal nitride coatings, which in general have excellent wear and corrosion resistance and therefore are widely used in industry.

Thin Film Growth Through Sputtering Technique and Its Applications 427

**b**

 TiN CrN ZrN TaN NbN

**c**

 TiN CrN ZrN TaN NbN

0 15 30 45 60 75

Ep (eV/at)

0 15 30 45 60 75

(eV/at)

E p

1600

3,0x10-6

4,0x10-6

5,0x10-6

6,0x10-6

K

wear

7,0x10-6

8,0x10-6

9,0x10-6

2000

HV0.025

2400

The critical passivation current density shows that in contrast to the other properties of the films, there was a decrease in the corrosion resistance of the films as the energy parameter increased with Ep. The critical passivation current density is proportional to the exposed area of the substrate due to the existence of pores or pinholes in the film ( Olaya *et al.*, 2005). The value of Icrit depends on the grain limits that join the columns due to the fact that they can contain vacancies, micropores, pinholes, and possibly microcracks, allowing the diffusion of electrolytes of the corrosive solution toward the substrate and increasing degradation of the coatings.

Therefore, the results reported in Fig. 25 c-d suggest that as the energy parameter increased there was more substrate area exposed at the bottom of the pinholes. These pinholes or defects are usually localized at the grain boundaries, which are defined by the crystal growth process, which consequently models the final film structure. The rise in the corrosion current with Ep was very small for the NbN and TiN films, but significantly higher for TaN, ZrN and CrN. However, the higher ion bombardment induced by the degree of magnetic field created more defects, increasing the residual stresses and therefore deteriorating the film–substrate adhesion and consequently the response of the coatings to the corrosion products, as was visually observed at the end of the corrosion test for the TaN film. This phenomenon can be explained by galvanic coupling produced by a difference in the corrosion potential between coated and uncoated specimens. The potential difference is characterized by anodic dissolution of the substrate material with a high anodic current density at the defect site, leading to an adhesion failure of the coating.

The critical passivation current density shows that in contrast to the other properties of the films, there was a decrease in the corrosion resistance of the films as the energy parameter increased with Ep. The critical passivation current density is proportional to the exposed area of the substrate due to the existence of pores or pinholes in the film ( Olaya *et al.*, 2005). The value of Icrit depends on the grain limits that join the columns due to the fact that they can contain vacancies, micropores, pinholes, and possibly microcracks, allowing the diffusion of electrolytes of the corrosive solution toward the substrate and increasing

Therefore, the results reported in Fig. 25 c-d suggest that as the energy parameter increased there was more substrate area exposed at the bottom of the pinholes. These pinholes or defects are usually localized at the grain boundaries, which are defined by the crystal growth process, which consequently models the final film structure. The rise in the corrosion current with Ep was very small for the NbN and TiN films, but significantly higher for TaN, ZrN and CrN. However, the higher ion bombardment induced by the degree of magnetic field created more defects, increasing the residual stresses and therefore deteriorating the film–substrate adhesion and consequently the response of the coatings to the corrosion products, as was visually observed at the end of the corrosion test for the TaN film. This phenomenon can be explained by galvanic coupling produced by a difference in the corrosion potential between coated and uncoated specimens. The potential difference is characterized by anodic dissolution of the substrate material with a high anodic current density at the defect site, leading to an

0 15 30 45 60 75 90

(eV/at)

E p **a**

 TiN CrN ZrN TaN NbN

degradation of the coatings.

adhesion failure of the coating.

residual Stress (GPa)

Thin Film Growth Through Sputtering Technique and Its Applications 429

growth parameters have on the crystallinity and the micro-structure of the thin films has been discussed, and based on the described models, examples have been provided of the thin films' growth under each of the growth parameters. The fundamental idea of presenting these basic theories is to introduce the science of materials to young researchers in the world, helping them to understand by means of examples the basic concepts that apply to the growth of thin films, with the hope that a better understanding of these theories

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Fig. 25. Summary of mechanical and chemical properties of films deposited through sputtering. a) residual stress b) hardness c) wear coefficient d) critical passivation current and d) corrosion potential

#### **8. Conclusion**

In this chapter, the main physical-chemical models utilized to explain the growth of thin films through the rf sputtering technique have been presented, and the influence that the growth parameters have on the crystallinity and the micro-structure of the thin films has been discussed, and based on the described models, examples have been provided of the thin films' growth under each of the growth parameters. The fundamental idea of presenting these basic theories is to introduce the science of materials to young researchers in the world, helping them to understand by means of examples the basic concepts that apply to the growth of thin films, with the hope that a better understanding of these theories will produce new and better applications.

