Preface

Development of the thin film and coating technologies (TFCT) associated with the develop‐ ment of the solid surface science. The science about solid surface, in turn, is associated with occurrence and development of vacuum technology (late nineteenth to early twentieth cen‐ tury). In the first years of the twentieth century due to the efforts of Irving Langmuir, the science of surface stood out in a special area of research. I. Langmuir was a pioneer in the development of experimental techniques necessary for high-vacuum studies. He introduced the concept of chemical bond adsorption, surface adsorption lattice, and accommodation co‐ efficient; performed the fundamental research of the work function metals, heterogeneous catalysis, and adsorption; and derived the laws of thermionic emission. I. Langmuir was awarded the Nobel Prize for "Outstanding Discoveries in the Field of Surface Chemistry De‐ velopment" in 1932.

Fundamental research on the semiconductor surface in the 1930s focused on the interface between the "metal-semiconductor" (probably this fact is the starting point of the creation of modern thin-film technology). These studies provide the first practical applications of semi‐ conductors in the early 1940s—the selenium rectifiers and detectors on basis lead sulfide, based on the properties of the "metal-semiconductor" interface. Further, the creation of bi‐ polar transistors with point contacts (1949). In the early 1960s, there was the creation of field-effect transistors (FET) based on Si with an inversion layer or structure-based "metaloxide-semiconductor." In FET operation, Si-SiO2 interface plays a crucial role. Then, the in‐ tensive study of surfaces and "film-substrate" interfaces followed, which led to the creation of a huge variety of semiconductor devices. Thus, we can say that TFCT developed in the end of the twentieth century made possible the technological revolution in electronics and through it the revolution in IT and communications.

All variety of TFCT can be easily classified on the thickness of the obtained structures. That is, all multilayer structures can be divided into three main groups: thick (up to 1000 nm), thin (10 to 1000 nm), and ultrathin (below 10 nm).

This book presents the TFCT to create films and coating thickness of less than 1000 nm (thin and ultrathin). There are many methods for producing such films and coatings. They can be divided into the following three groups:


At the end of the twentieth and at the beginning of the twenty-first century, TFCT penetrat‐ ed in many sectors of human life and industry: biology and medicine; nuclear, fusion, and hydrogen energy; protection against corrosion and hydrogen embrittlement; jet engine; space materials science; and many others. Unfortunately, not all of these areas are presented in this book.

Currently, TFCT along with nanotechnology (NT) is the most promising for the develop‐ ment of almost all industries. TFCT and NT interpenetrate and enrich each other. More and more often, published works, which researched and successfully used nanostructured thin films instead of the usual. Evidence of this can be found in some of these book chapters.

All the chapters in the book are divided into three sections. The first section contains papers in which the emphasis is on the use of some features of the methods that had not previously been detected or not used. The second and third sections present recent TFCT achievements in the sectors identified in the titles of these sections.

> **Nikolay N. Nikitenkov** Tomsk Polytechnic University Russia

**Novelties in the Traditional Technologies of Thin Films and Coatings**

At the end of the twentieth and at the beginning of the twenty-first century, TFCT penetrat‐ ed in many sectors of human life and industry: biology and medicine; nuclear, fusion, and hydrogen energy; protection against corrosion and hydrogen embrittlement; jet engine; space materials science; and many others. Unfortunately, not all of these areas are presented

Currently, TFCT along with nanotechnology (NT) is the most promising for the develop‐ ment of almost all industries. TFCT and NT interpenetrate and enrich each other. More and more often, published works, which researched and successfully used nanostructured thin films instead of the usual. Evidence of this can be found in some of these book chapters.

All the chapters in the book are divided into three sections. The first section contains papers in which the emphasis is on the use of some features of the methods that had not previously been detected or not used. The second and third sections present recent TFCT achievements

> **Nikolay N. Nikitenkov** Tomsk Polytechnic University

> > Russia

in the sectors identified in the titles of these sections.

in this book.

X Preface

#### **Molecular Precursor Method for Fabricating** *p***-Type Cu2O and Metallic Cu Thin Films Molecular Precursor Method for Fabricating** *p***-Type Cu<sup>2</sup> O and Metallic Cu Thin Films**

Hiroki Nagai and Mitsunobu Sato Hiroki Nagai and Mitsunobu Sato

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

http://dx.doi.org/10.5772/66476

#### **Abstract**

Functional thin films are used in various fields of our life. Many different methods are used to fabricate these films including physical vapor deposition (PVD) and chemical processes. The chemical processes can be used to manufacture thin films in a relatively cheap way, as compared to PVD methods. This chapter summarizes the procedures of the molecular precursor method (MPM), a chemical process, for fabrication of both metal oxide semiconductor Cu<sup>2</sup> O and metallic Cu thin films by utilizing Cu(II) complexes in coating solutions. The MPM, recently developed and reported by the present authors, represents a facile procedure for thin film fabrication of various metal oxides or phosphates. This method pertinent to the coordination chemistry and materials science including nanoscience and nanotechnology has provided various thin films of high quality. The MPM is based on the design of metal complexes in coating solutions with excellent stability, homogeneity, miscibility, coatability, etc., which are practical advantages. The metal oxides and phosphates are useful as the electron and/or ion conductors, semiconductors, dielectric materials, etc. This chapter will describe the principle and recent achievement, mainly on fabricating the *p*-type Cu<sup>2</sup> O and metallic Cu thin films of the MPM.

**Keywords:** molecular precursor method, thin film, *p*-type Cu<sup>2</sup> O, copper

### **1. Introduction**

A sustainable society requires innovative technology where many disciplines interact. Highly functionalized thin films in various devices such as computers, which were developed mainly for the semiconductor industry in the last century, are now widely used in various fields of our daily life. For example, the touch panel in mobile phone uses a transparent conductive thin film and an antireflection thin film on the glass. A product with various thin films makes

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2017 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

life more comfortable. Many different methods are used to fabricate such thin films including physical vapor deposition (PVD) such as laser ablation, molecular beam epitaxy, sputtering, and chemical processes [1]. The chemical processes can be used to manufacture thin films in a relatively cheap way compared to PVD methods. The chemical process is a processing technique for preparing thin films, ceramic coatings, and powders. However, it is usually difficult to fabricate high-quality thin films using chemical processes.

Transparent metal oxide thin films of *p*-type semiconductors deposited directly on various substrates offer several advantages in the design of heterojunctions with *n*-type oxide films for fabricating thin film devices [2, 3]. Over the past years, the quest to obtain high-quality cuprous oxide films has fueled the development of many physical processing techniques including sputtering, thermal oxidation, vacuum evaporation, molecular beam epitaxy, and electrodeposition [4–9]. However, reproducible formation of cuprous oxide films, uncontaminated by copper or the CuO phase, is a formidable challenge from the technical perspective. Most of the aforementioned physical processing techniques are not economically viable in large area applications. The chemical procedures, which save both energy and resources, have therefore emerged as an attractive alternative because of cost-effective production capabilities.

Recently, we achieved the fabrication of *p*-type Cu<sup>2</sup> O transparent thin films by means of a chemical process, the molecular precursor method (MPM), using the thermal reaction of molecular precursor films spin-coated on a Na-free glass substrate. A stable precursor solution for the spin-coating process was facilely prepared by reacting an isolated Cu(II) complex of ethylenediamine-*N, N, N', N'*-tetraacetic acid (EDTA, H<sup>4</sup> edta) with dibutylamine in ethanol [10, 11]. The MPM is one of the chemical processes used for thin film formation of metal oxides or phosphates [12–22]. The method is based on the formation of excellent precursor films involving anionic metal complexes and the alkylammonium cation. The stability, homogeneity, miscibility, and other characteristics of the precursor solutions, which can be used for various coating methods, are practical advantages, in contrast to the conventional sol-gel method.

This chapter summarizes the procedures used in the MPM for the fabrication of both metal oxide semiconductor Cu<sup>2</sup> O and metallic Cu thin films by utilizing Cu(II) complexes in coating solutions.

### **2. Molecular precursor method**

The decision by the Nobel Prize committee to award the Nobel Prize for chemistry in 1913 to Alfred Werner met with worldwide approval. In a statement, the committee said that Alfred Werner received the prize in recognition of his work on the linkage of atoms in molecules by which he has thrown new light on earlier investigations and opened up new fields of research especially in inorganic chemistry [23]. Today, the metal complexes are used in various applications such as catalysts, luminescence, and medicine. In 1996, one of the authors, M.S., focused on the thin film fabrication of various metal oxides and phosphate compounds using the metal complexes of stable [10–22]. This is the MPM, which is one of the chemical processes used for thin film fabrication. In those days, most of the researchers in the field of thin film formation by chemical processes preferred to use rather unstable metal complexes. It is easy to imagine the capability of polymers to form films because we use polymer films every day. In fact, well-adhered precursor films can be formed on various substrates by coating the solution dispersing the produced oligomers and polymers including metallic species provided by hydrolyzing the unstable metal complexes. These results led us to believe for a long time that only the oligomers and polymers can form precursor films, but the stable metal complexes having discrete molecular weight would not be useful in the formation of such thin films. The MPM was a challenge to this central belief.

life more comfortable. Many different methods are used to fabricate such thin films including physical vapor deposition (PVD) such as laser ablation, molecular beam epitaxy, sputtering, and chemical processes [1]. The chemical processes can be used to manufacture thin films in a relatively cheap way compared to PVD methods. The chemical process is a processing technique for preparing thin films, ceramic coatings, and powders. However, it is usually difficult

Transparent metal oxide thin films of *p*-type semiconductors deposited directly on various substrates offer several advantages in the design of heterojunctions with *n*-type oxide films for fabricating thin film devices [2, 3]. Over the past years, the quest to obtain high-quality cuprous oxide films has fueled the development of many physical processing techniques including sputtering, thermal oxidation, vacuum evaporation, molecular beam epitaxy, and electrodeposition [4–9]. However, reproducible formation of cuprous oxide films, uncontaminated by copper or the CuO phase, is a formidable challenge from the technical perspective. Most of the aforementioned physical processing techniques are not economically viable in large area applications. The chemical procedures, which save both energy and resources, have therefore emerged as an attractive alternative because of cost-effective production

chemical process, the molecular precursor method (MPM), using the thermal reaction of molecular precursor films spin-coated on a Na-free glass substrate. A stable precursor solution for the spin-coating process was facilely prepared by reacting an isolated Cu(II) complex

nol [10, 11]. The MPM is one of the chemical processes used for thin film formation of metal oxides or phosphates [12–22]. The method is based on the formation of excellent precursor films involving anionic metal complexes and the alkylammonium cation. The stability, homogeneity, miscibility, and other characteristics of the precursor solutions, which can be used for various coating methods, are practical advantages, in contrast to the conventional sol-gel

This chapter summarizes the procedures used in the MPM for the fabrication of both metal

The decision by the Nobel Prize committee to award the Nobel Prize for chemistry in 1913 to Alfred Werner met with worldwide approval. In a statement, the committee said that Alfred Werner received the prize in recognition of his work on the linkage of atoms in molecules by which he has thrown new light on earlier investigations and opened up new fields of research especially in inorganic chemistry [23]. Today, the metal complexes are used in various applications such as catalysts, luminescence, and medicine. In 1996, one of the authors, M.S., focused on the thin film fabrication of various metal oxides and phosphate compounds

O and metallic Cu thin films by utilizing Cu(II) complexes in coating

O transparent thin films by means of a

edta) with dibutylamine in etha-

to fabricate high-quality thin films using chemical processes.

4 Modern Technologies for Creating the Thin-film Systems and Coatings

Recently, we achieved the fabrication of *p*-type Cu<sup>2</sup>

of ethylenediamine-*N, N, N', N'*-tetraacetic acid (EDTA, H<sup>4</sup>

capabilities.

method.

solutions.

oxide semiconductor Cu<sup>2</sup>

**2. Molecular precursor method**

The MPM, pertinent to coordination chemistry and materials science including nanoscience and nanotechnology, has been used to fabricate various high-quality thin films with appropriate film thicknesses. As a result, the MPM represents a facile procedure for thin film fabrication of various metal oxides or phosphates, which are useful as electron and/or ion conductors, semiconductors, dielectric materials, etc [24, 25].

**Figure 1** shows the Co<sup>3</sup> O4 thin films, which were first fabricated using the molecular precursor solutions. To date, more than 40 kinds of metal oxides or phosphates have been easily fabricated. **Figure 2** shows the general protocol for fabricating the titanium dioxide thin films. First, a water-resistant coating solution was prepared by the reaction of a neutral [Ti(H<sup>2</sup> O)(edta)] complex with dipropylamine in ethanol, where edta represents ethylenediamine-*N, N, N′, N′*-tetraacetate anion. Molecular precursor solutions can be used in spin, dip, or spray coating on various material surfaces to form precursor thin films. To obtain the precursor film, the coated film was dried at around 70°C. The precursor films involving metal complexes should be amorphous, just as with the metal/organic polymers in the solgel processes; otherwise, it would not be possible to obtain the resulting metal oxide thin films spread homogeneously on substrates by using the following heat treatment. After this, the precursor film was heat-treated at appropriate temperatures for eliminating the components in the organic ligand from the metal complexes involved in the precursor films and to fabricate thin films of crystallized metal oxides or phosphates. The heat treatment of a well-adhered precursor film at 450°C in air is useful for the fabrication of transparent titania thin films.

**Figure 1.** The Co<sup>3</sup> O4 thin films which were first time to fabricate the thin film using the molecular precursor solutions.

**Figure 2.** Protocol for fabricating the thin film using the molecular precursor method.

Cuprous oxide, Cu<sup>2</sup> O, with a cubic structure is a potential candidate for *p*-type semiconductors having a band gap of 2.0 eV, which is the band gap of its single crystal. The thin films of Cu<sup>2</sup> O can be grown generally by dry processes in a vacuum chamber at high temperature, such as sputtering, thermal oxidation, and pulsed laser deposition [26–28]. Over the past years, the quest to obtain high-quality Cu<sup>2</sup> O films has fueled the development of many physical processing techniques including sputtering, thermal oxidation, vacuum evaporation, molecular beam epitaxy, and electrodeposition. However, reproducible formation of Cu<sup>2</sup> O films uncontaminated by the CuO phase is a formidable challenge from the technical perspective [29, 30]. In addition, most of the aforementioned physical processing techniques are not economically viable in large area applications. Solution-based processes have, therefore, emerged as attractive alternatives because of the ability for cost-effective production [31]. Armelao and coworkers have successfully employed a sol-gel solution containing dissolved copper acetate to produce a Cu<sup>2</sup> O thin film [32]. The semiconductive nature of the thin film was, however, unclear owing to the lack of Hall effect measurements in their work.

