**5. Unconventional lithography techniques**

**Figure 8.** Scanning electron microscope images of resist profile (a,b) and line metallic structure (c,d). Images (a) and (b) were obtained using a trilayer of SML/ZEP520A/PMMA electron beam resists. Images (c) and (d) correspond to liftoff process of 100-nm-thick gold using SML electron beam resist. Raith 150 ebeam system was used to expose resists.

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

Overlay comes into play when several fabrication levels are needed to fabricate a given structure. Overlay error determines the precision at which a pattern is placed accurately in wanted position on the chip. Mostly, placement accuracies of sub-30 nm up to 1 µm are required to connect micrometric or nanometric scale elements to the pads or between other elements of the structure. Alignment marks are usually used to achieve such precisions, and the patternplacement accuracy increases when the marks are as close as possible to the critical areas. Alignment system consists generally of a detector, which detects a given mark, and a software that analyzes and determines the center of the mark relatively to a reference. The accuracy depends not only on the detector signal but also on the quality of the marks. Depending on subsequent process requirements, these marks can be made from resist, metal or being etched. **Figure 9** illustrates transferred ellipsoidal on silicon waveguide gold obtained by liftoff of 30 nm-thick gold. An alignment of <30 nm was required to properly align gold nanostructure on

**Figure 9.** Scanning electron image of 11 ellipsoidal gold on Si waveguide for guided plasmonic applications.

the waveguide [21].

As we have previously shown, lithography is a key step in the processes of micro and nanotechnologies. The aim is to structure a pattern in a polymer that has been deposited on a substrate. The polymer is generally an organic or inorganic resist. The structure will allow to realize other steps such as etching or materials' deposit. Currently, this technology use conventional lithography techniques like optical lithography or electron beam lithography (EBL). Beyond these technologies, other methods are referenced as un-conventional to reduce the cost of production and permit large series. Among them, we have the nanoimprint lithography (NIL) (see **Figure 1**). This technique is based on printing in a polymer by using a mold which may be rigid or flexible. This method was developed in the 1990s because it allowed to obtain rapidly large area nanoscale patterns with low costs. The main steps are printing in the polymer with a mold, demolding, and transferring the pattern into the substrate. As shown in **Figure 10**, there are two main techniques of NIL. The first developed method is more commonly known as hot embossing thermal process. S.Y. Chou published the first results of this technique in 1995 [6].

**Figure 10.** Principle of Thermal NIL (a) and UV-NIL (b).

A few years later, a second method called UV-NIL was developed in the Philips Research Labs [7]. In this case, the photon energy is used to cure the photosensitive resist. This process requires transparent mold and offers other advantages than thermal NIL. In the following paragraphs, we will present NIL techniques, mainly thermal and UV-NIL, how to design the mold, and we will also discuss about a new method called nanosphere lithography (NSL).

### **5.1. Thermal NIL**

The principle of thermal NIL is to imprint in a thermoplastic polymer with a structured rigid mold [6]. The mold needs to have an antisticking treatment to avoid lifting printed patterns. This process uses a polymer heated at a temperature above its glass transition temperature (Tg) and a pressure between 10 and 200 bars during the imprint [22]. Thus, the mold is removed after cooling the substrate. **Figure 10a** shows the main steps of this method which provides resolution in the nanometer range. However, it has the drawbacks to operate with high temperatures and high pressures.

**Figure 11.** (a) Process scheme for Si mold developed by EBL and HSQ resist and (b) SEM image of a silicon mold for nanogap electrodes (150 nm).

The fundamental step is to design the rigid mold; the materials used are mainly Si, SiO2, or quartz; and the pattern is realized by an electron beam lithography to achieve high resolution. **Figure 11a** shows an example scheme to process Si mold developed for nanogap electrodes with high-resolution HSQ negative-tone resist. **Figure 11b** is a SEM image of master mold (Si) obtained with this process for a gap of 150 nm and usable for thermal NIL. More complete details of the thermal NIL process are given in Ref. [23].