#### **9. References**

428 Crystallization – Science and Technology

**d**

 TiN CrN ZrN TaN NbN

**e**

 TiN CrN ZrN TaN NbN

0 15 30 45 60 75

0 15 30 45 60 75

(eV/at)

E p

In this chapter, the main physical-chemical models utilized to explain the growth of thin films through the rf sputtering technique have been presented, and the influence that the

Fig. 25. Summary of mechanical and chemical properties of films deposited through sputtering. a) residual stress b) hardness c) wear coefficient d) critical passivation current

(eV/at)

E p

0,1

E

and d) corrosion potential

**8. Conclusion**

corr (mV)

Icrit (mA/cm2

)

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**16**

 *Mexico*

**Crystallization of Ge:Sb:Te Thin Films for**

J. J. Gervacio Arciniega, E. Prokhorov, F. J. Espinoza Beltran and G. Trapaga

Chalcogenide glasses are a chemical compound consisting of at least one chalcogen element, sulphur, selenium, or tellurium, in combination with other elements. These glasses obtained great attention after discovery between 1962 and 1969 by Kolomiets, Eaton, Ovshinsky and Pearson of the S-shape current-voltage characteristic in chalcogenide glasses and the switching phenomenon from high to low resistivity states (Popescu, 2005). In 1968 S. R. Ovshinsky demonstrated very short (about of 10-10 seconds) reversible electrical switching phenomena in Te81Ge15Sb2S2 thin films due to amorphous-crystalline phase transition (Ovshinsky, 1968). This work opened a new area of phase change technology, which now is one of the most important technologies for memory devices and applications in computers,

The first electrical phase-change memory devices used various binary, ternary and quaternary Te-based films with compositions made up of Ge:Te, Si:As:Te, Ge:As:Si:Te, etc., systems (Stand, 2005). The early phase change materials used in optical storage comprised simple alloys based primarily on compositions in the vicinity of the tellurium-germanium eutectic. Antimony was primarily used, although other elements including selenium, arsenic and bismuth were all shown to have beneficial effects. Based on the results obtained from films with Ge15Sb4Te81 composition (Ovshinsky, 1971) the first application of Ge:Sb:Te alloys

In the early 1990s, a second generation of high speed phase change materials based on Ge:Sb:Te alloys was reported by several optical memory research groups (Ohta et al 1989, Yamada et al, 1991, Gonzalez-Hernandez et al, 1992). These alloys have stoichiometric compositions along the GeTe–Sb2Te3 pseudobinary line of phase diagram such as Ge2Sb2Te5,

Ge:Sb:Te stoichiometric alloys have three phases: one amorphous and two crystal structures. The first crystalline phase is the rock salt NaCl-like and the second more stable phase is hexagonal. When the amorphous films are heated, the transition from amorphous to rock salt-like structure occurs at around 120-1700C; subsequent heating transforms this phase into a stable hexagonal structure at temperatures around 200-2500C. The exact transition temperature depends on the composition of the film. The hexagonal structure remains stable over a wide temperature range from 200-2500C to around the melting point (593-630 0C,

CD, DVD, phase-change random access memories, etc.

was reported for phase change rewritable optical disks.

**1. Introduction**

Ge1Sb2Te4 and Ge1Sb4Te7.

**Phase Change Memory Application**

*CINVESTAV, Unidad Queretaro, Juriquilla, Querétaro,*

## **Crystallization of Ge:Sb:Te Thin Films for Phase Change Memory Application**

J. J. Gervacio Arciniega, E. Prokhorov, F. J. Espinoza Beltran and G. Trapaga *CINVESTAV, Unidad Queretaro, Juriquilla, Querétaro, Mexico*

## **1. Introduction**

432 Crystallization – Science and Technology

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Hippler, Optical and chemical characterization of thin TiNx films deposited by DC-

Chalcogenide glasses are a chemical compound consisting of at least one chalcogen element, sulphur, selenium, or tellurium, in combination with other elements. These glasses obtained great attention after discovery between 1962 and 1969 by Kolomiets, Eaton, Ovshinsky and Pearson of the S-shape current-voltage characteristic in chalcogenide glasses and the switching phenomenon from high to low resistivity states (Popescu, 2005). In 1968 S. R. Ovshinsky demonstrated very short (about of 10-10 seconds) reversible electrical switching phenomena in Te81Ge15Sb2S2 thin films due to amorphous-crystalline phase transition (Ovshinsky, 1968). This work opened a new area of phase change technology, which now is one of the most important technologies for memory devices and applications in computers, CD, DVD, phase-change random access memories, etc.