Recently, *p*-type Cu<sup>2</sup> O transparent thin films were fabricated using the thermal reaction of molecular precursor films spin-coated on a Na-free glass substrate [10]. A stable precursor solution for the spin-coating process was facilely prepared by reacting an isolated Cu(II) complex of EDTA with dibutylamine in ethanol. The 50-nm-thick Cu<sup>2</sup> O thin film resulting from heat treatment of the precursor film at 450°C for 10 min in Ar gas at a flow rate of 1.0 L min−1 was characterized by X-ray diffraction (XRD). **Figure 3** shows the XRD pattern of the resultant thin film deposited on the Na-free glass substrate after heat treating the precursor film. The XRD pattern of the resultant thin film indicated a precise cubic lattice cell parameter with *a* = 0.4265(2) nm, with a crystallite size of 8(2) nm. No additional peaks from any possible contaminants such as Cu and CuO appeared in the XRD pattern of the Cu<sup>2</sup> O film. X-ray Molecular Precursor Method for Fabricating *p*-Type Cu2O and Metallic Cu Thin Films http://dx.doi.org/10.5772/66476 7

Cuprous oxide, Cu<sup>2</sup>

the past years, the quest to obtain high-quality Cu<sup>2</sup>

6 Modern Technologies for Creating the Thin-film Systems and Coatings

**Figure 2.** Protocol for fabricating the thin film using the molecular precursor method.

containing dissolved copper acetate to produce a Cu<sup>2</sup>

plex of EDTA with dibutylamine in ethanol. The 50-nm-thick Cu<sup>2</sup>

films of Cu<sup>2</sup>

mation of Cu<sup>2</sup>

in their work.

Recently, *p*-type Cu<sup>2</sup>

O, with a cubic structure is a potential candidate for *p*-type semicon-

O films has fueled the development

O thin film [32]. The semiconductive

O thin film resulting from

O film. X-ray

O can be grown generally by dry processes in a vacuum chamber at high tem-

O films uncontaminated by the CuO phase is a formidable challenge from the

O transparent thin films were fabricated using the thermal reaction of

ductors having a band gap of 2.0 eV, which is the band gap of its single crystal. The thin

perature, such as sputtering, thermal oxidation, and pulsed laser deposition [26–28]. Over

of many physical processing techniques including sputtering, thermal oxidation, vacuum evaporation, molecular beam epitaxy, and electrodeposition. However, reproducible for-

technical perspective [29, 30]. In addition, most of the aforementioned physical processing techniques are not economically viable in large area applications. Solution-based processes have, therefore, emerged as attractive alternatives because of the ability for cost-effective production [31]. Armelao and coworkers have successfully employed a sol-gel solution

nature of the thin film was, however, unclear owing to the lack of Hall effect measurements

molecular precursor films spin-coated on a Na-free glass substrate [10]. A stable precursor solution for the spin-coating process was facilely prepared by reacting an isolated Cu(II) com-

heat treatment of the precursor film at 450°C for 10 min in Ar gas at a flow rate of 1.0 L min−1 was characterized by X-ray diffraction (XRD). **Figure 3** shows the XRD pattern of the resultant thin film deposited on the Na-free glass substrate after heat treating the precursor film. The XRD pattern of the resultant thin film indicated a precise cubic lattice cell parameter with *a* = 0.4265(2) nm, with a crystallite size of 8(2) nm. No additional peaks from any pos-

sible contaminants such as Cu and CuO appeared in the XRD pattern of the Cu<sup>2</sup>

**Figure 3.** XRD pattern of the thin film adhered to a Na-free glass substrate after heat treatment at 450°C for 10 min in Ar gas at a flow rate of 1.0 L min–1.

 photoelectron spectroscopy (XPS) peaks attributed to the O 1s and Cu 2p3/2 level of the Cu<sup>2</sup> O film were observed at 532.6 and 932.4 eV, respectively. The peak position of the Cu 2p3/2 level is identical to the reported values, which were observed for Cu<sup>2</sup> O thin films prepared by other methods [31, 33]. In addition, no peak for the CuO phase was observed at its typical value of 944 eV [31]. The average grain size of the deposited Cu<sup>2</sup> O particles was ~200 nm, observed via field-emission scanning electron microscopy (FE-SEM). The optical band edge evaluated from the absorption spectrum of the transparent Cu<sup>2</sup> O thin film was 2.3 eV, assuming a direct transition semiconductor. The tensile strength of the films on the glass substrate was measured by a stud pull adhesion test. The tensile strength of the adhesion of the Cu<sup>2</sup> O thin film to the substrate was 83(2) MPa, indicating strong adhesion to the glass substrate.

**Figure 4** shows the Arrhenius plot of the Cu<sup>2</sup> O thin film on the Na-free glass substrate over the temperature range 160–300 K. Hall effect measurements of the thin film indicated that the single phase Cu<sup>2</sup> O thin film is a typical *p*-type semiconductor with a hole concentration of 1.7 × 1016 cm–3 and hole mobility of 4.8 cm<sup>2</sup> V–1 s–1 at ambient temperature. The activation energy from the valence band to the acceptor level was determined as 0.30 eV by least-square fitting using the free carrier concentration spectroscopy (FCCS) method [34].

**Figure 4.** Temperature dependency of the resistivity (― ◊ ―), carrier mobility (…○…), and carrier concentration (… ∆ …) of the thin film. Film was formed by heat treatment at 450°C for 10 min in Ar gas at a flow rate of 1.0 L min–1 on a 10 × 10 × 1.1 mm<sup>3</sup> Na-free glass substrate.

The resistivity of the Cu<sup>2</sup> O thin film fabricated using the MPM was lower than that of both films obtained by the oxidation of a copper film and by the dc reactive magnetron sputtering process [26–28]. It was observed that the carrier concentration tends to be high and the carrier mobility is low for the Cu<sup>2</sup> O thin film fabricated by the MPM, compared to the thin films formed by previously reported processes (**Table 1**).

The method described is the first example of fabrication and characterization of *p*-type Cu<sup>2</sup> O transparent thin films using a coating solution prepared from a starting Cu2+ complex of the H2 edta2− ligand. Reduction of the Cu2+ species occurred unambiguously through the heating process under Ar gas flow. The CuO phase could be found in the thin film by the prolonged reactions after fabricating the Cu<sup>2</sup> O thin film at the identical atmospheric condition and temperature. This result indicates that the formation mechanism of Cu<sup>2</sup> O in the MPM differs from that of the sol-gel method, in which the Cu<sup>2</sup> O phase was derived from the CuO phase in N<sup>2</sup> atmosphere at 90°C. It is notable that the quality of the resultant thin film fabricated


**Table 1.** Comparison of electrical properties of the thin film fabricated herein using a 0.3 mmol g–1 precursor solution under Ar gas at a flow rate of 1.0 L min–1, versus those of the films formed by other methods. All values were measured at 300 K.

by the MPM is excellent, but that fabricated by the CuO reduction was not sufficient for Hall effect measurement.

#### **3. Kinetic study of Cu2 O thin film fabrication**

The resistivity of the Cu<sup>2</sup>

10 × 10 × 1.1 mm<sup>3</sup>

H2

in N<sup>2</sup>

rier mobility is low for the Cu<sup>2</sup>

reactions after fabricating the Cu<sup>2</sup>

formed by previously reported processes (**Table 1**).

Na-free glass substrate.

8 Modern Technologies for Creating the Thin-film Systems and Coatings

from that of the sol-gel method, in which the Cu<sup>2</sup>

O thin film fabricated using the MPM was lower than that of both

O thin film fabricated by the MPM, compared to the thin films

O thin film at the identical atmospheric condition and tem-

O phase was derived from the CuO phase

O

O in the MPM differs

films obtained by the oxidation of a copper film and by the dc reactive magnetron sputtering process [26–28]. It was observed that the carrier concentration tends to be high and the car-

**Figure 4.** Temperature dependency of the resistivity (― ◊ ―), carrier mobility (…○…), and carrier concentration (… ∆ …) of the thin film. Film was formed by heat treatment at 450°C for 10 min in Ar gas at a flow rate of 1.0 L min–1 on a

The method described is the first example of fabrication and characterization of *p*-type Cu<sup>2</sup>

perature. This result indicates that the formation mechanism of Cu<sup>2</sup>

transparent thin films using a coating solution prepared from a starting Cu2+ complex of the

edta2− ligand. Reduction of the Cu2+ species occurred unambiguously through the heating process under Ar gas flow. The CuO phase could be found in the thin film by the prolonged

atmosphere at 90°C. It is notable that the quality of the resultant thin film fabricated

In order to clarify the precise mechanism of Cu<sup>2</sup> O formation from the Cu(II) complex, a kinetic study was performed using XRD [11]. In the study, it was clarified that the thermal reaction of the precursor film, which consists of a dibutylammonium salt of a [Cu(edta)]2− complex ion, first produced metallic Cu species in Ar gas containing <10 ppm of air as an impurity. The Cu phase appeared gradually, and the amount of the phase could be determined from the area of the (111) peak of Cu. The activation energy (1.5 × 10<sup>2</sup> kJ mol–1) of the reduction reaction from the Cu(II) complex to metallic Cu species was obtained by an Arrhenius plot over the temperature range 230–250°C. Above this temperature range, the Cu<sup>2</sup> O phase was formed by the oxidation of the Cu phase under Ar gas flow. The amount of the Cu<sup>2</sup> O phase could be determined from the area of the (111) peak. The activation energy (1.4 × 10<sup>2</sup> kJ mol–1) of Cu<sup>2</sup> O formation from the Cu phase was obtained by the Arrhenius plot over the temperature range 400–450°C. In order to examine the stability of the formed Cu<sup>2</sup> O phase, the oxidation reaction rate from Cu<sup>2</sup> O to the CuO phase in an identical atmosphere was also measured over the temperature range 450–475°C. The activation energy of the oxidation reaction from Cu<sup>2</sup> O to the CuO phase was determined to be 1.0 × 10<sup>2</sup> kJ mol–1. It was observed that the quality of the *p*-type Cu<sup>2</sup> O thin film is strongly dependent on the mechanism of the low-temperature formation.

The XPS spectra of the films are shown in **Figure 5**. The peak positions of the Cu 2p3/2 level in spectra (a)–(d) are 933.5, 932.7, 932.5, and 933.2 eV in **Figure 5**A, respectively. The peaks observed in (a) and (d) can be assigned to the Cu2+ ion, and the broad peak at 944 eV observed in (d) is typical of CuO. In contrast, the peaks observed in (b) and (c) can be assigned to metallic Cu and/or Cu<sup>+</sup> ions. These XPS results are consistent with the XRD results.

**Figure 5.** XPS spectra of (A) Cu 2p3/2, (B) C 1s, and (C) N 1s of the thin films heat-treated in Ar gas flow. The films were fabricated by heat treating the precursor film at (a) 200°C for 10 min, (b) 400°C for 0 min, (c) 450°C for 10 min, and (d) 450°C for 60 min.

In **Figure 5**B and C, the XPS spectra of the C1s and N1s peaks, respectively, are shown. No impurities such as nitrogen or carbon atoms can be found in the XPS spectra of the resultant Cu<sup>2</sup> O thin films, although the metallic Cu<sup>0</sup> thin film includes a certain amount of nitrogen and carbon atoms. It was thus clarified that single phase Cu<sup>2</sup> O formation was completed by the removal of organic residues in the Cu<sup>0</sup> thin film. The co-presence of nitrogen and carbon atoms was thus shown to have an important role in preventing the oxidation of the produced Cu<sup>2</sup> O phase. The presence of nitrogen and carbon atoms may also help in organizing the stepwise reactions.

In the sol-gel method for Cu<sup>2</sup> O thin film formation, the CuO film is annealed at 900°C for 5 h in nitrogen for the partial removal of oxygen atoms from the initial oxide thin film. In contrast, the MPM eliminates the organic components in order to form the Cu<sup>2</sup> O thin film at the abovementioned lower temperature. It is interesting that the difference between these two methods is in the kind of atoms that must be removed. Furthermore, it is important that the formation route of the *p*-type Cu<sup>2</sup> O thin films determines the quality of the thin films, as mentioned above. The MPM is additionally preferable in terms of saving energy by reducing the formation temperature.

It was first shown that the expected Cu<sup>2</sup> O formation using the MPM occurred via an unexpected intermediate Cu<sup>0</sup> phase formed by the thermal decomposition of the molecular precursor involving a Cu(II) complex salt. The XRD measurement of the crystallized thin films was useful in determining the activation energies of the redox reactions from the Cu(II) complex to Cu<sup>0</sup> , from Cu<sup>0</sup> to Cu<sup>2</sup> O, and from Cu<sup>2</sup> O to CuO (**Figure 6**). The redox reactions of the metals and organic ligands occurred stepwise with annealing of the thin films under moderate conditions. Consequently, the ligand in the molecular precursor plays an important role in fabricating excellent *p*-type Cu<sup>2</sup> O thin films. It is also suggested that kinetic studies on the thermal reactions of metal complexes in the solid state are essential for revealing the reaction mechanism of thin film fabrication.

Molecular Precursor Method for Fabricating *p*-Type Cu2O and Metallic Cu Thin Films http://dx.doi.org/10.5772/66476 11

**Figure 6.** A kinetics study was performed in order to clarify the reaction mechanism for fabricating the excellent p-type Cu<sup>2</sup> O thin films by using the molecular precursor method.

### **4. Fabrication of copper thin films**

In **Figure 5**B and C, the XPS spectra of the C1s and N1s peaks, respectively, are shown. No impurities such as nitrogen or carbon atoms can be found in the XPS spectra of the resultant

**Figure 5.** XPS spectra of (A) Cu 2p3/2, (B) C 1s, and (C) N 1s of the thin films heat-treated in Ar gas flow. The films were fabricated by heat treating the precursor film at (a) 200°C for 10 min, (b) 400°C for 0 min, (c) 450°C for 10 min, and (d)

atoms was thus shown to have an important role in preventing the oxidation of the produced

5 h in nitrogen for the partial removal of oxygen atoms from the initial oxide thin film. In

at the abovementioned lower temperature. It is interesting that the difference between these two methods is in the kind of atoms that must be removed. Furthermore, it is important that

mentioned above. The MPM is additionally preferable in terms of saving energy by reducing

sor involving a Cu(II) complex salt. The XRD measurement of the crystallized thin films was useful in determining the activation energies of the redox reactions from the Cu(II) complex

als and organic ligands occurred stepwise with annealing of the thin films under moderate conditions. Consequently, the ligand in the molecular precursor plays an important role in

thermal reactions of metal complexes in the solid state are essential for revealing the reaction

contrast, the MPM eliminates the organic components in order to form the Cu<sup>2</sup>

O phase. The presence of nitrogen and carbon atoms may also help in organizing the

thin film includes a certain amount of nitrogen

thin film. The co-presence of nitrogen and carbon

O thin film formation, the CuO film is annealed at 900°C for

O thin films determines the quality of the thin films, as

O formation using the MPM occurred via an unex-

O to CuO (**Figure 6**). The redox reactions of the met-

O thin films. It is also suggested that kinetic studies on the

phase formed by the thermal decomposition of the molecular precur-

O formation was completed by

O thin film

Cu<sup>2</sup>

450°C for 60 min.

Cu<sup>2</sup>

to Cu<sup>0</sup>

stepwise reactions.