### **5.2. UV-NIL**

The cross-linking of the resist with UV nanoimprint lithography is obtained by the photon energy. The mold can be flexible or rigid and necessarily transparent. **Figure 10b** summarizes this method used for both soft UV-NIL with a flexible mold and hard UV-NIL with a rigid mold in quartz. The UV transparent mold is imprint in a low-viscosity UV-curable resist UV at room temperature and low pressure between 0 and 1 bar [24]. The mold is firstly removed followed by etching of residual resist layer to permit a transfer of patterns in the substrate by liftoff technique or etching process.

For hard UV-NIL, the main goal is to design by ebeam lithography, the mold generally made in quartz. It is necessary to use a metallic or dielectric mask to control form factor of the pattern during the etch process [25, 26]. Before imprinting in UV-curable resist, a special treatment is applied to reduce surface energy of the mold. An antisticking is needed to avoid resist tear-off during "demold" step between mold and substrate [27]. The next step is to etch the residual layer and transfer in the substrate.

With soft UV-NIL, the flexible mold is generally in poly(dimethylsiloxane) PDMS [28] and obtained from a Si master mold. PDMS offers good chemical stability and high optical transparency [29]. The method to get the master mold is the same as described in the thermal NIL process. The next step is an antisticking treatment of the master mold, deposits PDMS mixed with his curing agent on master mold and bakes at 60° during 2 h. A full description of the process is made in reference [30]. **Figure 12** shows an example of dots AFM image in PDMS stamp and stamp PDMS imprint in a UV-curable resist from AMO for gold nanoparticle applications. An example of 30-nm-thick gold array for plasmonic application is illustrated in **Figure 12c**.

**Figure 12.** (a) AFM image of dots in PDMS stamp with a periodicity of 500 nm and height of 90 nm, (b) SEM image of an imprint in UV-curable resist (AMONIL from AMO GMBH) with PDMS flexible stamp obtained from the Si master mold and (c) 30-nm-thick gold nanoparticles array obtained by UV-NIL for plasmonic applications.

### **5.3. Nanosphere lithography**

**5.1. Thermal NIL**

temperatures and high pressures.

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

nanogap electrodes (150 nm).

**5.2. UV-NIL**

details of the thermal NIL process are given in Ref. [23].

liftoff technique or etching process.

layer and transfer in the substrate.

The principle of thermal NIL is to imprint in a thermoplastic polymer with a structured rigid mold [6]. The mold needs to have an antisticking treatment to avoid lifting printed patterns. This process uses a polymer heated at a temperature above its glass transition temperature (Tg) and a pressure between 10 and 200 bars during the imprint [22]. Thus, the mold is removed after cooling the substrate. **Figure 10a** shows the main steps of this method which provides resolution in the nanometer range. However, it has the drawbacks to operate with high

**Figure 11.** (a) Process scheme for Si mold developed by EBL and HSQ resist and (b) SEM image of a silicon mold for

The fundamental step is to design the rigid mold; the materials used are mainly Si, SiO2, or quartz; and the pattern is realized by an electron beam lithography to achieve high resolution. **Figure 11a** shows an example scheme to process Si mold developed for nanogap electrodes with high-resolution HSQ negative-tone resist. **Figure 11b** is a SEM image of master mold (Si) obtained with this process for a gap of 150 nm and usable for thermal NIL. More complete

The cross-linking of the resist with UV nanoimprint lithography is obtained by the photon energy. The mold can be flexible or rigid and necessarily transparent. **Figure 10b** summarizes this method used for both soft UV-NIL with a flexible mold and hard UV-NIL with a rigid mold in quartz. The UV transparent mold is imprint in a low-viscosity UV-curable resist UV at room temperature and low pressure between 0 and 1 bar [24]. The mold is firstly removed followed by etching of residual resist layer to permit a transfer of patterns in the substrate by