The first electrical phase-change memory devices used various binary, ternary and quaternary Te-based films with compositions made up of Ge:Te, Si:As:Te, Ge:As:Si:Te, etc., systems (Stand, 2005). The early phase change materials used in optical storage comprised simple alloys based primarily on compositions in the vicinity of the tellurium-germanium eutectic. Antimony was primarily used, although other elements including selenium, arsenic and bismuth were all shown to have beneficial effects. Based on the results obtained from films with Ge15Sb4Te81 composition (Ovshinsky, 1971) the first application of Ge:Sb:Te alloys was reported for phase change rewritable optical disks.

In the early 1990s, a second generation of high speed phase change materials based on Ge:Sb:Te alloys was reported by several optical memory research groups (Ohta et al 1989, Yamada et al, 1991, Gonzalez-Hernandez et al, 1992). These alloys have stoichiometric compositions along the GeTe–Sb2Te3 pseudobinary line of phase diagram such as Ge2Sb2Te5, Ge1Sb2Te4 and Ge1Sb4Te7.

Ge:Sb:Te stoichiometric alloys have three phases: one amorphous and two crystal structures. The first crystalline phase is the rock salt NaCl-like and the second more stable phase is hexagonal. When the amorphous films are heated, the transition from amorphous to rock salt-like structure occurs at around 120-1700C; subsequent heating transforms this phase into a stable hexagonal structure at temperatures around 200-2500C. The exact transition temperature depends on the composition of the film. The hexagonal structure remains stable over a wide temperature range from 200-2500C to around the melting point (593-630 0C,

Crystallization of Ge:Sb:Te Thin Films for Phase Change Memory Application 435

Sb2Te3 pseudobinary line (Ge2Sb2Te5 Ge1Sb2Te4 and Ge1Sb4Te7) including GeTe and Sb2Te3 (which lie at the end of GeTe-Sb2Te3 pseudobinary line), taking into account the processes during the incubation time. It is necessary to note that GeTe and Sb2Te3, as well as Ge:Sb:Te, can be used in phase-change data storage. Additionally, the crystallization kinetics of Ge4Sb1Te5 films will be present in this chapter. The structure of this alloy is unclear. In contrast to other Ge:Sb:Te ternary compounds, it does not belong to the GeTe – Sb2Te3 homologous series, although it lies on (or very close to) the GeTe-Sb2Te3 pseudobinary line in the Ge:Sb:Te phase diagram between GeTe and Ge2Sb2Te5. It is suggested (Coombs et al., 1995) that this alloy is a solid solution between GeTe and a compound with a composition close to Ge4Sb1Te5. This material, in contrast to stoichiometric alloys, demonstrates only one amorphous-NaCl-type phase transition, but it has the largest optical (Kato et al., 1999) and electrical (Morales-Sanchez et al., 2005) contrast between the crystalline and amorphous

Necessary to note that as-prepared and melt-quenched Ge:Sb:Te amorphous materials show different crystallization kinetics (Nobukuni et al., 1999, Park et al., 1999, Khulbe et al., 2000, Wei et al., 2003, Kalb et al., 2004, Raoux et al., 2008). It has been proposed in the literature that for melt-quenched amorphous materials may exist: embryos following the condensation and evaporation of embryos to form crystalline clusters (Khulbe et al., 2000), preexisting clusters (Park et al., 1999), crystal nuclei (Wei et al., 2003), sinks and voids after repeated overwriting (Nobukuni et al., 1999), locally ordered regions with structure similar to that of crystalline Sb (Naito et al., 2004) or a crystalline amorphous border (Raoux et al., 2008). In spite of different models, crystallization in the melt-quenched amorphous material can start from these nucleation centers, which decrease incubation time and increase crystallization speed. But in

Studies of crystallization kinetics of phase change materials are mostly analyzed using the Johnson–Mehl–Avrami-Kolmogorov (JMAK) model for isothermal annealing (see for example: (Weidenhof et al., 2001, Ruitenberg et al., 2002, Trappe et al., 2000, Morales-Sanchez et al., 2010)), which permitted to determine the activation energy for the crystallization process. This method will be used for investigation of the crystallization

Figure 1 shows in-situ optical reflection (using a laser diode emitting at 650 nm) as a function of temperature with a heating rate of 50C/min for all investigated films with composition indicated on the graph. An abrupt change in reflection is associated with the onset of phase crystallization. Stoichiometric Ge2Sb2Te5, Ge1Sb2Te4 and Ge1Sb4Te7 materials demonstrate two changes: the first corresponds to amorphous-NaCl-type transition and the second to NaCl-type-hexagonal transformation. In contrast, GeTe, Ge4Sb1Te5 and Sb2Te3 show only one amorphous-crystalline transition. The highest crystallization temperature (Tc) has been observed in GeTe, the lowest in Sb2Te3. The crystallization kinetics of the GeTe-Sb2Te3 pseudobinary line will be analyzed according to this Figure from the highest to the

An important aspect to develop thin films of phase change materials is the deposition method. It is well known from the literature that deposition technology strongly affects

this chapter only crystallization of as-prepared materials will be discussed.

states, compared with other Ge:Sb:Te ternary alloys.

kinetics in the compositions mentioned above.

lowest crystallization temperature.