In the sol-gel method for Cu<sup>2</sup>

the formation temperature.

pected intermediate Cu<sup>0</sup>

, from Cu<sup>0</sup>

the formation route of the *p*-type Cu<sup>2</sup>

It was first shown that the expected Cu<sup>2</sup>

to Cu<sup>2</sup>

fabricating excellent *p*-type Cu<sup>2</sup>

mechanism of thin film fabrication.

O, and from Cu<sup>2</sup>

O thin films, although the metallic Cu<sup>0</sup>

10 Modern Technologies for Creating the Thin-film Systems and Coatings

the removal of organic residues in the Cu<sup>0</sup>

and carbon atoms. It was thus clarified that single phase Cu<sup>2</sup>

From the kinetic study of the Cu<sup>2</sup> O thin film formation, it was elucidated that the Cu<sup>0</sup> species formed as an intermediate was oxidized to the resultant Cu<sup>2</sup> O thin film during the heat treatment, and the oxidizing agent is the oxygen present in the commercially available Ar gas as an impurity (<2 ppm) [11]. However, the intermediate Cu thin film obtained through the reaction using the precursor film is not electrically conductive. Therefore, in order to fabricate transparent metal copper thin films, we examined novel precursor solutions [35, 36]. A novel precursor solution containing a Cu2+ complex of EDTA and a Cu2+ complex of propylamine derived from formic acid, and the amine was prepared by mixing the two precursor solutions. The concentration of total copper in the ethanolic precursor solution was adjusted to 0.35 mmol g–1. The spin-coating method was used for precursor film formation on a Na-free glass substrate. The spin-coated precursor films were preheated in a drying oven at 70°C for 10 min and then, heat-treated at 350°C for 15 min under an Ar gas flow of 1.5 L min–1 to fabricate thin films in a tubular furnace with a quartz glass tube. The resultant thin film is hereby denoted as **A**. The rate of temperature increase was controlled by a proportional-integral-derivative program preinstalled in the furnace. Before increasing the temperature, the tubular furnace was filled with Ar gas. The thickness of the resultant films was measured using a stylus profilometer. A flat- and same-sized quartz glass plate was placed on the resulting thin film **A** in the tubular furnace and then post-annealed at 350°C for 20 and 40 min in an Ar gas flow of 1.5 L min–1. The resulting thin film is hereby denoted as **APn** (n = post-annealing time). The XRD patterns of the resultant thin films **A, AP20**, and **AP40** with a thickness of 40 nm over the 2θ range 30°–50° are shown in **Figure 7**. The peaks at 2θ = 36.6° and 42.5° for **A** can be assigned to the (111) and (200) phases of Cu<sup>2</sup> O, respectively, and an additional peak at 43.5° for **A** is assigned to the (111) phase of copper (JCPDS card No. 04–0836). The peak at 2θ = 36.9° for **AP20** is assigned to the (111) phase of Cu<sup>2</sup> O and that at 43.7° can be assigned to the (111) phase of copper.

**Figure 7.** XRD patterns of A, AP20, and AP40 thin films. Parallel beam optics (θ = 0.3°) was employed for calculations. The lines indicate the following: ……A, ----AP20, and ━ AP40. The peaks are denoted as follows: ▽ Cu; | Cu<sup>2</sup> O.

The single peak at 2θ = 43.7° for **AP40** is attributed the (111) phase of copper. Thus, the Cu<sup>2</sup> O phase decreased gradually with increasing post-annealing time, and no peak owing to any impurity phase such as Cu<sup>2</sup> O and CuO could be detected in the annealed **AP40** film. The cell parameter of the Cu cubic lattice in **A**, which was determined by the Wilson & Pike method, is *a* = 3.71(3) Å, and the crystallite size of Cu crystals formed in the film can be determined as 11(1) nm; the estimated standard deviations are presented in parentheses. The cell parameter of the Cu cubic lattice in **AP40** is *a* = 3.72(3) Å, and the crystallite size of the Cu crystals can be determined as 13(1) nm. The surface morphology of the **A** and **AP40** thin films was observed using FE-SEM. The grain size of the Cu particles also increased from 50 nm (A) to 70 nm (AP40) upon annealing. **Figure 8** shows the Auger spectra of **A, AP20**, and **AP40** thin films. Broad peaks were observed at 264 eV for carbon, 509 eV for oxygen, and 764, 835, and 914 eV for copper atoms. The kinetic energies of the copper atoms are identical to those in other films fabricated by the MPM. The result indicates that the amount of carbon atoms in the **AP40** thin film was reduced to half by the post-annealing treatment of **A**. **Figure 9** shows

and then, heat-treated at 350°C for 15 min under an Ar gas flow of 1.5 L min–1 to fabricate thin films in a tubular furnace with a quartz glass tube. The resultant thin film is hereby denoted as **A**. The rate of temperature increase was controlled by a proportional-integral-derivative program preinstalled in the furnace. Before increasing the temperature, the tubular furnace was filled with Ar gas. The thickness of the resultant films was measured using a stylus profilometer. A flat- and same-sized quartz glass plate was placed on the resulting thin film **A** in the tubular furnace and then post-annealed at 350°C for 20 and 40 min in an Ar gas flow of 1.5 L min–1. The resulting thin film is hereby denoted as **APn** (n = post-annealing time). The XRD patterns of the resultant thin films **A, AP20**, and **AP40** with a thickness of 40 nm over the 2θ range 30°–50° are shown in **Figure 7**. The peaks at 2θ = 36.6° and 42.5° for **A** can be

for **A** is assigned to the (111) phase of copper (JCPDS card No. 04–0836). The peak at 2θ = 36.9°

**Figure 7.** XRD patterns of A, AP20, and AP40 thin films. Parallel beam optics (θ = 0.3°) was employed for calculations. The lines indicate the following: ……A, ----AP20, and ━ AP40. The peaks are denoted as follows: ▽ Cu; | Cu<sup>2</sup>

O, respectively, and an additional peak at 43.5°

O and that at 43.7° can be assigned to the (111)

O.

assigned to the (111) and (200) phases of Cu<sup>2</sup>

12 Modern Technologies for Creating the Thin-film Systems and Coatings

for **AP20** is assigned to the (111) phase of Cu<sup>2</sup>

phase of copper.

**Figure 8.** The Auger spectra of A and A20, and AP40 thin films. The lines indicate the following: ……A, ----AP20, and ━AP40, respectively.

**Figure 9.** The curve fitting results of carbon state in AP40 thin film. The lines indicate the following: ……A, —Carbon (neutral), -·-·-·Carbon (tetravalent), respectively.

the curve fitting results of the Auger spectrum for **A** over the range 230–300 eV, corresponding to the carbon atoms, along with the curves of the neutral and tetravalent carbon atoms. The curve fitting results suggest that the carbon atom in **A** was 85% agreement to the neutral. The electrical resistivity of the **A, AP20**, and **AP40** thin films is 7.5 × 10–4, 2.8 × 10–4, and 4.7 × 10–5 Ω cm, respectively. Thus, highly conductive translucent copper thin films could be obtained in commercially available Ar gas. A plausible scheme for copper lattice formation, which can be deduced from the XRD and Auger spectra, is presented in **Scheme 1**.

**Scheme 1.** Plausible scheme for the formation of a Cu thin film from the precursor film.

The scheme indicates that four Cu complexes are required to construct one FCC copper unit cell. During the heat treatment of the precursor complexes in Ar gas flow containing <2 ppm of oxygen as impurity, neighboring complexes react with each other. The valency of copper was reduced from +2 to 0 by the thermal decomposition of the complexes of EDTA and butylamine ligands in Ar gas. In the process, Cu<sup>2</sup> O involving Cu and the neutral carbon atom is produced in the **A** thin film. During the reaction from the **A** thin film to the resultant **AP40** thin film by post-annealing, when the oxygen content is below 2 ppm in the Ar gas, it cannot react with the film, because the quartz glass plate placed on the **A** thin film can prevent the film from encountering the oxygen molecule. In fact, the copper thin film, which was separately prepared by a vacuum plating method, was not oxidized by post-annealing under an identical condition. Thus, it is accepted that the reduction reaction occurred because of the materials inside the **A** thin film. Under these conditions, only one candidate that can act as a reductant for Cu2+ ion remains on the carbon atoms in film **A**.

The polycrystalline Cu lattices were gradually structured by reducing the valency of the Cu2+ ion with carbon atoms, and the Cu grains were simultaneously grown by annealing. This reaction mechanism involving the reduction reaction caused by carbon atoms may be comparable to the modern and indirect steel-making system using corks. The tensile strength of the **AP40** adhered onto the Na-free glass substrate was 36(12) MPa as determined from the stud pull adherence tests, indicating strong adhesion to the glass substrate. The tensile strength of the Cu film deposited onto an identical Na-free glass substrate by a vacuum plating method was 1.7(5) MPa after an identical heat treatment of the **AP40** thin film. Thus, the tensile strength of the **AP40** thin film on the Na-free substrate was more than 20 times higher than that of the Cu thin film deposited by the vacuum plating method. The covalent bonds between the trace amounts of Cu2+ ion present locally at the interface between the thin film **AP40** and the O2− ions belonging to the Na-free glass molecules may assist in the formation of a robust interface between the Na-free glass substrate and the **AP40** thin film. In fact, the tensile strength of the adhesion of the Cu<sup>2</sup> O thin film to the substrate fabricated using the MPM was 83(2) MPa.

**Figure 10** presents the transmittance and reflectance spectra of the thin films. The transmittance spectra of **A, AP20**, and **AP40** are not significantly different in the UV-Vis region, and the transparency of **AP40** is more than ~30% in the visible region. The infrared reflectance of **AP40** is higher than 40% and reached 100% in the far-infrared region, whereas the reflectance of **A** was low, 20–30%, over this region. The MPM can facilely control the film thickness by adjusting the concentration of Cu ion in the precursor solution under identical spin-coating conditions. When the Cu thin film is 100-nm-thick, the conductivity is 1.8 × 10–5 Ω cm, and the transparency in the visible region is below 5%. Thus, a thicker film indicates higher conductivity, but reduced transparency.

the curve fitting results of the Auger spectrum for **A** over the range 230–300 eV, corresponding to the carbon atoms, along with the curves of the neutral and tetravalent carbon atoms. The curve fitting results suggest that the carbon atom in **A** was 85% agreement to the neutral. The electrical resistivity of the **A, AP20**, and **AP40** thin films is 7.5 × 10–4, 2.8 × 10–4, and 4.7 × 10–5 Ω cm, respectively. Thus, highly conductive translucent copper thin films could be obtained in commercially available Ar gas. A plausible scheme for copper lattice formation, which can be

**Figure 9.** The curve fitting results of carbon state in AP40 thin film. The lines indicate the following: ……A, —Carbon

deduced from the XRD and Auger spectra, is presented in **Scheme 1**.

**Scheme 1.** Plausible scheme for the formation of a Cu thin film from the precursor film.

(neutral), -·-·-·Carbon (tetravalent), respectively.

14 Modern Technologies for Creating the Thin-film Systems and Coatings

Recently, we attempted to embed copper in narrow trenches (0.2–1.0 μm wide and 5.0 μm deep) by using the MPM. A new precursor solution was prepared by dispersing the Cu nanopowder (20–40 nm) into the abovementioned Cu precursor solution. Si substrates with the trenches were immersed in this precursor solution under ultrasonic vibration for 1 min and then slowly withdrawn from the solution. The dip coating and heat treatment steps were repeated twice. The cross-sectional FE-SEM images of the treated substrate indicate that the embedded copper fills the trenches without voids.

**Figure 10.** (A) The transmittance spectra of the resulting thin films. The lines indicate the following: ……A, ----AP20, and ━AP40, respectively. (B) The reflectance spectra of the resulting thin films. The lines indicate the following: ……A, ----AP20, and ━AP40, respectively.

### **5. Conclusion**

Thermal reactions of metal complex films useful for ceramic thin film production such as Cu<sup>2</sup> O did not attract much attention for a long time, with the exception of the CVD procedure. However, we indicated that the MPM provides facile and unique routes to obtain *p*-type Cu<sup>2</sup> O and metallic Cu thin film with excellent adhesion to glass substrates, through the thermal reactions of metal complexes.

The importance of the metal complex in the MPM was presented by using the unprecedented thin film fabrication of *p*-type Cu<sup>2</sup> O, along with the recently elucidated reaction mechanism. In general, impurities such as nitrogen and carbon atoms interfere seriously with the functions of semiconductor devices. However, the present molecular precursor involving nitrogen and carbon atoms is necessary for fabricating an excellent *p*-type Cu<sup>2</sup> O thin film using this solutionbased process. Additionally, in the chemical fabrication of copper thin films with high conductivity, an organic ligand that reacts with the central Cu(II) ion is essential, and the atoms derived from the ligand, which can act as a reductant, prevent the produced copper from oxidation.

It is important that most of the originally included atoms in the MPM system are not involved in the resultant thin films if the amounts and treatment are appropriate. Therefore, the role of the ligand of the metal complex resembles that of auxiliary lines to solve geometrical problems in mathematics. In the sol-gel process, the similarity of the gel composition and the final oxides is desirable and is supposed to be an advantage of the method, though the rearrangement of the polymerized amorphous species to the crystalline requires much energy. From this point of view, the concept of the MPM is quite different from that of the conventional sol-gel method and has many potential applications.

### **Author details**

Hiroki Nagai and Mitsunobu Sato\*

\*Address all correspondence to: ft10302@ns.kogakuin.ac.jp

Department of Applied Physics, School of Advanced Engineering, Kogakuin University, Tokyo, Japan

### **References**

**5. Conclusion**

reactions of metal complexes.


16 Modern Technologies for Creating the Thin-film Systems and Coatings

thin film fabrication of *p*-type Cu<sup>2</sup>

Cu<sup>2</sup>

Thermal reactions of metal complex films useful for ceramic thin film production such as

**Figure 10.** (A) The transmittance spectra of the resulting thin films. The lines indicate the following: ……A, ----AP20, and ━AP40, respectively. (B) The reflectance spectra of the resulting thin films. The lines indicate the following: ……A,

and metallic Cu thin film with excellent adhesion to glass substrates, through the thermal

The importance of the metal complex in the MPM was presented by using the unprecedented

general, impurities such as nitrogen and carbon atoms interfere seriously with the functions of semiconductor devices. However, the present molecular precursor involving nitrogen and

based process. Additionally, in the chemical fabrication of copper thin films with high conductivity, an organic ligand that reacts with the central Cu(II) ion is essential, and the atoms derived from the ligand, which can act as a reductant, prevent the produced copper from oxidation.

It is important that most of the originally included atoms in the MPM system are not involved in the resultant thin films if the amounts and treatment are appropriate. Therefore, the role of the ligand of the metal complex resembles that of auxiliary lines to solve geometrical problems in mathematics. In the sol-gel process, the similarity of the gel composition and the final oxides is desirable and is supposed to be an advantage of the method, though the rearrangement of the polymerized amorphous species to the crystalline requires much energy. From this point of view, the concept of the MPM is quite different from that of the conventional

carbon atoms is necessary for fabricating an excellent *p*-type Cu<sup>2</sup>

sol-gel method and has many potential applications.