For hard UV-NIL, the main goal is to design by ebeam lithography, the mold generally made in quartz. It is necessary to use a metallic or dielectric mask to control form factor of the pattern during the etch process [25, 26]. Before imprinting in UV-curable resist, a special treatment is applied to reduce surface energy of the mold. An antisticking is needed to avoid resist tear-off during "demold" step between mold and substrate [27]. The next step is to etch the residual

With soft UV-NIL, the flexible mold is generally in poly(dimethylsiloxane) PDMS [28] and obtained from a Si master mold. PDMS offers good chemical stability and high optical Nanosphere lithography (NSL) is a simple technique to implement and inexpensive. It is also called as colloidal lithography and allows well-ordered nanoparticles in a plane and on large surface. The structure obtained by NSL based on a self-assembling nanosphere achieves a colloidal mask in two dimensions (**Figure 13**). This technique has demonstrated to be well suited for the fabrication of size-tunable nanoparticles in the 20–1000 nm range [8]. This method can also be used to obtain silicon mold for NIL application. To file a nanosphere solution onto the substrate, several methods exist as spin coating, [31] drop coating file [32], and thermo-electrically cooled angle coating [33]. Nanospheres meet into 2D hexagonal mesh on the substrate due to capillary forces during the solvent evaporation. After the step of deposition of self-assembling nanospheres, a thickness material layer is evaporated by electron beam through the nanosphere mask. Then, nanosphere mask is removed with a solvent.

NSL mask fabrication may depend on the number of layers required to obtain nanostructure networks [34]. **Figure 14** shows an example with a single and a double NSL layers and nanostructure arrays obtained with theses configurations. When a gold layer is deposited through a single monolayer by self-assembled nanospheres onto the substrate and NSL mask is removed, an array of triangular nanoparticles is obtained (**Figure 14**, left). For a configuration with two monolayers of nanospheres deposited and assembled, this is obtained by an increasing of the nanosphere concentration. When the second layer is assembled on the first, in order to obtain a significant part of a double layer of hexagonally assembled nanospheres, the free interstices where the material can be deposited on the substrate form an homogeneous pattern of hexagonal nanoparticles (**Figure 14**, right).

**Figure 13.** Principle of nanosphere lithography.

**Figure 14.** Schematic representation of gold nanostructures obtained with a single (left) and double layer (right).

### **6. Conclusion**

Combined with deposition techniques, conventional lithography and un-conventional lithography are powerful tools to fabricate thin film functional devices. These tools offer tremendous opportunities to enhance not only device performances and reduce effective cost, but also to discover and explore new functionalities. It provides, in particular, the possibility to better control of pattern transfer in terms of resolution, density, and resist profile.

### **Acknowledgements**

to obtain a significant part of a double layer of hexagonally assembled nanospheres, the free interstices where the material can be deposited on the substrate form an homogeneous pattern

**Figure 14.** Schematic representation of gold nanostructures obtained with a single (left) and double layer (right).

Combined with deposition techniques, conventional lithography and un-conventional lithography are powerful tools to fabricate thin film functional devices. These tools offer

of hexagonal nanoparticles (**Figure 14**, right).

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

**Figure 13.** Principle of nanosphere lithography.

**6. Conclusion**

This work was done within the Center for Nanoscience and Nanotechnology clean room and partly supported by the RENATECH network and the General Council of Essonne.

### **Author details**

Abdelhanin Aassime\* and Frederic Hamouda

\*Address all correspondence to: abdelhanin.aassime@u-psud.fr

Centre de Nanosciences et de Nanotechnologies, C2N-Site d'Orsay, Université Paris-Sud, Orsay Cedex, France

### **References**


lithography. Microelectronic Engineering. 1993;21(1–4):293–296. doi:10.1016/0167-9317 (93)90076-H

[19] Rooks M, Belic N, Kratschmer E, Viswanathan R. Experimental optimization of the electron-beam proximity effect forward scattering parameter. Journal of Vacuum Science and Technology B. 2005;23(6):2769–2774. doi:10.1116/1.2062431

[6] Chou SY, Krauss PR, Renstrom PJ. Imprint of sub-25 nm vias and trenches in polymers.