**2. Deposition of phase change films**

depending on composition) (Gonzalez-Hernandez et al, 1992). The principle of phasechange memory operation is based on a reversible phase-transformation from the amorphous (high resistance/low reflectivity state) to crystalline NaCl-type (low resistance/high reflectivity state) under short laser or electrical pulses. It is necessary to note that for phase-change memory applications only the amorphous-NaCl-type transition has been used, probably because the heating time produced by the laser or electrical pulse is too short to form the stable hexagonal structure (Yamada, 1991). This transformation is always accompanied by abrupt changes in reflection (about 20-30 %) and resistivity (about 3 orders of magnitude) in chalcogenide films. The reversible transformation from crystalline to the amorphous structurally disordered state can be obtained by increasing the local temperature of the Ge:Sb:Te layer above its melting point by short intense laser, or electrical pulses and the subsequent quenching with a cooling rate about 1010 deg/s (Yamada et al, 1991).

The phase transition between the amorphous and NaCl-type crystalline phase in stoichiometric materials is fast because atoms in the amorphous state do not need to travel long distances to take their position in the crystal lattice. On the other hand, nonstoichiometric compounds require long-range diffusion when they crystallize from an amorphous state (Yamada et al, 1991, Yamada 1996), which in general slows down the crystallization process. For example, 30 at. % of the deviation from Ge2Sb2Te5 stoichiometric composition increased the crystallization time from 220 to 500 ns (Kyrsta et al., 2001).

For several years, Ge:Sb:Te alloys along the pseudobinary GeTe–Sb2Te3 line have been used as the main material for optical phase-change data storage devices and are currently being investigated for nonvolatile electronic storage purposes (which utilize the difference in the electrical resistance of the two phases) due to the high reflectivity and resistivity contrast between the amorphous and crystalline phases. Many investigations have been carried out on this type of phase-change materials with the purpose of improving their properties and to increase its storage capacity, stability, speed and versatility. These investigations have shown that the crystallization of the phase-change (or amorphous-NaCl-type transition) can be considered as a rate-limiting process to obtain a fast data transfer. That is why there are a considerable number of experimental and theoretical studies investigating the amorphousto-crystalline NaCl-type phase transformation. In spite of the large number of publications, the activation energy for crystallization of Ge:Sb:Te materials reported in the literature shows a large discrepancy; the activation energy of crystallization for Ge2Sb2Te5 reported in the literature varied between 0.8 and 2.9 eV, for example. Such dispersion in the kinetic parameters might depend on the differences in the deposition methods, type of substrate, dielectric cover layer, film thickness and/or parameters of deposition processes, which lead to differences in the crystallization temperature, activation energy, and the like. In addition, experimental data show relatively long incubation times for crystallization of Ge:Sb:Te films and a relatively large amount of crystallized material during this period of time (Morales-Sanchez et al 2010). Incubation time manifests itself as the time necessary to reach critical values of nucleus for the crystallization to occur (Senkader et al, 2004). The existence of a non-negligible amount of crystallite during the incubation time could also be responsible for the dispersion in the kinetic parameters reported in the literature.

Based on such antecedents, we aimed in this chapter at investigating the crystallization kinetics of thin films on a glass substrate obtained by the same deposition process (DC sputtering), with the same thickness (around 200 nm), and composition along the GeTe–

depending on composition) (Gonzalez-Hernandez et al, 1992). The principle of phasechange memory operation is based on a reversible phase-transformation from the amorphous (high resistance/low reflectivity state) to crystalline NaCl-type (low resistance/high reflectivity state) under short laser or electrical pulses. It is necessary to note that for phase-change memory applications only the amorphous-NaCl-type transition has been used, probably because the heating time produced by the laser or electrical pulse is too short to form the stable hexagonal structure (Yamada, 1991). This transformation is always accompanied by abrupt changes in reflection (about 20-30 %) and resistivity (about 3 orders of magnitude) in chalcogenide films. The reversible transformation from crystalline to the amorphous structurally disordered state can be obtained by increasing the local temperature of the Ge:Sb:Te layer above its melting point by short intense laser, or electrical pulses and

the subsequent quenching with a cooling rate about 1010 deg/s (Yamada et al, 1991).

composition increased the crystallization time from 220 to 500 ns (Kyrsta et al., 2001).

the dispersion in the kinetic parameters reported in the literature.