O, along with the recently elucidated reaction mechanism. In

O thin film using this solution-

O did not attract much attention for a long time, with the exception of the CVD procedure. However, we indicated that the MPM provides facile and unique routes to obtain *p*-type Cu<sup>2</sup>

O


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[14] Sato M., Hara H., Kuritani M., Nishide T. Novel route to Co<sup>3</sup>

[15] Sato M., Tanji T., Hara H., Nishide T., Sakashita Y. SrTiO<sup>3</sup>

plexes. Sol. Energy Mater. Sol. Cells. 2008;92:1136–1144.

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2008;43:6902–6911.

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strates via N-alkyl substituted amine salt of Co(III)-EDTA complex. Sol. Energy Mater.

synthesis from a non-polymerized precursor system of a stable Ti(IV) complex and Sr(II)

[16] Nagai H., Mochizuki C., Hara H., Takano I., Sato M. Enhanced UV-sensitivity of Visresponsive anatase thin films fabricated by using precursor solutions involving Ti com-

[17] Nagai H., Hasegawa M., Hara H., Mochizuki C., Takano I., Sato M. An important factor controlling the photoreactivity of titania: O-deficiency of anatase thin film. J. Mater. Sci.

[18] Nagai H., Aoyama S., Hara H., Mochizuki C., Takano I., Baba N., Sato M. Rutile thin film responsive to visible light and with high UV light sensitivity. J. Mater. Sci.

[19] Nagai H., Aoyama S., Hara H., Mochizuki C., Takano I., Honda T., Sato M. Photoluminescence and photoreactivity affected by oxygen defects in crystal-oriented rutile thin film

[20] Likius DS., Nagai H., Aoyama S., Mochizuki C., Hara H., Baba N., Sato M. Percolation threshold for electrical resistivity of Ag-nanoparticle/titania composite thin films fabri-

[21] Mochizuki C., Sasaki Y., Hara H., Sato M., Hayakawa T., Yang F., Hu X., Shen H., Wang S. Crystallinity control of apatite through Ca-EDTA complexes and porous composites

[22] Honda T., Oda T., Mashiyama Y., Hara H., Sato M. Fabrication of c-axis oriented Ga-doped MgZnO-based UV transparent electrodes by molecular precursor method.

[23] Werner HA. A forerunner to modern inorganic chemistry. Angew. Chem. Int. Ed.

[24] Nagai H., Sato M. Heat Treatment in Molecular Precursor Method for Fabricating Metal Oxide Thin Films. In: Czerwinski F., editor. Heat Treatment – Conventional and Novel

[25] Nagai H., Sato M. Highly Functionalized Lithium-Ion Battery. In: Yang D., editor. Alkali-

[26] Jayatissa HA., Guo K., Jayasuriya CA. Fabrication of cuprous and cupric oxide thin films

fabricated by molecular precursor method. J. Mater. Sci. 2010;45:5704–5710.

cated using molecular precursor method. J. Mater. Sci. 2012;47:3890–3899.

with PLGA. J. Biomed. Mater. Res. Part B. 2009;90:290–301.

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by heat treatment. Appl. Surf. Sci. 2009;255:9474–9479.

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O4

thin films on glass sub-

film fabrication and powder


**Provisional chapter**

### **Modification of Oxide Thin Films with Low-Energy Ion Bombardment Ion Bombardment**

**Modification of Oxide Thin Films with Low-Energy** 

Oscar Rodríguez de la Fuente Oscar Rodríguez de la Fuente Additional information is available at the end of the chapter

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/66651

### **Abstract**

We review in this chapter the use of low-energy ion bombardment (LEIB) in oxide thin films. In most cases, radiation effects in oxides are ultimately based on the preferential sputtering of the oxygen anions, yielding a chemically reduced oxide. The physics governing the processes in the low ion energy range will be briefly commented here. Also, general uses and applications of LEIB are reviewed here, focusing later in those specific applications on oxide layers. LEIB in oxides has supported, for instance, the fabrication of conductive transparent layers on top of semiconductors or the formation of self-organized morphological surface patterns. Finally, we show a novel application of LEIB when applied on single-crystalline surfaces of some oxides, which is the formation of an epitaxial thin film of the corresponding suboxide. For instance, we show how ion bombardment transform the surface of TiO<sup>2</sup> (110) into an epitaxial TiO(001) thin film.

**Keywords:** low-energy ion bombardment, oxide, thin film, defect, epitaxy

### **1. Introduction**

Ion irradiation of solids has been a research topic for decades. In the keV range, the maximum penetration depth in the material does not exceed a few nanometers, and ions in this range of energies have traditionally assisted in the preparation and analysis of surfaces: it has served as a cleaning tool in vacuum conditions and, combined with surface analysis techniques, in compositional depth analysis or to assist in the growth of thin films. But low-energy ion bombardment (LEIB) has been also a tool to controllably modify surfaces or thin films. It has several advantages over other surface modification methods: (a) its low penetration depth, which allows the modification of a shallow surface layer or a thin film, (b) the high degree of control by choosing the type of ion, energy, flux, and dose received by the sample, and (c) the possibility of modifying very small surface areas by using masks or focused beams.

© 2016 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2017 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

We focus in this chapter in this low-energy range (from a few hundred eV to a few keV), showing the capabilities of the technique for a specific family of systems: thin oxide films. Oxides are materials which display a huge plethora of physical and chemical properties, and in the form of thin films, they have been extensively studied in the last two decades. Its ample tolerance to defects favors the availability of compounds with different stoichiometries and properties. This is a source of complexity but also provides opportunities to explore many different properties in various fields, such as catalysis, magnetism, sensors, or electronics. LEIB stands as a very attractive experimental technique to modify the properties of oxide thin films by creating defects or inducing structural and compositional transitions.

Oxides are present in a large number of technological devices, and a growing number of applications will probably be found in the coming decades. For this to occur, a broad insight of their properties and a deep understanding of the role played by defects will be required. As in the case of silicon technology, where defect engineering is responsible for the precise and powerful control of electronic properties, defect engineering in oxides is for sure the tool to be exploited to enhance the performance of any present or potential device based in oxides. This is by no means a novel statement, for it has been well known for many decades that the presence of defects in oxides determines many of their properties. However, two challenges nowadays are (a) to acquire a more profound knowledge of defects properties in oxides and (b) to be able to introduce them in the surface of the material in a controlled way. LEIB is a good tool for this purpose but goes beyond the simple generation of defects. It also induces the formation of new structural phases or the generation of textured and nanostructured surfaces, as we will describe in this chapter.

#### **1.1. Historical remarks**

Ion sputtering, or the ejection of atoms from surfaces with energetic ions, is a phenomenon reported more than 150 years ago. The effect of sputtering was already described in 1853 by Grove [1], when he observed that the cathode metal located inside discharge tubes gradually coated the inner walls. The coating was caused by the positive ions of the discharges hitting the negatively charged metal cathode, sputtering atoms away from it subsequently coating the glass. However, this was not unambiguously verified until 1902, about 50 years later, when Goldstein [2] performed the correct experiments. In the meantime, other hypotheses were considered (such as that thermal evaporation induced by the discharge). But the effect of atomic collisions from positively charged ions was finally acknowledged. It was in the mid-1900s that a quantitative description of the sputtering process was reached and a sound theory was developed. Since then, hundreds of works have continued improving the existing theories and performing new experiments about surface sputtering. Some of these works are reviewed in the excellent work by Sigmund [3]. The purpose of many of these investigations has been the application of low-energy ion bombardment on a growing number of materials, where of course we can meet oxides.

#### **1.2. Organization of the chapter**

In this work, we start introducing the most important theoretical concepts related to ion bombardment, focusing around the low-energy range. As we are dealing with oxides, the concept of preferential sputtering of multicomponent materials is discussed. Then, we present some of the many applications or current uses of the technique for all types of materials, such as the modification of the surface topography or the controlled introduction of surface defects. Next, we invoke those uses reported in the literature about oxide surfaces and thin films, grouped in two big sets: (a) those examples exclusively describing morphological, structural, and chemical changes of the surfaces and (b) those works reporting other associated physicochemical properties, such as electrical conductivity. Finally, we present two applications recently carried out in our research group, which is the formation of singlecrystalline epitaxial thin films of a suboxide when the surface of the corresponding oxide is ion bombarded: we report the formation of a TiO(001)/TiO<sup>2</sup> (110) thin film and a Fe3 O4 (111)/α-Fe<sup>2</sup> O3 (0001) bilayer.

### **2. Theoretical background**

We present in this section a few theoretical concepts about some of the most relevant physical processes taking place during the collision of ions with solid surfaces. For a more profound study of the physics involved, the reader is referred to other works in the literature [3, 4].

#### **2.1. Ion stopping**

We focus in this chapter in this low-energy range (from a few hundred eV to a few keV), showing the capabilities of the technique for a specific family of systems: thin oxide films. Oxides are materials which display a huge plethora of physical and chemical properties, and in the form of thin films, they have been extensively studied in the last two decades. Its ample tolerance to defects favors the availability of compounds with different stoichiometries and properties. This is a source of complexity but also provides opportunities to explore many different properties in various fields, such as catalysis, magnetism, sensors, or electronics. LEIB stands as a very attractive experimental technique to modify the properties of oxide thin films

Oxides are present in a large number of technological devices, and a growing number of applications will probably be found in the coming decades. For this to occur, a broad insight of their properties and a deep understanding of the role played by defects will be required. As in the case of silicon technology, where defect engineering is responsible for the precise and powerful control of electronic properties, defect engineering in oxides is for sure the tool to be exploited to enhance the performance of any present or potential device based in oxides. This is by no means a novel statement, for it has been well known for many decades that the presence of defects in oxides determines many of their properties. However, two challenges nowadays are (a) to acquire a more profound knowledge of defects properties in oxides and (b) to be able to introduce them in the surface of the material in a controlled way. LEIB is a good tool for this purpose but goes beyond the simple generation of defects. It also induces the formation of new structural phases or the generation of textured and nanostructured surfaces, as we will describe in this chapter.

Ion sputtering, or the ejection of atoms from surfaces with energetic ions, is a phenomenon reported more than 150 years ago. The effect of sputtering was already described in 1853 by Grove [1], when he observed that the cathode metal located inside discharge tubes gradually coated the inner walls. The coating was caused by the positive ions of the discharges hitting the negatively charged metal cathode, sputtering atoms away from it subsequently coating the glass. However, this was not unambiguously verified until 1902, about 50 years later, when Goldstein [2] performed the correct experiments. In the meantime, other hypotheses were considered (such as that thermal evaporation induced by the discharge). But the effect of atomic collisions from positively charged ions was finally acknowledged. It was in the mid-1900s that a quantitative description of the sputtering process was reached and a sound theory was developed. Since then, hundreds of works have continued improving the existing theories and performing new experiments about surface sputtering. Some of these works are reviewed in the excellent work by Sigmund [3]. The purpose of many of these investigations has been the application of low-energy ion bombardment on a growing number of materials,

In this work, we start introducing the most important theoretical concepts related to ion bombardment, focusing around the low-energy range. As we are dealing with oxides, the

by creating defects or inducing structural and compositional transitions.

22 Modern Technologies for Creating the Thin-film Systems and Coatings

**1.1. Historical remarks**

where of course we can meet oxides.

**1.2. Organization of the chapter**

Several mechanisms are active when an ion enters a solid material, dissipating its initial energy. Ion stopping in a solid can take place by interactions with both the electrons and ions of the crystal. The physical quantity describing the interaction of the ion with the solid is the stopping power or cross section *S*(*E*), which depends on the ion energy, and so it can be divided into two terms, an electronic (*e*) and a nuclear (*n*) term:

$$S(E) = S\_\epsilon(E) + S\_\mu(E) = -\frac{1}{N} \left( \left( \frac{dE}{dx} \right)\_\epsilon + \left( \frac{dE}{dx} \right)\_n \right) \tag{1}$$

As Eq. (1) states, *S*(*E*) can be also described in terms of the energy loss rate *dE/dx*, being *N* the number density of atoms in the crystal. It is interesting to note that, as a universal reference, energy losses are typically of several hundreds of eV/nm in the low-energy range.

Collisions of the incident ion with electrons are more frequent, but in the range of low-tomedium energies, they have a little effect on the main trajectory of the ion (because of their small mass). In these ranges, electronic stopping has a comparable effect to that of a viscous flow continuously reducing the kinetic energy of the ion. In contrast, the interaction with the ions is more discrete, taking place between elastic consecutive collisions. For these nuclear losses, a simple elastic model using conservation of momentum describes the maximum transferred energy *Tmax* from an incident atom to the struck atom as

$$T\_{\text{max}} = \frac{4 \, E\_{\text{o}} \, M\_{\text{i}} \, M\_{\text{z}}}{(M\_{\text{i}} + M\_{\text{z}})^2} \tag{2}$$

where *M1* and *M2* are the masses of both atoms and *E0* is the kinetic energy of the incident one. Depending on the nature of the collision, the effect produced has been typically classified in three regimes. In the *single knock-on* regime, the incident ion generates a small cascade of events, where atoms sequentially transfer their energy. Eventually, one of the atoms surpasses the lattice binding forces and is ejected out of the solid. But if the ion is sufficiently energetic, secondary cascades are generated, each of which can produce their own sputtered atoms. But the whole process can still be described as a sum of individual events. This is called the *linear cascade* regime. If, instead, the energy density of the collision is very high and most of the atoms located in a region are severely displaced from their equilibrium positions, looking more like a melted fluid, the regime is called *thermal spike*. The range where a given event is placed depends not only on the ion energy but also on the ion mass. As can be deduced from Eq. (2), a light ion (H<sup>+</sup> , He<sup>+</sup> , etc.) is very inefficient in its energy transfer *Tmax* and enters the spike regime at very high ion energies. Most of the cases considered in this chapter, with medium-mass ions (mostly Ar<sup>+</sup> ) with up to a few keV of energy, lie in the *single knock-on* or *linear cascade* regime.

#### *2.1.1. Ion range*

To evaluate the spatial range of the area of the ion-bombarded material, the most important concept is that of the ion range *R*. The ion range in a solid is the integrated distance traveled by the ion before it stops after reducing its energy to zero, and so it is defined as

$$\mathcal{R} = \int\_{\mathcal{E}, \, d\mathcal{E} \, \mathrm{d}\mathbf{x}}^{0} \mathrm{d}E \tag{3}$$

Another useful magnitude is the projected range *Rp* , defined as the distance of *R* projected along the incoming direction of the ion before entering the solid. In the case of normal incidence, *Rp* is just the depth at which the ion remains implanted, which is not necessarily the distance traveled *R*. **Figure 1** shows schematically these parameters.

#### **2.2. Partial sputtering yield**

The sputtering yield is the average number of atoms leaving the solid per incoming ion. An atom is sputtered off the solid when it receives an impact which overcomes the threshold energy and has the possibility to exit the solid and enter the vacuum region. While the former condition is achieved by many atoms involved in cascade collisions, the last condition is only fulfilled by a few atoms close to the surface with their linear momentum pointing away toward vacuum.