[7] Haisma J, Verheijen M, VandenHeuvel K, VandenBerg J. Mold-assisted nanolithography: a process for reliable pattern replication. Journal of Vacuum Science and Technol-

[8] Haynes CL, Van Duyne RP. Nanosphere lithography: a versatile nanofabrication tool for studies of size-dependent nanoparticle optics. Journal of Physical Chemistry B.

[9] Mizoguchi H, Saitoh T, Matsunaga T. Development of light sources for lithography at

[10] Trouiller Y. From 120 to 32 nm CMOS technology: development of OPC and RET to rescue optical lithography. Comptes Rendus Physique. 2006;7(8):887–895. doi:10.1016/

[11] Pease RF, Chou SY. Lithography and other patterning techniques for future electronics. Proceedings of the IEEE. 2008;96(2):248–270. doi:10.1109/JPROC.2007.911853

[12] Bae WJ, Trikeriotis M, Rodrigues R, Zettel MF, Piscani E, Ober CK, Giannelis EP, Zimmerman P. High index nanocomposite photoresist for 193 nm lithography. Advances in Resist Materials and Processing Technology XXVI. 2009;7273:727326. doi:

[13] Kulmala TS, Vockenhuber M, Buitrago E, Fallica R, Ekinci Y. Toward 10 nm half-pitch in extreme ultraviolet lithography: results on resist screening and pattern collapse mitigation techniques. Journal of Micro-Nanolithography MEMS and MOEMS.

[14] Harriott LR, Berger SD, Biddick C, Blakey MI, Bowler SW, Brady K, Camarda RM, Connelly WF, et al. The SCALPEL proof of concept system. Microelectronic Engineer-

[15] Dhaliwal RS, Enichen WA, Golladay SD, Gordon MS, Kendall RA, Lieberman JE, Pfeiffer HC, Pinckney DJ, Robinson CF, Rockrohr JD, Stickel W, Tressler EV. Prevail electron projection technology approach for next-generation lithography. IBM Journal

[16] Broers AN. Resolution limits for electron beam lithography. IBM Journal of Research

[17] Vieu C, Carcenac F, Pepin A, Chen Y, Mejias M, Lebib A, Manin-Ferlazzo L, Couraud L, Launois H. Electron beam lithography: resolution limits and applications. Applied

[18] Dubonos SV, Gaifullin BN, Raith HF, Svintsov AA, Zaitsev SI. Evaluation, verification and error determination of proximity parameters alpha, beta and eta in electron beam

Surface Science. 2000;164:111–117. doi:10.1016/S0169-4332(00)00352-4

Applied Physics Letters. 1995;67(21):3114–3116. doi:10.1063/1.114851

present and for the future. Komatsu Technical Report. 2013;59(166).

ogy B. 1996;14(6):4124–4128. doi:10.1116/1.588604

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

2001;105(24):5599–5611. doi:10.1021/jp010657m

2015;14(3):033507. doi:10.1117/1.JMM.14.3.033507

of Research and Development. 2001;45(5):615–638.

and Development. 1988;32(4):502–513.

ing. 1997;35(1–4):477–480. doi:10.1016/S0167-9317(96)00189-X

j.crhy.2006.10.001

10.1117/12.814154


bodies. Microelectronic Engineering. 2010;87(5):1001–1004. doi:10.1016/j.mee.2009. 11.114


#### **In Situ Engineering and Characterization of Correlated Materials with Integrated OMBE–ARPES In Situ Engineering and Characterization of Correlated Materials with Integrated OMBE–ARPES**

Dawei Shen, Haifeng Yang and Zhengtai Liu Dawei Shen, Haifeng Yang and Zhengtai Liu

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

http://dx.doi.org/10.5772/65711

### **Abstract**

bodies. Microelectronic Engineering. 2010;87(5):1001–1004. doi:10.1016/j.mee.2009.