The phase transition between the amorphous and NaCl-type crystalline phase in stoichiometric materials is fast because atoms in the amorphous state do not need to travel long distances to take their position in the crystal lattice. On the other hand, nonstoichiometric compounds require long-range diffusion when they crystallize from an amorphous state (Yamada et al, 1991, Yamada 1996), which in general slows down the crystallization process. For example, 30 at. % of the deviation from Ge2Sb2Te5 stoichiometric

For several years, Ge:Sb:Te alloys along the pseudobinary GeTe–Sb2Te3 line have been used as the main material for optical phase-change data storage devices and are currently being investigated for nonvolatile electronic storage purposes (which utilize the difference in the electrical resistance of the two phases) due to the high reflectivity and resistivity contrast between the amorphous and crystalline phases. Many investigations have been carried out on this type of phase-change materials with the purpose of improving their properties and to increase its storage capacity, stability, speed and versatility. These investigations have shown that the crystallization of the phase-change (or amorphous-NaCl-type transition) can be considered as a rate-limiting process to obtain a fast data transfer. That is why there are a considerable number of experimental and theoretical studies investigating the amorphousto-crystalline NaCl-type phase transformation. In spite of the large number of publications, the activation energy for crystallization of Ge:Sb:Te materials reported in the literature shows a large discrepancy; the activation energy of crystallization for Ge2Sb2Te5 reported in the literature varied between 0.8 and 2.9 eV, for example. Such dispersion in the kinetic parameters might depend on the differences in the deposition methods, type of substrate, dielectric cover layer, film thickness and/or parameters of deposition processes, which lead to differences in the crystallization temperature, activation energy, and the like. In addition, experimental data show relatively long incubation times for crystallization of Ge:Sb:Te films and a relatively large amount of crystallized material during this period of time (Morales-Sanchez et al 2010). Incubation time manifests itself as the time necessary to reach critical values of nucleus for the crystallization to occur (Senkader et al, 2004). The existence of a non-negligible amount of crystallite during the incubation time could also be responsible for

Based on such antecedents, we aimed in this chapter at investigating the crystallization kinetics of thin films on a glass substrate obtained by the same deposition process (DC sputtering), with the same thickness (around 200 nm), and composition along the GeTe–

Sb2Te3 pseudobinary line (Ge2Sb2Te5 Ge1Sb2Te4 and Ge1Sb4Te7) including GeTe and Sb2Te3 (which lie at the end of GeTe-Sb2Te3 pseudobinary line), taking into account the processes during the incubation time. It is necessary to note that GeTe and Sb2Te3, as well as Ge:Sb:Te, can be used in phase-change data storage. Additionally, the crystallization kinetics of Ge4Sb1Te5 films will be present in this chapter. The structure of this alloy is unclear. In contrast to other Ge:Sb:Te ternary compounds, it does not belong to the GeTe – Sb2Te3 homologous series, although it lies on (or very close to) the GeTe-Sb2Te3 pseudobinary line in the Ge:Sb:Te phase diagram between GeTe and Ge2Sb2Te5. It is suggested (Coombs et al., 1995) that this alloy is a solid solution between GeTe and a compound with a composition close to Ge4Sb1Te5. This material, in contrast to stoichiometric alloys, demonstrates only one amorphous-NaCl-type phase transition, but it has the largest optical (Kato et al., 1999) and electrical (Morales-Sanchez et al., 2005) contrast between the crystalline and amorphous states, compared with other Ge:Sb:Te ternary alloys.

Necessary to note that as-prepared and melt-quenched Ge:Sb:Te amorphous materials show different crystallization kinetics (Nobukuni et al., 1999, Park et al., 1999, Khulbe et al., 2000, Wei et al., 2003, Kalb et al., 2004, Raoux et al., 2008). It has been proposed in the literature that for melt-quenched amorphous materials may exist: embryos following the condensation and evaporation of embryos to form crystalline clusters (Khulbe et al., 2000), preexisting clusters (Park et al., 1999), crystal nuclei (Wei et al., 2003), sinks and voids after repeated overwriting (Nobukuni et al., 1999), locally ordered regions with structure similar to that of crystalline Sb (Naito et al., 2004) or a crystalline amorphous border (Raoux et al., 2008). In spite of different models, crystallization in the melt-quenched amorphous material can start from these nucleation centers, which decrease incubation time and increase crystallization speed. But in this chapter only crystallization of as-prepared materials will be discussed.