As we are dealing with oxides, the concept of preferential sputtering is the most important one. Ion irradiation of multicomponent surfaces always involves compositional changes. Many works in the literature deal with these types of processes, which we will briefly expose here (the reader is referred to Ref. [5] for an excellent review of the topic). The most important physical quantity in a multicomponent material is the partial sputtering yield *Yi* of element *i*, which is the average number of ejected *i* atoms per incoming ion. Mathematically, it can be defined as

$$Y\_{\parallel} = \int\_{0}^{\omega} p\_{\parallel}(\mathbf{z}) \, \mathbf{N}\_{\parallel}(\mathbf{z}) \, d\mathbf{z} \tag{4}$$

 **Figure 1.** Scheme of the main parameters from Eqs. (3) to (6) for an ion impinging a surface at an angle *θ*. The center of the ellipsoid represents the region where the ion deposits most of its energy.

where the partial sputtering yield depends on the density *Ni* (*z*) of atoms of type *i* (in atoms per unit volume) at a given depth *z* below the surface and with a sputtering probability *pi* (*z*), which decays to zero with increasing depth.

#### *2.2.1. Preferential sputtering of oxygen*

Depending on the nature of the collision, the effect produced has been typically classified in three regimes. In the *single knock-on* regime, the incident ion generates a small cascade of events, where atoms sequentially transfer their energy. Eventually, one of the atoms surpasses the lattice binding forces and is ejected out of the solid. But if the ion is sufficiently energetic, secondary cascades are generated, each of which can produce their own sputtered atoms. But the whole process can still be described as a sum of individual events. This is called the *linear cascade* regime. If, instead, the energy density of the collision is very high and most of the atoms located in a region are severely displaced from their equilibrium positions, looking more like a melted fluid, the regime is called *thermal spike*. The range where a given event is placed depends not only on the ion energy but also on the ion mass. As can be deduced from Eq. (2), a

at very high ion energies. Most of the cases considered in this chapter, with medium-mass ions

To evaluate the spatial range of the area of the ion-bombarded material, the most important concept is that of the ion range *R*. The ion range in a solid is the integrated distance traveled

> 0 \_\_\_\_\_ *dE*

along the incoming direction of the ion before entering the solid. In the case of normal inci-

The sputtering yield is the average number of atoms leaving the solid per incoming ion. An atom is sputtered off the solid when it receives an impact which overcomes the threshold energy and has the possibility to exit the solid and enter the vacuum region. While the former condition is achieved by many atoms involved in cascade collisions, the last condition is only fulfilled by a few atoms close to the surface with their linear momentum pointing away toward

As we are dealing with oxides, the concept of preferential sputtering is the most important one. Ion irradiation of multicomponent surfaces always involves compositional changes. Many works in the literature deal with these types of processes, which we will briefly expose here (the reader is referred to Ref. [5] for an excellent review of the topic). The most important physi-

is the average number of ejected *i* atoms per incoming ion. Mathematically, it can be defined as

 0 -∞ *pi* (*z*) *Ni*

is just the depth at which the ion remains implanted, which is not necessarily the

by the ion before it stops after reducing its energy to zero, and so it is defined as

, etc.) is very inefficient in its energy transfer *Tmax* and enters the spike regime

*dE* /*dx* (3)

, defined as the distance of *R* projected

of element *i*, which

(*z*) *dz* (4)

) with up to a few keV of energy, lie in the *single knock-on* or *linear cascade* regime.

light ion (H<sup>+</sup>

(mostly Ar<sup>+</sup>

dence, *Rp*

vacuum.

*2.1.1. Ion range*

, He<sup>+</sup>

**2.2. Partial sputtering yield**

*R* = ∫*Ei*

24 Modern Technologies for Creating the Thin-film Systems and Coatings

*Yi* = ∫

Another useful magnitude is the projected range *Rp*

distance traveled *R*. **Figure 1** shows schematically these parameters.

cal quantity in a multicomponent material is the partial sputtering yield *Yi*

In a simple binary oxide, the ratio between the partial sputtering yields can be described as Eq. (4)

$$\frac{Y\_M}{Y\_O} = \frac{N\_M}{N\_O} \left(\frac{M\_O}{M\_M}\right)^{2m} \left(\frac{lI\_O}{lI\_M}\right)^{1\cdot 2m} \tag{5}$$

where *Ni* are the atomic densities, *Mi* are the masses, and *Ui* are the binding energies of the species (*M* stands for the metal cation and *O* for the oxygen anion). *m* is an exponent which characterizes the type of interatomic potential describing the collision and varies from 1 to 0 in the range from high to very low energies, respectively. Eq. (5) predicts that oxygen atom (the lightest atom and also frequently the most weakly bound to the crystal lattice) will be more easily ejected off the solid. However, Eq. (5) only describes a transient state. If the ion penetration depth is much smaller than the thickness of the sample bombarded, a steady state (after a sufficiently large ion fluence) will be reached where the bombarded material loses elements in a ratio equal to its bulk composition. That is a simple consequence of mass conservation, and it is achieved if mass diffusion to long distances is not relevant.

According to the effects previously addressed, a clear consequence of ion sputtering is the formation of a modified layer, not only in terms of its structure or its density of defects but also in terms of chemical composition. A fact to take into account is that the depth of the altered layer may be larger than the penetration depth of the incoming ions. The reason can rely on thermal or bombardment-induced diffusion. Dissipated energy, either from the ion beam or from mechanical agitation from collisions, enhances diffusion. The diffusion is driven by a chemical gradient, which can be present even from many atomic layers below *Rp* (so, relatively far from the direct influence of the impinging ion).

### **3. Applications of low-energy ion bombardment**

### **3.1. Surface cleaning and sample preparation**

Probably, the most extended application of ion bombardment is the cleaning of the surfaces prior to their analysis or before another physical or chemical process. For instance, those surfaces of hard drilling tools which are to be coated with nitrides by means of physical vapor deposition (PVD) are usually exposed to Ar<sup>+</sup> prior to the coating. The reasons are the cleaning of the surface and the enhancement of atomic rugosity, which will favor mechanical adhesion of the coating, and thus will warrant a longer life of the tool. This process is one of the many examples in surface science and engineering which requires LEIB. Another extended use occurs for transmission electron microscopy (TEM), which requires thin (electron transparent) samples. The final stage is usually carried out with ion bombardment, usually named ion milling.

#### **3.2. Secondary ion mass spectroscopy (SIMS)**

Ion bombardment, sometimes in conjunction with surface analysis techniques, has been traditionally used to obtain surface composition. Secondary ion mass spectroscopy (SIMS) analyses, with the help of a mass spectrometer, the nature of the atoms sputtered from a surface with an ion beam [6]. It is a useful technique to carry out depth profiling studies, especially if combined with XPS or Auger spectroscopy. However, the change in the chemical state of the surface, induced by the ion beam, must always be taken into account (except if the surface layer or layers to be analyzed are single component).

### **3.3. Improvement of thin-film growth**

The simultaneous combination of thin-film growth (by means of any vacuum technique) and LEIB has been frequently employed to explore different states of the film grown and to improve the required properties. The effect of the incoming energetic ions has a profound influence in the morphology of the film, as well as in its defect density, which in turn modifies many other properties. The technique, usually called ion-beam-assisted deposition (IBAD), can produce a number of beneficial changes in several characteristics of the film, such as density, texture, residual stresses, adhesion, or crystalline order. Of course, this technique has been applied to thin films of oxides [7]. In most cases, the main role of the ions of the beam is to deposit energy, creating defects and inducing diffusion and mobility. But in some other cases (this is usually named reactive IBAD), the second role of the atoms in the beam is to chemically react with surface species to form compounds.

### **3.4. Controlled generation of surface defects**

in a ratio equal to its bulk composition. That is a simple consequence of mass conservation,

According to the effects previously addressed, a clear consequence of ion sputtering is the formation of a modified layer, not only in terms of its structure or its density of defects but also in terms of chemical composition. A fact to take into account is that the depth of the altered layer may be larger than the penetration depth of the incoming ions. The reason can rely on thermal or bombardment-induced diffusion. Dissipated energy, either from the ion beam or from mechanical agitation from collisions, enhances diffusion. The diffusion is driven by a chemi-

Probably, the most extended application of ion bombardment is the cleaning of the surfaces prior to their analysis or before another physical or chemical process. For instance, those surfaces of hard drilling tools which are to be coated with nitrides by means of physical vapor

of the surface and the enhancement of atomic rugosity, which will favor mechanical adhesion of the coating, and thus will warrant a longer life of the tool. This process is one of the many examples in surface science and engineering which requires LEIB. Another extended use occurs for transmission electron microscopy (TEM), which requires thin (electron transparent) samples. The final stage is usually carried out with ion bombardment, usually named

Ion bombardment, sometimes in conjunction with surface analysis techniques, has been traditionally used to obtain surface composition. Secondary ion mass spectroscopy (SIMS) analyses, with the help of a mass spectrometer, the nature of the atoms sputtered from a surface with an ion beam [6]. It is a useful technique to carry out depth profiling studies, especially if combined with XPS or Auger spectroscopy. However, the change in the chemical state of the surface, induced by the ion beam, must always be taken into account (except if the surface

The simultaneous combination of thin-film growth (by means of any vacuum technique) and LEIB has been frequently employed to explore different states of the film grown and to improve the required properties. The effect of the incoming energetic ions has a profound influence in the morphology of the film, as well as in its defect density, which in turn modifies

(so, relatively far

prior to the coating. The reasons are the cleaning

and it is achieved if mass diffusion to long distances is not relevant.

cal gradient, which can be present even from many atomic layers below *Rp*

from the direct influence of the impinging ion).

26 Modern Technologies for Creating the Thin-film Systems and Coatings

**3.1. Surface cleaning and sample preparation**

deposition (PVD) are usually exposed to Ar<sup>+</sup>

**3.2. Secondary ion mass spectroscopy (SIMS)**

layer or layers to be analyzed are single component).

**3.3. Improvement of thin-film growth**

ion milling.

**3. Applications of low-energy ion bombardment**

If bulk defects control many properties of solids (dislocations in the case of mechanical properties of metals or point defects in the case of optical properties of many oxides, to mention two time-honored examples), surface defects may as well control some surface physicochemical properties. The most paradigmatic example is that of the so-called active centers [8], which control some catalytic conversions and which have been proposed to be surface steps [9] or vacancies. A profound knowledge of the physical and chemical properties of surface defects requires a well-known process to generate them controllably, and LEIB is a good tool for that purpose [10–12]. Under the correct parameters (energy, dose, temperature, etc.), it can generate a broad spectrum of surface defects (surface vacancies, adatoms, linear steps, vacancy islands, etc.), which can be subsequently studied for whatever phenomenon of interest [13, 14]. Also, the kinetics of flattening of an initially rough surface, with the roughness induced with LEIB, has been studied for different types of materials (among them, also oxides). For instance, for the case of TiO<sup>2</sup> (110) and after low ion doses (only 0.5 monolayers were sputtered), a power scaling law has been found for the high-temperature evolution of the width *w* of the surface terraces as *w* ~ *t* 0.24 [15].

### **3.5. Surface nanopatterning**

One of the first phenomena observed on ion-bombarded surfaces was the formation, in some cases, of ordered structures with a well-defined periodicity of a few tens of nanometers [16]. These patterns, usually in the form of ripples, may develop spontaneously from an initially flat surface as a result of several competing mechanisms which are active in a wide variety of materials. This phenomenon soon caught the attention of the scientific community for its interest and potential applications. The method just requires an ion beam homogeneously irradiating a surface and can easily pattern macroscopic areas of the order of mm<sup>2</sup> or cm2 (once the correct processing conditions have been achieved). Thus, the absence of masks, complicated nanolithographic methods, or focused ion beams is very attractive from the technological point of view.

As already mentioned, the self-organized nanostructures spontaneously emerge under those experimental conditions which promote a clearly balanced competition between opposed kinetic processes occurring at the surface. Ion bombardment induces surface disordering and roughening, while surface diffusion favors smoothing and defect annihilation. These two main opposed trends, acting under nonequilibrium conditions and summed with other mechanisms or condition characteristic of each type of surface, determine the evolution and the final morphology of the surface.

The family of theories describing surface patterning is too ample to be mentioned here, but we can briefly comment here on the most acknowledged model, which accounts for the formation of ripples under an off-normal incident ion beam. This is the Bradley-Harper (BH) model [17], which successfully explains the dynamics of ripple formation in isotropic and amorphous surfaces: the ripples are parallel to the incident direction of the beam under high incidence angles and perpendicular to the beam direction under angles close to be grazing. To understand the BH model, one must consider the distribution of surface defects created by ion collisions. This is described by the Sigmund model [18], which accounts, under a continuum model approach, for an ellipsoidal average distribution of energy deposited under ion impacts formulated as

$$\mathbb{E}(\mathbf{z}',\boldsymbol{\rho}) = \frac{\varepsilon}{(2\pi)^{13}a\beta^2} \exp\left(-\frac{(\mathbf{z}'-\mathbf{a})^2}{2} \cdot \frac{\boldsymbol{\rho}^2}{2\beta^2}\right) \tag{6}$$

Under cylindrical coordinates, *z*′ and *ρ* are the radial and longitudinal components, oriented along the initial incidence ion direction. *α* and *β* are the longitudinal and lateral spreadings of the spatial distribution, *a* is the center of the ellipsoidal distribution, and *ε* represents the total energy deposited. **Figure 1** graphically displays these parameters. That part of the energy deposited close to the surface will create a surface vacancy by sputtering an atom off the solid. Bradley and Harper used the Sigmund distribution to account for the creation of surface defects, describing their effect on surface topography *h*(*x,y,t*) with a partial differential rate equation, where *x* represents the direction parallel to the projection of the ion beam on the surface:

$$\frac{\partial h(\mathbf{x}, y, t)}{\partial t} = \mathbf{-Y}\_0(\Theta) + \frac{\partial Y\_0}{\partial \theta} \frac{\partial h}{\partial \mathbf{x}} + \mathbf{v}\_0 \frac{\partial^2 h}{\partial \mathbf{x}^2} + \mathbf{v}\_\perp \frac{\partial^2 h}{\partial y^2} \mathbf{-K} \, (\nabla^2)^2 h \tag{7}$$

The first term in Eq. (7) is the erosion rate *Y0* for a flat surface, whose explicit dependency with the incidence angle *θ* is considered in the second term. The third and fourth terms take into account the contribution to roughening with the most important mechanism: the curvature-dependent sputtering yield. The concave regions of the surface have a larger probability to be even more eroded by the ion beam. Surface atoms at the valleys are more easily sputtered than those at the crests, so that once a valley is initiated, its height difference with the crests continues growing. This is, very naively described, the surface instability necessary to start the formation of ripples. The last term considers surface smoothing by atomic self-diffusion.

The BH model is relatively simple, and more advanced equations, including higher order terms or surface anisotropy, have been developed in the last two decades. But the BH model, as described by Eq. (7), agrees well with many experimental results, including the rotation of the ripples depending on the incidence angle of the ion beam. The BH theory only considers amorphous and isotropic solids and ignores the existence of crystallographic directions and surface steps. Surprisingly, the BH theory has proven successful in many crystalline surfaces and has been a good starting point in other cases. However, surface anisotropy induces new effects and mechanisms to be taken into account. For instance, surface steps may limit diffusion via Ehrlich-Schwöbel barriers, or may determine etching rates at low grazing angles, as we will see later. In any case, ion beam patterning has been observed in different types of materials, including oxides. We will refer to those works on surface nanopatterning of oxides in a specific section later.