[31] Hulteen JC, Van Duyne RP. Nanosphere lithography: a materials general fabrication process for periodic particle array surfaces. Journal of Vacuum Science and Technology

[32] Hulteen JC, Treichel DA, Smith MT, Duval ML, Jensen TR, Van Duyne RP. Nanosphere lithography: size-tunable silver nanoparticle and surface cluster arrays. Journal of

[33] Micheletto R, Fukuda H, Ohtsu M. A simple method for the production of a twodimensional, ordered array of small latex particles. Langmuir. 1995;11(9):3333–3337.

[34] Kadiri H, Kostcheev S, Turover D, Salas-Montiel R, Nomenyo K, Gokarna A, Lerondel G. Topology assisted self-organization of colloidal nanoparticles: application to 2D large-scale nanomastering. Beilstein Journal of Nanotechnology. 2014;5(XX):1203–1209.

Physical Chemistry B. 1999;103(19):3854–3863. doi:10.1021/jp9904771

A. 1995;13(XX):1553–1558. doi:10.1116/1.579726

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

11.114

doi:10.1021/la00009a012

doi:10.3762/bjnano.5.132

Oxide molecular beam epitaxy has emerged as an effective technique to fabricate complex oxide thin films and novel superlattices with atomic‐level precision. In this chapter, we first briefly introduce the oxide molecular beam epitaxy technique and then show how to use this technique to achieve high‐quality thin films with good stoichi‐ ometry. Moreover, we exhibit that the combination of oxide molecular beam epitaxy and *in situ* angle‐resolved photoemission spectroscopy is indeed a versatile toolkit to tailor and characterize properties of novel quantum materials.

**Keywords:** oxide molecular beam epitaxy (OMBE), correlated materials, angle‐re‐ solved photoemission spectroscopy (ARPES)

### **1. Introduction**

In transition metal oxides, the subtle interplay among charge, orbital, lattice and spin degrees of freedom gives rise to a spectrum of fascinating physical phenomena, including high‐ temperature superconductivity [1], metal‐insulator transition [2], colossal magnetoresistance [3], and so on. Remarkably, in thin film interfaces and ultrathin films of correlated oxides, emergent physics, which does not exist in bulk crystals, occurs [4, 5]. As a well‐known example, two‐dimensional electron gas with high mobility amazingly emerges at the interface of two‐ band insulators LaAlO3 and SrTiO3 [6]. This emergent electron gas was even found to be superconducting [7]. Another example is that strong ferroelectricity and ferromagnetism were found in EuTiO3/DyScO3 superlattices [8]. As Nobel laureate Herbert Kroemer said that 'the interface is the device' [9], these emergent physics may potentially revolutionize our modern technologies.

© 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.

In order to access these thin film‐based physics, the first and most important step is to grow these oxide thin film structures with high quality. This needs exquisite control of growth, and usually is very challenging. In the past several decades, fortunately, reactive oxide molecular beam epitaxy (OMBE) has been proved to be an effective technique in the growth of some oxides with high quality, though being not easy [10–12]. Recently, the *in situ* combined system of OMBE and angle‐resolved photoemission spectroscopy (ARPES) [13–17] have shown great potential in exploring intricate many‐body physics based on oxide film structures, which further intensifies the power of OMBE.

In this chapter, we first present basics of OMBE technique. Then, we show how to grow high‐ quality films with good stoichiometry, and the power of the integrated OMBE‐ARPES in studying and designing many‐body interactions in complex oxides.