Studies of crystallization kinetics of phase change materials are mostly analyzed using the Johnson–Mehl–Avrami-Kolmogorov (JMAK) model for isothermal annealing (see for example: (Weidenhof et al., 2001, Ruitenberg et al., 2002, Trappe et al., 2000, Morales-Sanchez et al., 2010)), which permitted to determine the activation energy for the crystallization process. This method will be used for investigation of the crystallization kinetics in the compositions mentioned above.

Figure 1 shows in-situ optical reflection (using a laser diode emitting at 650 nm) as a function of temperature with a heating rate of 50C/min for all investigated films with composition indicated on the graph. An abrupt change in reflection is associated with the onset of phase crystallization. Stoichiometric Ge2Sb2Te5, Ge1Sb2Te4 and Ge1Sb4Te7 materials demonstrate two changes: the first corresponds to amorphous-NaCl-type transition and the second to NaCl-type-hexagonal transformation. In contrast, GeTe, Ge4Sb1Te5 and Sb2Te3 show only one amorphous-crystalline transition. The highest crystallization temperature (Tc) has been observed in GeTe, the lowest in Sb2Te3. The crystallization kinetics of the GeTe-Sb2Te3 pseudobinary line will be analyzed according to this Figure from the highest to the lowest crystallization temperature.

## **2. Deposition of phase change films**

An important aspect to develop thin films of phase change materials is the deposition method. It is well known from the literature that deposition technology strongly affects

Crystallization of Ge:Sb:Te Thin Films for Phase Change Memory Application 437

The thin films have been prepared on glass and Si substrates by DC magnetron sputtering from one composite target. The deposition conditions were: 4.5x10-6 mbar base pressure, 2.0x10-5 mbar work pressure, 180 cm3/min Ar gas input flow, 0.49 W/cm2 DC power density, and 10 min deposition time. Before films sputtering, the targets were cleaned for 10 minutes by sputtering to remove the oxide from the surface. According to XRD measurements all films were obtained in amorphous state. Energy dispersive spectroscopy has shown that deviations of films compositions from the targets were approximately 2%. Figure 2a shows typical surface topography obtained from AFM measurements of asprepared amorphous thin films. All amorphous films demonstrated vary smooth surface with the surface roughness between 1 and 3 nm. In contrast, in the crystalline films (annealed to the temperature of 2000C) the topography image shows grains with average size between 20-120 nm (Fig. 2b). The dimensions of grains are dependent on the films

GeTe is a promising candidate for the application of in phase-change technology due to its higher crystallization temperature (onset of crystallization temperature at about 1840C) when compared to other alloys, since it offers a significant improvement in data-retention at high temperature (Perniola et al., 2010, Fantini et al., 2010). Amorphous films, in contrast to Ge:Sb:Te ternary alloys, crystallize in the rhombohedral phase. This phase can be viewed as a rock salt structure distorted along the [111]-direction (Caravati et al., 2010). Additionally, the central atom is displaced along the [111]-direction from the center of the rhombohedron. Upon crystallization, in a rhombohedral (distorted rock salt) structure appeared about 10% of vacancies occurring on Ge sites (Kolobov 2004). GeTe is a classical ferroelectric material that shows displacive type ferroelectric-paraelectric transition from a rhombohedral ferroelectric phase to a rock salt type structure with paraelectric properties. Two sublattices form this rock salt structure: Ge atoms compose one of the lattices and tellurium atoms the

The NaCl type crystalline structure of GeTe is an unstable phase. This crystalline structure exists for temperatures above 300°C. The ferroelectric-paraelectric transition occurs in the interval of 327-4270C depending on various factors, such as exact composition, carrier density, etc. For example, changes in Te composition from 50 to 50.7 at. % shift the transition

Activation energy of amorphous to rhombohedral phase crystallization reported in the literature demonstrates large dispersion: 1.7 eV (Libera et al., 1993), 1.77 eV (Lu et al., 1995), 1.96 eV (Fantini et al., 2010), 2.5 eV (Fan et al., 2004), 3.9 eV (Matsushita et al., 1989). Such dispersion can be dependent on the differences in the deposition methods, type of substrate, method of results interpretation, etc. According to the literature and to our measurements (Fig. 1), GeTe has the highest crystallization temperature for materials on the GeTe-Sb2Te3 pseudobinary line and must demonstrate the highest activation energy of crystallization.

Figure 3 shows X-ray patterns of GeTe film measured at the temperature indicated on the graph. At temperatures below 1800C, the material shows only wide bands, which are characteristic of amorphous materials. At higher temperatures, a rhombohedral GeTe phase appears. At a temperature of 3500C two phases have been observed: rhombohedral and

composition.