### **4. LEIB of oxide surfaces and thin films**

mechanisms or condition characteristic of each type of surface, determine the evolution and

The family of theories describing surface patterning is too ample to be mentioned here, but we can briefly comment here on the most acknowledged model, which accounts for the formation of ripples under an off-normal incident ion beam. This is the Bradley-Harper (BH) model [17], which successfully explains the dynamics of ripple formation in isotropic and amorphous surfaces: the ripples are parallel to the incident direction of the beam under high incidence angles and perpendicular to the beam direction under angles close to be grazing. To understand the BH model, one must consider the distribution of surface defects created by ion collisions. This is described by the Sigmund model [18], which accounts, under a continuum model approach, for an ellipsoidal average distribution of energy deposited under

, *<sup>ρ</sup>*) <sup>=</sup> \_\_\_\_\_\_\_\_ *<sup>ε</sup>*

(2*π*)1.5 *<sup>α</sup> <sup>β</sup>*<sup>2</sup> exp(-

Under cylindrical coordinates, *z*′ and *ρ* are the radial and longitudinal components, oriented along the initial incidence ion direction. *α* and *β* are the longitudinal and lateral spreadings of the spatial distribution, *a* is the center of the ellipsoidal distribution, and *ε* represents the total energy deposited. **Figure 1** graphically displays these parameters. That part of the energy deposited close to the surface will create a surface vacancy by sputtering an atom off the solid. Bradley and Harper used the Sigmund distribution to account for the creation of surface defects, describing their effect on surface topography *h*(*x,y,t*) with a partial differential rate equation, where *x* represents the direction parallel to the projection of the ion beam on the surface:

(*z*′ - *a*)2 \_\_\_\_\_ <sup>2</sup> *<sup>α</sup>*<sup>2</sup> - *<sup>ρ</sup>*<sup>2</sup> \_\_\_

∂<sup>2</sup> \_\_\_*h* <sup>∂</sup>*x*<sup>2</sup> <sup>+</sup> *<sup>v</sup>*<sup>⊥</sup>

with the incidence angle *θ* is considered in the second term. The third and fourth terms take into account the contribution to roughening with the most important mechanism: the curvature-dependent sputtering yield. The concave regions of the surface have a larger probability to be even more eroded by the ion beam. Surface atoms at the valleys are more easily sputtered than those at the crests, so that once a valley is initiated, its height difference with the crests continues growing. This is, very naively described, the surface instability necessary to start the formation of ripples. The last term considers surface smoothing by atomic

The BH model is relatively simple, and more advanced equations, including higher order terms or surface anisotropy, have been developed in the last two decades. But the BH model, as described by Eq. (7), agrees well with many experimental results, including the rotation of the ripples depending on the incidence angle of the ion beam. The BH theory only considers amorphous and isotropic solids and ignores the existence of crystallographic directions and surface steps. Surprisingly, the BH theory has proven successful in many crystalline surfaces

∂<sup>2</sup> \_\_\_*h*

<sup>2</sup> *<sup>β</sup>*<sup>2</sup>) (6)

<sup>∂</sup>*y*<sup>2</sup> - *<sup>K</sup>* (∇2)2 *<sup>h</sup>* (7)

for a flat surface, whose explicit dependency

the final morphology of the surface.

28 Modern Technologies for Creating the Thin-film Systems and Coatings

ion impacts formulated as

*E*(*z*′

<sup>∂</sup>*h*(*x*, *<sup>y</sup>*, *<sup>t</sup>*) \_\_\_\_\_\_\_\_

self-diffusion.

<sup>∂</sup>*<sup>t</sup>* <sup>=</sup> -*Y*<sup>0</sup>

The first term in Eq. (7) is the erosion rate *Y0*

(*θ*) + <sup>∂</sup>*Y*\_\_\_0 ∂*θ* \_\_\_ ∂*h* <sup>∂</sup>*<sup>x</sup>* <sup>+</sup> *<sup>v</sup>*<sup>∥</sup> Oxides are a colossal family of materials, both in the sense of the existing diversities and of the properties displayed. The specific and complex nature of the bond between the oxygen and the metallic cation favors the existence of a vast set of functionalities in oxides. Complicated interactions are present, which cross-link the different properties of the oxide. Also, there is a huge variety of defects in oxides, which are rather abundant, on the other hand. This circumstance enormously obscures the profound understanding of the physical mechanisms involved in oxide properties but also assists in the modification or control of these properties.

Oxide thin films lie at the core of many technological devices [19]. Indium-tin-oxide (ITO) conductive and transparent coatings are a good example of this statement. The modification of the surface morphology and the controlled introduction of defects in thin oxide films are tools to improve their response in their respective applications. LEIB can assist in this task.

### **4.1. Modification of the morphology**

Oxides have been also used to spontaneously generate ordered nanostructures in their surfaces with LEIB. Indeed, the first reported example took place for a glass surface bombarded with Ar<sup>+</sup> at 4 keV by Navez in 1956 [16]. The authors discovered at that moment the formation of ripples separated by tens of nanometers. For incidence angles close to normal incidence and up to *θ* = 80°, the direction of the ripples was perpendicular to the incidence angle. For grazing angles, in contrast, the direction of the ripples was parallel to the ion beam. This was probably the first time that the ripple rotation mechanism, later explained in the BH theory, was observed. Recent works on amorphous SiO<sup>2</sup> have found that, while the quantitatively results are similar [20], there are differences depending on the type of silica (fused silica, amorphized silica, or thermally grown SiO<sup>2</sup> ). The wavelength dependence on the energy and the wavelength coarsening vary with the substrate, which in turn depends on their surface energies (see **Figure 2**). It is interesting to note that, for low incidence angles, the surfaces remain smooth.

Although we focus in this chapter in the use of low-energy ions, medium energies have been used to nanostructure oxide surfaces too, and we would like to briefly mention it. However, in those cases where the energy is of the order of 10–100 keV, new mechanisms arise in the

 **Figure 2.** Ripple wavelength dependence on ion energy (left) and total ion dose (right) for three different SiO<sup>2</sup> surfaces. The total dose for the graph on the left is *Φ* = 1 × 10<sup>18</sup> cm−2. The incidence angle for both graphs is 45°. Reproduced with permission from Ref. [20].

evolution of the surface morphology, which are far from the mechanisms invoked by BH theories and similar. That is the case of anatase TiO<sup>2</sup> , for example, where different types of ions have been used to bombard the samples at different temperatures [21]. At these energies, the ion projected range *Rp* lies deep below the surface (of the order of tens of nanometers), and the agglomeration of point defects evolves into small voids in the first stages of ion bombardment. Further bombardment and void growth transform the initially flat surface into a morphology of void and mounds or nanorods aligned with the ion beam. In contrast, for the case of low-energy ions, the evolution of morphology does not include void formation, but is more limited to the formation and evolution of defects in the upper layers of the material.

Since TiO<sup>2</sup> (110) is the most studied oxide surface, it can be understood that the nanopatterning of this surface has been addressed several times. In the case of low-energy ions, the morphology of the surface is, in many cases, rough and disordered if the bombardment takes place at room temperature. Diffusion is more limited in oxides than in metals, and there is no active mechanism to induce self-ordering at low T. In the case of TiO<sup>2</sup> (110) bombarded with Ar<sup>+</sup> with an energy between 0.5 and 1.5 eV at room temperature, a disordered set of mounds is distributed across the surface [22]. These mounds are proposed to be chemically reduced oxide Ti<sup>2</sup> O3 regions, which align themselves along a given crystallographic direction under a thermal treatment between 700 K and 900 K.

However, at higher temperatures, the formation of ordered ripples in TiO<sup>2</sup> (110) at grazing incidence has been reported in Ref. [23]. The TiO<sup>2</sup> (110) orientation has two nonequivalent in-plane surface directions: (001) and (1–10). At low incidence angles, the existence of surface steps is determinant in ripple formation. Indeed, sputtering from terraces is about one order of magnitude smaller than from step edges. Also, the sputtering yield or erosion rate of the different steps differs, and the consequence is that the total sputtering yield is larger when the beam is directed along the [1–10] direction than when it is oriented along the [001] direction. Despite all, at sufficiently long bombardment times, the beam direction does not affect too much the morphology of the formed ripples, somehow demonstrating the universality of the method for patterning large areas in different materials.

Another phenomenon which also recalls universality is the temperature-dependent reorientation of the ripples observed in TiO<sup>2</sup> (110) [24], in the same way as for many metal surfaces. While at low (150 K) and high (620 and 720 K) temperatures ripples are aligned along the beam direction, at 300 K they are aligned perpendicularly, as illustrated in **Figure 3**. The same group has used LEIB as a tool to prepare TiO<sup>2</sup> (110) surfaces with steps aligned along special orientations [25]. The method can be viewed as the equivalent to surface texturing with ion beams when the grains with higher sputtering yields destabilize. In this case, the ion beam under a low incidence angle preferentially erodes some specific steps, creating [1–10] steps. These are thermodynamically unstable but kinetically stabilized within this approach.

evolution of the surface morphology, which are far from the mechanisms invoked by BH

The total dose for the graph on the left is *Φ* = 1 × 10<sup>18</sup> cm−2. The incidence angle for both graphs is 45°. Reproduced with

 **Figure 2.** Ripple wavelength dependence on ion energy (left) and total ion dose (right) for three different SiO<sup>2</sup>

ions have been used to bombard the samples at different temperatures [21]. At these energies, the ion projected range *Rp* lies deep below the surface (of the order of tens of nanometers), and the agglomeration of point defects evolves into small voids in the first stages of ion bombardment. Further bombardment and void growth transform the initially flat surface into a morphology of void and mounds or nanorods aligned with the ion beam. In contrast, for the case of low-energy ions, the evolution of morphology does not include void formation, but is more limited to the formation and evolution of defects in the upper layers

(110) is the most studied oxide surface, it can be understood that the nanopattern-

regions, which align themselves along a given crystallographic direction under a

ing of this surface has been addressed several times. In the case of low-energy ions, the morphology of the surface is, in many cases, rough and disordered if the bombardment takes place at room temperature. Diffusion is more limited in oxides than in metals, and there is no

 with an energy between 0.5 and 1.5 eV at room temperature, a disordered set of mounds is distributed across the surface [22]. These mounds are proposed to be chemically reduced

in-plane surface directions: (001) and (1–10). At low incidence angles, the existence of surface steps is determinant in ripple formation. Indeed, sputtering from terraces is about one order of magnitude smaller than from step edges. Also, the sputtering yield or erosion rate of the different steps differs, and the consequence is that the total sputtering yield is larger when the beam is directed along the [1–10] direction than when it is oriented along the [001] direction. Despite all, at sufficiently long bombardment times, the beam direction does not affect too

active mechanism to induce self-ordering at low T. In the case of TiO<sup>2</sup>

However, at higher temperatures, the formation of ordered ripples in TiO<sup>2</sup>

, for example, where different types of

surfaces.

(110) bombarded with

(110) at grazing

(110) orientation has two nonequivalent

theories and similar. That is the case of anatase TiO<sup>2</sup>

30 Modern Technologies for Creating the Thin-film Systems and Coatings

of the material.

permission from Ref. [20].

Since TiO<sup>2</sup>

Ar<sup>+</sup>

oxide Ti<sup>2</sup>

O3

thermal treatment between 700 K and 900 K.

incidence has been reported in Ref. [23]. The TiO<sup>2</sup>

 **Figure 3.** (a), (b), (c) and (d): STM images (200 × 200 nm2 ) of Ar<sup>+</sup> -modified TiO<sup>2</sup> (110) surfaces, where the bombardment has been carried out at different temperatures under an incidence angle of 75°, an ion energy of 2 keV, and a total dose of 9 × 1016 ions/cm<sup>2</sup> . The LEED pattern in (a) shows the crystalline order of the surface even after ion bombardment. All ripples are elongated along the beam projection, except for the case where bombardment has been done at RT. The black vertical arrow in (a) marks the projection of the beam on the surface and the horizontal black arrows mark the indicated crystallographic direction. Reproduced with permission from Ref. [24].

This work illustrates how grazing LEIB can be used to tune and texture certain surface steps, which can be relevant for applications sensitive to specific step orientations.

#### **4.2. Modification of the electronic properties**

Oxygen vacancy creation has been traditionally used to modify electronic properties of oxides. It is indeed a route to doping, since excess electrons (due to oxygen vacancies) can be transferred to the conduction band, transforming the material into a better conductor. This type of approaches has been sometimes referred to as vacancy engineering. The control to achieve the desired vacancy concentration can be gained through post-growth thermal annealings or during the growth, controlling experimental parameters such as the oxygen partial pressure or the laser energy if the material is grown by sputtering or by pulsed laser epitaxy (PLE) [26]. As already discussed, LEIB can be also helpful to control oxygen vacancies. Indeed, many of the most recent applications of LEIB on oxide surfaces are related to the modification of electronic and transport properties for different applications. In turn, other fundamental properties, such as the optical response, are modified as well. Most of the studies about oxides modified with LEIB have been accomplished on SrTiO3 and TiO<sup>2</sup> .

It is easy to understand that one of the oxides where LEIB has been most often employed to exploit its properties is SrTiO3 (STO), a transparent insulating perovskite with many intriguing characteristics and also frequently used as a substrate to grow other materials on it. A very interesting phenomenon is that LEIB can generate a conductive layer at its surface, and that has been the subject of several studies. There is a clear correlation between conductivity and the concentration of oxygen vacancies created by ion bombardment in STO [27]. The resistivity values of a modified nanometric layer obtained after bombarding a SrTiO3 (100) surface can follow a ~ *T*2.5 law, dominated by the mobility of the carriers, and attain very low values (between 2 × 10−4 and 6 × 10−4 Ω cm at room temperature) without losing transparency [28]. In order to compare, some of the best conducting STO samples were previously produced by boron implantation at 100 keV, resulting in relatively thick and nontransparent conducting layers with resistivity values of about 0.01 Ω cm [29]. Thus, LEIB on STO is able to produce transparent conductive layers competitive with ITO layers, with an even smaller thickness.

This type of conductive layers formed by LEIB has also been reported in rutile TiO<sup>2</sup> [30], reporting again high carrier densities and mobilities. The authors account for a high crystallographic order in the ion-bombarded region of the TiO<sup>2</sup> (100) surface, which is not compatible with the rutile structure and which is attributed to the ordering of the vacancies produced by ion bombardment.

With respect to optical properties, hole levels in self-trapped states localized in the gap can be stabilized by conduction carriers (generated by oxygen deficiency) in Ar<sup>+</sup> -bombarded single-crystalline STO [31]. Their recombination generates the emission of blue light at room temperature, and the emitting regions can be conveniently patterned. Cross-sectional TEM images show the formation of amorphous and modified (oxygen deficient) layers under the action of the ion beam (**Figure 4**).