### **2. Basics of OMBE technique**

MBE is a vacuum deposition technique in which well‐defined thermal beams of atoms or molecules react at a crystalline surface to produce an epitaxial film. Originally, it was developed to fabricate GaAs and (Al, Ga)As films [18], and soon successfully expanded to other semi‐ conductors as well as metals and insulators. In addition to molecular beams coming from individual heated element source, gas molecular may also be introduced into MBE. Including gas oxidants (e.g., oxygen or ozone) can make an OMBE, which is now applied to grow oxides [11, 12].

In 1985, Betts and Pitt began to use this technique to grow LiNbO3 films [19]. Later, motivated by the discovery of high‐temperature superconductivity, OMBE was used to grow complex cuprate thin films. Up to now, it has been broadly employed to fabricate a pool of oxides, including oxide superconductors (e.g., (Ba, K)BiO3, (La, Sr)2CuO4, Bi2Sr2Ca*n*‐1Cu*n*O2*n* + 4, etc.), ferroelectrics (e.g., LiTaO3, PbTiO3, etc.), ferromagnets (e.g., (La, Ca)MnO3, EuO, etc.), multi‐ ferroics (e.g., BiFeO3, YMnO3, etc.) and superlattices of these phases [11, 12].

While conventional MBE growth occurs in an ultra‐high vacuum, in OMBE growth the induction of active gas oxidants can pose new challenges in the instrumentation as well as the film growth [11, 17]. The presence of oxidant species requires the hardware to be necessarily compatible with an oxidizing environment, thus high‐temperature components (e.g., heater filaments, effusion cells and substrate holders, etc.) need to be made of highly oxidant‐resistive materials. Moreover, adequate pumping is needed to deal with the oxidant gas load. Further‐ more, oxygen acts as another variable which needs to be optimized in the growth, and oxygen inside films is tricky to study and manipulate. In addition, the oxidants can oxidize the cell materials such that one cannot get well‐controlled fluxes as planned during growth. These challenges make the use of OMBE in the growth of oxides less mature than the use of MBE in semiconductor growth [11].

**Figure 1** shows the schematic of a typical OMBE system. Single‐element evaporators are used to generate atomic beams for OMBE growth. Knudsen cells and crucibles are chosen for elements with desired fluxes below 2000°C, while electron beam evaporators are adopted for refractory elements (e.g., tungsten, ruthenium and iridium) which require higher temperature to provide the fluxes necessary for the growth. The atomic beams impinge upon the substrate unless they are blocked by shutters which are positioned at the output end of each cell and remotely controlled by a computer. The utilization of shutters enables the elemental fluxes to be supplied to in a continuous or a sequential way. The fluxes can be adjusted by changing the cell temperature, and are *in situ* measured by a quartz crystal microbalance (QCM). Reflective high‐energy electron diffraction (RHEED) is used in OMBE for the *in situ* characterization of the growing surface. Due to the grazing angle diffraction, it can provide surface‐sensitive information including thin film crystallinity, roughness, in‐plane lattice constants, growth mechanism and phase purity. If intermediate products or impurity phases are formed, the growth conditions would be adjusted accordingly. Distilled ozone was used as the oxidant. Compared to oxygen, ozone has stronger oxidizing ability and thus needs lower pressure.

In order to access these thin film‐based physics, the first and most important step is to grow these oxide thin film structures with high quality. This needs exquisite control of growth, and usually is very challenging. In the past several decades, fortunately, reactive oxide molecular beam epitaxy (OMBE) has been proved to be an effective technique in the growth of some oxides with high quality, though being not easy [10–12]. Recently, the *in situ* combined system of OMBE and angle‐resolved photoemission spectroscopy (ARPES) [13–17] have shown great potential in exploring intricate many‐body physics based on oxide film structures, which

In this chapter, we first present basics of OMBE technique. Then, we show how to grow high‐ quality films with good stoichiometry, and the power of the integrated OMBE‐ARPES in