**3. Crystallization of GeTe**

other (Rabe et al., 1987).

temperature from 4270C to 3560C (Okura, 1992).

the microstructure and as a results the optical, electrical and crystallization (especially the crystallization temperature) properties of phase-change materials. For example, the crystallization temperature of the films with the same composition but obtained by different deposition methods can differ for more than 200C, and no suitable explanation has yet been proposed (Morales-Sanchez et al., 2005). That is why, to compare crystallization properties all studied films have to be obtained with the same deposition process (DC sputtering), with the same thickness (around 200 nm), and composition along the GeTe–Sb2Te3 pseudobinary line. This method has been chosen because sputtering can produce films with the desired stoichiometry. Compared, for example, with thermal evaporation, in sputtering the thickness and chemical composition of thin films are easier to control.

Fig. 1. Dependencies of optical reflection on temperature for films with compositions indicated on the graph.

Fig. 2. Surface topography of as-prepare amorphous (a) and crystalline (b) Ge2Sb2Te5 films.

The thin films have been prepared on glass and Si substrates by DC magnetron sputtering from one composite target. The deposition conditions were: 4.5x10-6 mbar base pressure, 2.0x10-5 mbar work pressure, 180 cm3/min Ar gas input flow, 0.49 W/cm2 DC power density, and 10 min deposition time. Before films sputtering, the targets were cleaned for 10 minutes by sputtering to remove the oxide from the surface. According to XRD measurements all films were obtained in amorphous state. Energy dispersive spectroscopy has shown that deviations of films compositions from the targets were approximately 2%.

Figure 2a shows typical surface topography obtained from AFM measurements of asprepared amorphous thin films. All amorphous films demonstrated vary smooth surface with the surface roughness between 1 and 3 nm. In contrast, in the crystalline films (annealed to the temperature of 2000C) the topography image shows grains with average size between 20-120 nm (Fig. 2b). The dimensions of grains are dependent on the films composition.

## **3. Crystallization of GeTe**

436 Crystallization – Science and Technology

the microstructure and as a results the optical, electrical and crystallization (especially the crystallization temperature) properties of phase-change materials. For example, the crystallization temperature of the films with the same composition but obtained by different deposition methods can differ for more than 200C, and no suitable explanation has yet been proposed (Morales-Sanchez et al., 2005). That is why, to compare crystallization properties all studied films have to be obtained with the same deposition process (DC sputtering), with the same thickness (around 200 nm), and composition along the GeTe–Sb2Te3 pseudobinary line. This method has been chosen because sputtering can produce films with the desired stoichiometry. Compared, for example, with thermal evaporation, in sputtering the thickness and chemical composition of thin films are easier

Fig. 1. Dependencies of optical reflection on temperature for films with compositions

a b

Fig. 2. Surface topography of as-prepare amorphous (a) and crystalline (b) Ge2Sb2Te5 films.

to control.

indicated on the graph.

GeTe is a promising candidate for the application of in phase-change technology due to its higher crystallization temperature (onset of crystallization temperature at about 1840C) when compared to other alloys, since it offers a significant improvement in data-retention at high temperature (Perniola et al., 2010, Fantini et al., 2010). Amorphous films, in contrast to Ge:Sb:Te ternary alloys, crystallize in the rhombohedral phase. This phase can be viewed as a rock salt structure distorted along the [111]-direction (Caravati et al., 2010). Additionally, the central atom is displaced along the [111]-direction from the center of the rhombohedron. Upon crystallization, in a rhombohedral (distorted rock salt) structure appeared about 10% of vacancies occurring on Ge sites (Kolobov 2004). GeTe is a classical ferroelectric material that shows displacive type ferroelectric-paraelectric transition from a rhombohedral ferroelectric phase to a rock salt type structure with paraelectric properties. Two sublattices form this rock salt structure: Ge atoms compose one of the lattices and tellurium atoms the other (Rabe et al., 1987).

The NaCl type crystalline structure of GeTe is an unstable phase. This crystalline structure exists for temperatures above 300°C. The ferroelectric-paraelectric transition occurs in the interval of 327-4270C depending on various factors, such as exact composition, carrier density, etc. For example, changes in Te composition from 50 to 50.7 at. % shift the transition temperature from 4270C to 3560C (Okura, 1992).

Activation energy of amorphous to rhombohedral phase crystallization reported in the literature demonstrates large dispersion: 1.7 eV (Libera et al., 1993), 1.77 eV (Lu et al., 1995), 1.96 eV (Fantini et al., 2010), 2.5 eV (Fan et al., 2004), 3.9 eV (Matsushita et al., 1989). Such dispersion can be dependent on the differences in the deposition methods, type of substrate, method of results interpretation, etc. According to the literature and to our measurements (Fig. 1), GeTe has the highest crystallization temperature for materials on the GeTe-Sb2Te3 pseudobinary line and must demonstrate the highest activation energy of crystallization.