This work illustrates how grazing LEIB can be used to tune and texture certain surface steps,

Oxygen vacancy creation has been traditionally used to modify electronic properties of oxides. It is indeed a route to doping, since excess electrons (due to oxygen vacancies) can be transferred to the conduction band, transforming the material into a better conductor. This type of approaches has been sometimes referred to as vacancy engineering. The control to achieve the desired vacancy concentration can be gained through post-growth thermal annealings or during the growth, controlling experimental parameters such as the oxygen partial pressure or the laser energy if the material is grown by sputtering or by pulsed laser epitaxy (PLE) [26]. As already discussed, LEIB can be also helpful to control oxygen vacancies. Indeed, many of the most recent applications of LEIB on oxide surfaces are related to the modification of electronic and transport properties for different applications. In turn, other fundamental properties, such as the optical response, are modified as well. Most of the studies about oxides

It is easy to understand that one of the oxides where LEIB has been most often employed to

ing characteristics and also frequently used as a substrate to grow other materials on it. A very interesting phenomenon is that LEIB can generate a conductive layer at its surface, and that has been the subject of several studies. There is a clear correlation between conductivity and the concentration of oxygen vacancies created by ion bombardment in STO [27]. The resistivity values of a modified nanometric layer obtained after bombarding a SrTiO3

surface can follow a ~ *T*2.5 law, dominated by the mobility of the carriers, and attain very low values (between 2 × 10−4 and 6 × 10−4 Ω cm at room temperature) without losing transparency [28]. In order to compare, some of the best conducting STO samples were previously produced by boron implantation at 100 keV, resulting in relatively thick and nontransparent conducting layers with resistivity values of about 0.01 Ω cm [29]. Thus, LEIB on STO is able to produce transparent conductive layers competitive with ITO layers, with an even smaller

This type of conductive layers formed by LEIB has also been reported in rutile TiO<sup>2</sup>

reporting again high carrier densities and mobilities. The authors account for a high crystal-

with the rutile structure and which is attributed to the ordering of the vacancies produced by

With respect to optical properties, hole levels in self-trapped states localized in the gap can

gle-crystalline STO [31]. Their recombination generates the emission of blue light at room temperature, and the emitting regions can be conveniently patterned. Cross-sectional TEM images show the formation of amorphous and modified (oxygen deficient) layers under the

be stabilized by conduction carriers (generated by oxygen deficiency) in Ar<sup>+</sup>

and TiO<sup>2</sup>

(STO), a transparent insulating perovskite with many intrigu-

.

(100)

[30],


(100) surface, which is not compatible

which can be relevant for applications sensitive to specific step orientations.

**4.2. Modification of the electronic properties**

32 Modern Technologies for Creating the Thin-film Systems and Coatings

modified with LEIB have been accomplished on SrTiO3

lographic order in the ion-bombarded region of the TiO<sup>2</sup>

exploit its properties is SrTiO3

thickness.

ion bombardment.

action of the ion beam (**Figure 4**).

 **Figure 4.** (a) TEM cross-sectional image from an Ar<sup>+</sup> -modified STO single crystal. (b)–(d) electron diffraction images of the (b) red, (c) blue, and (c) yellow regions of the material. Diffraction patterns from (c) to (d) are qualitatively similar, but the contrast in the real space image suggests that a modification indeed exists: the ion beam transforms the upper layer of pristine STO into an amorphous film (red region) and induces the formation of an oxygen-deficient SrTiO3−x region with a thickness around 15–20 nm (blue region). Reproduced with permission from Ref. [31].

Under certain conditions, LEIB can also modify the optical properties of TiO<sup>2</sup> (110) substrates, greatly enhancing its absorbance in the visible range and its luminescence [32]. The formation of self-organized crystalline nanodots at the surface and the generation of Ti interstitials justify this optical response. In general, this type of investigations opens a path to explore the field of oxide optoelectronics.

Another potential application where LEIB on oxides has found a niche is the generation of substrates suitable for resistive switching (RS) processes, which lie at the core of the nonvolatile data-storage memristor technology. This technology, which is still in its exploratory stages, is based in the existence of on/off states which depend on the electrical resistance level of the bit (rather than on the electrostatic or magnetic state). The state can be switched with external electric fields in oxides, which force a metal-insulator transition (MIT). TiO<sup>2</sup> is a good candidate for this purpose, but it has been found that a previous forming step involving LEIB radically improves the RS process [33]. The conductive 2D layer generated consists in a set of self-organized grains with a locally reduced chemical composition. This distribution of grains is a kind of template where bipolar switching between a semiconductor and a metallic state (induced with a conductive atomic force microscopy (AFM) tip) is localized at the grains. **Figure 5** illustrates a progressive chemical reduction and a declining electrical resistance of the TiO<sup>2</sup> surface during ion bombardment. The same type of mechanism has been found in STO substrates [34] with rather low-energy Ar<sup>+</sup> ions (66–200 eV). LEIB greatly improves the resistive memory effect, if compared with pristine STO. In these cases, thermal annealings would create vacancies and would also improve the electrical conductance for RS processes. But the influence of LEIB is limited to the upper layers of the material and can be thus combined with focused ion beam methods or shadow masking to pattern surfaces.

 **Figure 5.** (a) Quantitative changes in the concentration of Ti cations with the sputtering time, as measured with XPS. (b) Behavior of the resistance (in Ω/μm<sup>2</sup> ) of the modified layer with sputtering time. Adapted with permission from Ref. [33].

Oxides exhibit outstanding magnetic properties, and in this respect, LEIB can be employed to modify the microstructure and consequently the magnetic response of the oxide layer. Bilayers of Ni80Fe20/α-Fe<sup>2</sup> O3 grown on SiO<sup>2</sup> exhibit, for instance, a different coercivity depending on the growth mode [35]. If ion bombardment is carried out on the α-Fe<sup>2</sup> O3 layer before the growth of the Ni80Fe20 film, the uniaxial anisotropy of this layer is reduced, exhibiting a lower coercivity. This is a clear example of how the coupling between a ferromagnetic layer and an antiferromagnetic layer can be tuned with LEIB.

Regarding the modification of interfaces, we can mention that STO, in combination with other insulating oxides such as LaAlO3 (LAO), forms a heterostructure with remarkable conductive, magnetic, and superconductive properties at the interface. The surprising origin is the stabilization of an emerging 2D electron gas (2DEG). There are examples of the modification of the 2DEG in LAO/STO interfaces [36]. They have shown that the conducting state at the interface can be completely suppressed with Ar<sup>+</sup> bombardment at 150 eV. The conductivity can be then restored with a high-temperature annealing in the presence of oxygen. The authors suggest that the strain induced by Ar implantation causes local destabilization and partial amorphization. These disordered regions localize defect states where charges are trapped, in the same way as in other types of semiconductors. The annealing eliminates the Ar and heals defects, restoring the initial conductivity. The LEIB approach has been also used to create a 2D electron gas in an Ar<sup>+</sup> -bombarded SrTiO3 (100) surface [37], which shows an increase of the low-temperature magnetoresistance when the magnetic field *H* points away from the surface. This and other effects are closely related to electron confinement, generated in this latter case by the modifications induced by ion bombardment.

### **5. Formation of epitaxial thin films of oxides**

data-storage memristor technology. This technology, which is still in its exploratory stages, is based in the existence of on/off states which depend on the electrical resistance level of the bit (rather than on the electrostatic or magnetic state). The state can be switched with external elec-

for this purpose, but it has been found that a previous forming step involving LEIB radically improves the RS process [33]. The conductive 2D layer generated consists in a set of self-organized grains with a locally reduced chemical composition. This distribution of grains is a kind of template where bipolar switching between a semiconductor and a metallic state (induced with a conductive atomic force microscopy (AFM) tip) is localized at the grains. **Figure 5** illus-

during ion bombardment. The same type of mechanism has been found in STO substrates

effect, if compared with pristine STO. In these cases, thermal annealings would create vacancies and would also improve the electrical conductance for RS processes. But the influence of LEIB is limited to the upper layers of the material and can be thus combined with focused ion

Oxides exhibit outstanding magnetic properties, and in this respect, LEIB can be employed to modify the microstructure and consequently the magnetic response of the oxide layer.

 **Figure 5.** (a) Quantitative changes in the concentration of Ti cations with the sputtering time, as measured with XPS. (b)

the growth of the Ni80Fe20 film, the uniaxial anisotropy of this layer is reduced, exhibiting a lower coercivity. This is a clear example of how the coupling between a ferromagnetic layer

Regarding the modification of interfaces, we can mention that STO, in combination with other

ions (66–200 eV). LEIB greatly improves the resistive memory

exhibit, for instance, a different coercivity depend-

(LAO), forms a heterostructure with remarkable conductive,

) of the modified layer with sputtering time. Adapted with permission from Ref. [33].

O3

layer before

trates a progressive chemical reduction and a declining electrical resistance of the TiO<sup>2</sup>

is a good candidate

surface

tric fields in oxides, which force a metal-insulator transition (MIT). TiO<sup>2</sup>

[34] with rather low-energy Ar<sup>+</sup>

Bilayers of Ni80Fe20/α-Fe<sup>2</sup>

Behavior of the resistance (in Ω/μm<sup>2</sup>

insulating oxides such as LaAlO3

O3

and an antiferromagnetic layer can be tuned with LEIB.

grown on SiO<sup>2</sup>

ing on the growth mode [35]. If ion bombardment is carried out on the α-Fe<sup>2</sup>

beam methods or shadow masking to pattern surfaces.

34 Modern Technologies for Creating the Thin-film Systems and Coatings

We present, in this final section of the chapter, a rather specific but quite surprising capability of ion bombardment, which we have successfully proven in a rutile TiO<sup>2</sup> (110) substrate. As we have seen in previous sections, LEIB is a commonly used process to modify the physicochemical properties of oxide surfaces. In this respect, the literature has usually described the resulting material as a defective and chemically reduced version of the oxide, lacking a profound knowledge of the real structure. As all properties ultimately depend on the structure, the situation is highly undesirable. We show here a case where the structure of the abovementioned ion-bombarded surface is well defined after the modification.

The purpose of this section is to show that the use of LEIB on a surface of a single-crystalline oxide can induce the formation of a single-crystalline and epitaxial thin layer of the corresponding suboxide. For the case of TiO<sup>2</sup> (110), we show here that high doses of Ar<sup>+</sup> bombardment of the single crystal produces a 10 nm thick film of TiO(001) [38].

It is well known that the progressive depletion of oxygen in polycrystalline materials during ion bombardment can, in some cases, induce the formation of a new phase with the cation in a lower oxidation state. In other words, the formation of a crystalline suboxide (i.e., with well-defined Bragg reflections) can be promoted. The main message of this section is that if the starting point is a single-crystalline surface (instead of ion bombarding a polycrystalline material) and if there is some kind of structural affinity or matching with a particular crystallographic orientation of the suboxide, this suboxide can nucleate and grow epitaxially coupled to the original oxide, being a thin epitaxial film the outcome of the process.

#### **5.1. Formation of a TiO(001) thin film on TiO<sup>2</sup> (110)**

Titanium dioxide is, with its different structures, a wideband semiconductor with many applications in technology. It has been also, along the years, a benchmark for the fundamental study of oxide surfaces. In fact, as we mentioned before, TiO<sup>2</sup> (110) is the most studied oxide surface [39]. Regarding its optical properties, it is very transparent and has a high refractive index, which together with its stability, non-toxicity and the easy synthesis of small particles, justifies its extensive use as a white pigment. It also displays a very interesting photocatalytic activity. This property, along with its particular interaction with water, has fostered the investigation and the use of TiO<sup>2</sup> as a self-cleaning coating [40]. As we have previously discussed, also in the field of resistive switching, the investigation with TiO<sup>2</sup> has found a satisfactory feedback.

The starting point of the modification we describe here is a rutile TiO<sup>2</sup> (110) clean and flat surface, which is ion bombarded with Ar<sup>+</sup> at 3 keV at room temperature with doses up to 8 × 1016 ions cm−2. Auger electron spectroscopy (AES) shows a clear chemical reduction: a decrease of the O/Ti ratio and a shift of the TiLMM and TiLMV transitions to higher kinetic energies. The incidence of the ion beam is normal, and so the modified surfaces show a rough topography, but with no recognizable pattern or symmetry as seen with atomic force microscopy (AFM). However, X-ray diffraction of the bombarded surfaces already shows the emergence of a new reflection (not present in the pristine sample) compatible with the cubic rock salt titanium monoxide phase with its (001) crystallographic direction oriented along the surface normal (**Figure 6a**). Low-energy electron diffraction (LEED) measurements performed before and after the modification show the transformation of the rectangular surface diffraction pattern corresponding to TiO<sup>2</sup> (110) to a different diffuse LEED pattern displaying square symmetry, with its main directions rotated 45° with respect to the TiO<sup>2</sup> (110) directions (**Figure 6b**–**d**).

Grazing incidence XRD (GIXRD) measurements help determine the crystalline orientations of both the dioxide and monoxide phases. HK scans in reciprocal space (those which explore the in-plane directions) show all the allowed reflections (not subject to extinction rules) of TiO(001) and TiO<sup>2</sup> (110), confirming the 45° rotation of the square surface lattice of TiO(001) with respect to the rectangular surface lattice of TiO<sup>2</sup> (110) (blue and black lattices, respectively, of **Figure 6e**). L scans (to explore the out-of-plane direction) determine that the periodicity of the TiO phase is indeed the corresponding one. These observations confirm the ion-induced transformation of the upper layers of TiO<sup>2</sup> (110) into TiO(001), with the particularity that both lattices are in registry. Being both lattices rotated 45°, the matching is almost perfect along the [001] direction of TiO<sup>2</sup> (110). Along the [1–10] the mismatch is large (about 10%), which clarifies the observations carried out with LEED, XRD, and scanning transmission electronic microscopy (STEM), all of them revealing the tilting of the lattice along that direction. This tilting, probably stabilized by the existence of misfit dislocations, helps relieve that large mismatch along that particular direction. STEM cross-sectional images of the modified layer show the homogeneity of the modified layer, revealing a constant thickness of 10 nm and its relatively good crystallinity.

Density functional theory (DFT) calculations confirm the high stability of the interface for a particular structural configuration, with an energy estimated to be *γint* = 3.2 J m−2. The good structural and chemical matchings justify this low value. Also, the calculations suggest that a contraction of the TiO bonds at the interface (the out-of-plane TiO distance slightly increases with the interface distance) is due to the charge transfer from TiO to TiO<sup>2</sup> , resulting in an interface with metallic character.