Figure 3 shows X-ray patterns of GeTe film measured at the temperature indicated on the graph. At temperatures below 1800C, the material shows only wide bands, which are characteristic of amorphous materials. At higher temperatures, a rhombohedral GeTe phase appears. At a temperature of 3500C two phases have been observed: rhombohedral and

Crystallization of Ge:Sb:Te Thin Films for Phase Change Memory Application 439

homogeneity of the system, random and uniform nucleation, and nucleation rate taking place at the very beginning of the transformation is time independent and applied for

Fig. 4. Dependencies of the volume fraction *x* versus time obtained for GeTe samples from

The value of activation energy for GeTe obtained using Equation (2) was very high: close to 12 eV. Such high value of activation energy has been obtained because GeTe, as other Ge:Sb:Te materials, shows long incubation times *τ* for crystallization, namely, the annealing time required to reach a critical nuclei size or to observe an abrupt increase in the crystalline volume fraction. In this case, the nucleation rate cannot be considered to be time independent for the entire crystallization process as is considered in classical JMAK models

In the case of GeTe, the volume fraction of crystalline phase approximately equals zero during all the incubation time (see Fig. 4). In such cases than the volume fraction of crystalline phase does not demonstrate a substantial increase during the incubation time it is possible to define the beginning of the transformation after the incubation time (Weidenhof

where *τ* is the incubation time. According to Equation (3) the plot *ln[-ln(1- x)]* versus *ln(t-τ)* must be a straight line with slope *n*. Figure 5 shows such modification of an Avrami plot for

The calculated activation energy for GeTe as determined from the modification of the Avrami Equation (3) using the Arrhenius type relation for *K* was 3.98±0.12 eV. The value of the Avrami exponent was about 1.6, which corresponds to diffusion controlled growth from

GeTe. The insert shows the rate constant as a function of the reciprocal temperature.

small dimension grains with decreasing nucleation rate (Christian, 1975).

, (3)

reflection measurements at the temperature indicated on the graph.

et al., 2001). The JMAK equation can now be expressed as:

( ) 1 exp[ ( ) ] *n xt Kt*

isotropic growth at a constant rate.

(Senkader et al., 2004).

NaCl-type, and at higher temperatures the material transforms into a NaCl-type phase. This phase exists only at temperatures above 4000C and in the process of cooling, again transforms into the stable rhombohedral phase.

Crystallization kinetics in GeTe, as in other materials, has been investigated using optical reflection during isothermal measurements. Reflection measurements were made with a 650 nm wavelength laser diode and a PIN diode as detector. In the reflectance measurements, the generally employed assumption has been used; the reflection signals are linearly related to the total transformed crystalline volume fractions *x* (Weidenhof et al., 2001):

$$\mathbf{x} = \left[ R(t) - Ra \right] / \left[ Rc - Ra \right] \,, \tag{1}$$

where R(t), the experimental measurement value of reflectance, Ra and Rc are reflectance of amorphous and fully crystalline phases, respectively.

Fig. 3. XRD patterns of GeTe film measured at the temperature indicated on the graph. Blue patterns correspond to rhombohedral phase. Red patterns correspond to NaCl-type phase.

Figure 4 shows the evolution of the total crystalline volume fraction *x*, for GeTe samples, calculated from reflection measurements for films isothermally annealed at the temperatures indicated on the graph.

In this chapter, crystallization kinetics will be analyzed using the Johnson–Mehl–Avrami-Kolmogorov (JMAK) model for isothermal annealing. According to the classical JMAK model, the transformed volume fraction *x* can be determined by the following expression:

$$\mathbf{x}(t) = 1 - \exp(-Kt^{\eta}) \, , \tag{2}$$

where *K=γexp(-E/kT), γ* is the frequency factor, *E* is the effective activation energy, *t* is the annealing time, and *n* the Avrami exponent, which provides information about the mechanisms of crystallization. The value of *n* can be evaluated from the slope of the *ln[-ln(1- x)]* versus *ln(t)* plot, which in materials with random nucleation and isotropic growth should be linear. The JMAK model is based on several assumptions such as:

NaCl-type, and at higher temperatures the material transforms into a NaCl-type phase. This phase exists only at temperatures above 4000C and in the process of cooling, again

Crystallization kinetics in GeTe, as in other materials, has been investigated using optical reflection during isothermal measurements. Reflection measurements were made with a 650 nm wavelength laser diode and a PIN diode as detector. In the reflectance measurements, the generally employed assumption has been used; the reflection signals are linearly related

where R(t), the experimental measurement value of reflectance, Ra and Rc are reflectance of