Modification of Oxide Thin Films with Low-Energy Ion Bombardment http://dx.doi.org/10.5772/66651 37

[39]. Regarding its optical properties, it is very transparent and has a high refractive index, which together with its stability, non-toxicity and the easy synthesis of small particles, justifies its extensive use as a white pigment. It also displays a very interesting photocatalytic activity. This property, along with its particular interaction with water, has fostered the investigation

8 × 1016 ions cm−2. Auger electron spectroscopy (AES) shows a clear chemical reduction: a decrease of the O/Ti ratio and a shift of the TiLMM and TiLMV transitions to higher kinetic energies. The incidence of the ion beam is normal, and so the modified surfaces show a rough topography, but with no recognizable pattern or symmetry as seen with atomic force microscopy (AFM). However, X-ray diffraction of the bombarded surfaces already shows the emergence of a new reflection (not present in the pristine sample) compatible with the cubic rock salt titanium monoxide phase with its (001) crystallographic direction oriented along the surface normal (**Figure 6a**). Low-energy electron diffraction (LEED) measurements performed before and after the modification show the transformation of the rect-

pattern displaying square symmetry, with its main directions rotated 45° with respect to the

Grazing incidence XRD (GIXRD) measurements help determine the crystalline orientations of both the dioxide and monoxide phases. HK scans in reciprocal space (those which explore the in-plane directions) show all the allowed reflections (not subject to extinction rules) of

of **Figure 6e**). L scans (to explore the out-of-plane direction) determine that the periodicity of the TiO phase is indeed the corresponding one. These observations confirm the ion-induced

lattices are in registry. Being both lattices rotated 45°, the matching is almost perfect along

clarifies the observations carried out with LEED, XRD, and scanning transmission electronic microscopy (STEM), all of them revealing the tilting of the lattice along that direction. This tilting, probably stabilized by the existence of misfit dislocations, helps relieve that large mismatch along that particular direction. STEM cross-sectional images of the modified layer show the homogeneity of the modified layer, revealing a constant thickness of 10 nm and its

Density functional theory (DFT) calculations confirm the high stability of the interface for a particular structural configuration, with an energy estimated to be *γint* = 3.2 J m−2. The good structural and chemical matchings justify this low value. Also, the calculations suggest that a contraction of the TiO bonds at the interface (the out-of-plane TiO distance slightly increases

with the interface distance) is due to the charge transfer from TiO to TiO<sup>2</sup>

(110), confirming the 45° rotation of the square surface lattice of TiO(001)

(110). Along the [1–10] the mismatch is large (about 10%), which

as a self-cleaning coating [40]. As we have previously discussed, also in

has found a satisfactory feedback.

(110) to a different diffuse LEED

(110) (blue and black lattices, respectively,

, resulting in an

(110) into TiO(001), with the particularity that both

at 3 keV at room temperature with doses up to

(110) clean and flat

and the use of TiO<sup>2</sup>

TiO<sup>2</sup>

TiO(001) and TiO<sup>2</sup>

the [001] direction of TiO<sup>2</sup>

relatively good crystallinity.

interface with metallic character.

the field of resistive switching, the investigation with TiO<sup>2</sup>

36 Modern Technologies for Creating the Thin-film Systems and Coatings

angular surface diffraction pattern corresponding to TiO<sup>2</sup>

with respect to the rectangular surface lattice of TiO<sup>2</sup>

surface, which is ion bombarded with Ar<sup>+</sup>

(110) directions (**Figure 6b**–**d**).

transformation of the upper layers of TiO<sup>2</sup>

The starting point of the modification we describe here is a rutile TiO<sup>2</sup>

 **Figure 6.** (a) *θ−2θ* scan of a TiO<sup>2</sup> (110) single crystal after ion bombardment. A new reflection from TiO(002), not present before, emerges. (b) LEED image of the TiO<sup>2</sup> (110) surface before LEIB, (c) during LEIB, and (d) after LEIB and a soft thermal annealing. New spots (marked with arrows), with a different symmetry, appear. They become sharper after the annealing. All diffractograms are taken at *E* = 63 eV. (e) X-ray diffraction HK map in reciprocal space of the ionbombarded TiO<sup>2</sup> (110) surface, where both sets of reflections are seen: black for TiO<sup>2</sup> (110) and blue for TiO(001). The two surface lattices, with their corresponding lattice vectors, are marked. (f) Scanning transmission electron microscopy (STEM) image of the irradiated layer. The transformed surface is seen. (g) Higher magnification STEM image and fast Fourier transform (top right) of the image. The yellow rectangle is a detail of the interface. Adapted with permission from Ref. [38].

All results together indicate that the initial stages of ion bombardment induce the formation of a chemically reduced and defective version of the dioxide at the initial stages, which transforms into a very disordered (or even amorphous) phase for intermediate doses, as the rutile structure is not stable for a high density of defects. The absence of a LEED pattern after medium doses supports the existence of this disordered intermediate phase. At sufficiently high doses, while preferential oxygen sputtering continues operative, the TiO(001) phase emerges, favored by the good registry between both oxides at the interface.

An interesting fact is that the thickness of the modified layer (10 nm) is larger than the average depth where most of the ion damage is generated (around 4 nm, according to SRIM simulations). This implies that diffusion is active, probably assisted by radiation, and also enhanced by a locally higher temperature due to the dissipation of the energetic ions. Most probably, mass transport is governed by interstitial Ti cations [41, 42] generated by ion bombardment, which migrate to the interface transforming it locally into TiO. The interface thus advances toward the bulk. When the diffusion length of the species is not sufficiently long to compensate for the receding surface (atoms are being continuously sputtered by the ion beam), then a steady state is reached, and the thickness of the modified layer saturates.

### **Acknowledgements**

I would like to thank all my former and present collaborators, especially those participating in the works reviewed in the last section of this chapter. Financial support from Projects MAT2012-38045-C04-03 and FIS2014-61839-EXP from the Spanish MINECO is also acknowledged.

### **Author details**

Oscar Rodríguez de la Fuente

Address all correspondence to: oscar.rodriguez@fis.ucm.es

Materials Physics Department, Universidad Complutense de Madrid, Madrid, Spain

### **References**


An interesting fact is that the thickness of the modified layer (10 nm) is larger than the average depth where most of the ion damage is generated (around 4 nm, according to SRIM simulations). This implies that diffusion is active, probably assisted by radiation, and also enhanced by a locally higher temperature due to the dissipation of the energetic ions. Most probably, mass transport is governed by interstitial Ti cations [41, 42] generated by ion bombardment, which migrate to the interface transforming it locally into TiO. The interface thus advances toward the bulk. When the diffusion length of the species is not sufficiently long to compensate for the receding surface (atoms are being continuously sputtered by the ion beam), then

I would like to thank all my former and present collaborators, especially those participating in the works reviewed in the last section of this chapter. Financial support from Projects MAT2012-38045-C04-03 and FIS2014-61839-EXP from the Spanish MINECO is also

Materials Physics Department, Universidad Complutense de Madrid, Madrid, Spain

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### **Conventional and Un-Conventional Lithography for Fabricating Thin Film Functional Devices Conventional and Un-Conventional Lithography for Fabricating Thin Film Functional Devices**

Abdelhanin Aassime and Frederic Hamouda Abdelhanin Aassime and Frederic Hamouda

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

http://dx.doi.org/10.5772/66028

#### **Abstract**

Thin film devices are conquering many aspects of today's life, and continuous shrinking of building block dimensions of these structures enhances their performances and makes them economically attractive. This chapter is an overview of some conventional and unconventional lithography techniques used to fabricate thin film functional structures. Several aspects of pattern transfer were addressed with emphasis on the limits of these lithography techniques. We have thus highlighted the issue of pitch resolution for optical lithography and discussed some aspect of proximity effects for electron beam lithography. Pattern transfer from resist image to the wafer was also discussed. Considered as unconventional, we discussed several aspects linked to thin film fabrication using nanoimprint and nanosphere lithography techniques.

**Keywords:** optical lithography, electron beam lithography, nanoimprint lithography, nanosphere lithography, liftoff

### **1. Introduction**

The continuous trend toward miniaturized and high-performance systems has been leading research and development in novel materials and devices with superior and new functionalities. In this regard, the high sensitivity of modern technologies at submicron scale opens prospects for realization of thin film functional devices such as capacitors for power components [1], sensing devices for biomedical applications [2] and magnetic thin films structures for data processing technology [3]. These devices were mostly fabricated following top-down view where deposition techniques were combined with lithography and eventually etching. In this chapter, we will mainly focus on depicting conventional and un-conventional lithog-

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

raphy (see **Figure 1**) used for fabricating thin film functional devices when combined with physical vapor deposition (PVD) technique.

**Figure 1.** Scheme showing different lithography techniques.

The first part of this chapter reports on most known conventional lithography techniques. These techniques consist of transferring a given pattern to a thin resist by means of local interaction between the resist and the beam of particles like photons or electrons. Once the resist is structured, functional devices can be obtained by liftoff technique after PVD deposition of thin films [4] or by etching [5]. Existing conventional lithography can be divided in two categories depending on the way of writing (transferring) the pattern: parallel and series writing methods. For parallel method such as optical lithography with different wavelengths, transfer to resist is done at once by using a mask containing the pattern. This technique has the advantage of being fast but limited by the mask fabrication step and diffraction effects. For series writing (maskless), such as electron beam lithography (EBL) and laser lithography, pattern is transferred to the resist pixel-by-pixel which makes them more precise but unfortunately slower. In the industry, parallel lithography technique is commonly used for mass production and serial method is mostly used for mask making. The aim of this first part is to present the principle of these various fabrication techniques with the focus on pitch resolution for optical lithography and proximity effects for electron beam lithography (EBL). Pattern transfer to the wafer using both lithography techniques will also be addressed.

In this second part, we report alternative techniques considered as unconventional lithography. Among them, we find nanoimprint lithography (NIL). This method is based on the printing patterns in a polymer with a mold which can be rigid or flexible, and after separation, transfer the pattern in the substrate. The first results were achieved in 1995 with the thermal NIL and rigid molds, more commonly known as hot embossing [6]. It permits to carry out high-density nanostructures on large areas and at low cost. A second method based on UV has been developed in 1996 at Philips Research Labs, which offers the advantage to work at room temperature and low pressure [7]. This technique uses the photon energy to crosslink the resist. Molds are transparent and can be flexible or rigid. Then, two ways have been developed: soft UV-NIL and hard UV-NIL. Each has advantages and will be addressed in this chapter with different steps to implement these techniques. Application examples will be also presented.

Among emerging methods, we find nanosphere lithography (NSL), also called as colloidal lithography. It is a low-cost simple technique to implement and permits to pattern well-ordered 2D nanoparticle arrays on large surface [8]. Another advantage of the NSL is to be a good candidate for the fabrication of diameter tunable nanoparticles in a wide range of 20–1000 nm. We will present the different technical steps to obtain a two-dimensional colloid mask.

### **2. Optical lithography**

raphy (see **Figure 1**) used for fabricating thin film functional devices when combined with

The first part of this chapter reports on most known conventional lithography techniques. These techniques consist of transferring a given pattern to a thin resist by means of local interaction between the resist and the beam of particles like photons or electrons. Once the resist is structured, functional devices can be obtained by liftoff technique after PVD deposition of thin films [4] or by etching [5]. Existing conventional lithography can be divided in two categories depending on the way of writing (transferring) the pattern: parallel and series writing methods. For parallel method such as optical lithography with different wavelengths, transfer to resist is done at once by using a mask containing the pattern. This technique has the advantage of being fast but limited by the mask fabrication step and diffraction effects. For series writing (maskless), such as electron beam lithography (EBL) and laser lithography, pattern is transferred to the resist pixel-by-pixel which makes them more precise but unfortunately slower. In the industry, parallel lithography technique is commonly used for mass production and serial method is mostly used for mask making. The aim of this first part is to present the principle of these various fabrication techniques with the focus on pitch resolution for optical lithography and proximity effects for electron beam lithography (EBL). Pattern

transfer to the wafer using both lithography techniques will also be addressed.

In this second part, we report alternative techniques considered as unconventional lithography. Among them, we find nanoimprint lithography (NIL). This method is based on the printing patterns in a polymer with a mold which can be rigid or flexible, and after separation, transfer the pattern in the substrate. The first results were achieved in 1995 with the thermal NIL and rigid molds, more commonly known as hot embossing [6]. It permits to carry out high-density nanostructures on large areas and at low cost. A second method based on UV has been developed in 1996 at Philips Research Labs, which offers the advantage to work at room

physical vapor deposition (PVD) technique.

44 Modern Technologies for Creating the Thin-film Systems and Coatings

**Figure 1.** Scheme showing different lithography techniques.

Optical lithography consists mainly on a light source illuminating, through an ensemble of optical lenses and apertures, a mask containing pattern aimed to be transferred to a given substrate. The mask can be set in the vicinity of the substrate (contact/proximity mode) or at a certain distance from it (projection mode). The mask is called "binary" in the sense that the light either passes through metal-free area (1) or being reflected in metal-covered surface of the mask (0). Pattern transfer is intermediated by a specific resist which covers the substrate. The interaction between light and resist results in local modification of molecules arrangement of the resist which can be revealed in a specific solution (development process). **Figure 2** depicts the mask replication in contact/proximity and projection modes. The contact mode is a configuration where the mask is directly touching the resist, whereas a gap of few microns separates them in the proximity mode.

**Figure 2.** Mask replication in contact/proximity mode (a) and projection configuration (b). The condenser collects light from the source and illuminates the mask pattern. An additional imaging lens is needed to de-magnify the mask pattern (up to ∼1:4).

The fundamental limit of optical lithography is not determined by the optical system alone but rather by an overall contribution from the optics, resist and subsequent process steps. As depicted in **Figure 3**, there exist in general two kinds of resolution, one which is linked to feature size and the other to pitch. While the feature size determines the critical dimension that can be obtained (e.g., size of the transistor), the pitch determines its density on the wafer (e.g., number of transistors per wafer).

**Figure 3.** Example of pattern to be exposed showing a hole diameter as smallest feature and related pitch.

For clarity, we will focus only on projection systems in this paragraph. The pitch resolution (R) of these systems is usually expressed in terms of source wavelength *λ* and numerical aperture ( = ()) as

$$R = k1 \frac{\mathcal{X}}{n \sin \theta} \tag{1}$$

where *k1* a process-dependent constant with values in 0.5–1 range, *n* is the refractive index of the media between the mask and the wafer, and *θ* is half acceptance angle of the lens (see **Figure 2**).

To improve the pitch resolution, it is necessary to decrease (*λ, k*1) and to enhance the numerical aperture. Historically, this improvement was driven by decrease in wavelength *λ* of the source. From mercury lamps at 365nm–435nm wavelength, to excimer laser sources with Krypton Fluoride (248 nm), Argon Fluoride (193 nm), and molecular Fluorine (157 nm) [9]. Furthermore, optical resolution limit has been pushed toward sub-100 nm features using resolution enhancement techniques (RET) such as optical proximity correction, high numerical aperture, and phase-shift masks; 65 nm device geometry (nodes) was indeed achieved using wavelength as large as 193 nm [10, 11]. While numerical aperture higher than 1 is not possible in conventional air-media imaging, using water between the last imaging lens and the wafer has pushed down the limit to sub-45 nm [12]. This technique, known as immersion lithography, is a potential candidate to take over the actual 193 nm technology in the industry. To complete the picture, extreme UV (EUV) is another immerging technique on which relies the future of next generation of circuit components to push further resolution down to sub-20 nm [13].