MBE is a vacuum deposition technique in which well‐defined thermal beams of atoms or molecules react at a crystalline surface to produce an epitaxial film. Originally, it was developed to fabricate GaAs and (Al, Ga)As films [18], and soon successfully expanded to other semi‐ conductors as well as metals and insulators. In addition to molecular beams coming from individual heated element source, gas molecular may also be introduced into MBE. Including gas oxidants (e.g., oxygen or ozone) can make an OMBE, which is now applied to grow oxides

In 1985, Betts and Pitt began to use this technique to grow LiNbO3 films [19]. Later, motivated by the discovery of high‐temperature superconductivity, OMBE was used to grow complex cuprate thin films. Up to now, it has been broadly employed to fabricate a pool of oxides, including oxide superconductors (e.g., (Ba, K)BiO3, (La, Sr)2CuO4, Bi2Sr2Ca*n*‐1Cu*n*O2*n* + 4, etc.), ferroelectrics (e.g., LiTaO3, PbTiO3, etc.), ferromagnets (e.g., (La, Ca)MnO3, EuO, etc.), multi‐

While conventional MBE growth occurs in an ultra‐high vacuum, in OMBE growth the induction of active gas oxidants can pose new challenges in the instrumentation as well as the film growth [11, 17]. The presence of oxidant species requires the hardware to be necessarily compatible with an oxidizing environment, thus high‐temperature components (e.g., heater filaments, effusion cells and substrate holders, etc.) need to be made of highly oxidant‐resistive materials. Moreover, adequate pumping is needed to deal with the oxidant gas load. Further‐ more, oxygen acts as another variable which needs to be optimized in the growth, and oxygen inside films is tricky to study and manipulate. In addition, the oxidants can oxidize the cell materials such that one cannot get well‐controlled fluxes as planned during growth. These challenges make the use of OMBE in the growth of oxides less mature than the use of MBE in

**Figure 1** shows the schematic of a typical OMBE system. Single‐element evaporators are used to generate atomic beams for OMBE growth. Knudsen cells and crucibles are chosen for

ferroics (e.g., BiFeO3, YMnO3, etc.) and superlattices of these phases [11, 12].

studying and designing many‐body interactions in complex oxides.

further intensifies the power of OMBE.

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

**2. Basics of OMBE technique**

semiconductor growth [11].

[11, 12].

**Figure 1.** Schematic illustration of an OMBE system with an ozone‐generating system. Reproduced with permission from Ref. [17].

**Figure 2** displays the photo of DCA R450 OMBE system in Shanghai Institute of Microsystem and Information Technology (SIMIT). It is equipped with 10 changeable effusion cells, and one four‐seat e‐beam evaporator for at most four refractory elements, which can cover all transition metals of interests. It also has *in situ* QCM to measure the fluxes and real‐time RHEED to directly monitor the growth. Ozone was obtained from self‐made ozone‐generating‐and‐ distilling system. Ozone generator would generate a small amount of ozone out of oxygen gas, then silica gel cooled down with liquid nitrogen would absorb ozone. Warming up the silica gel would give out the ozone gas to be used in OMBE growth.

**Figure 2.** OMBE system in Shanghai Institute of Microsystem and Information Technology (SIMIT), Chinese Academy of Sciences.

### **3. Fabricating oxide thin films with good stoichiometry by OMBE**

In this section, we show generally how to grow high‐quality oxide thin film with good stoichiometry by OMBE. Starting with choosing the proper substrate, we mainly talk about two growth methods commonly used—absorption‐controlled growth and shutter growth to achieve high‐quality films.

Epitaxial thin films cannot be obtained without using proper substrates. The substrate not only guides the growth of thin film with the right crystalline structure but also provides the knob of strain which can essentially tune the electronic structure of the material [20]. **Figure 3** displays lattice constants of single‐crystalline perovskite substrates which are commercially available and commonly used to grow perovskite or layered perovskite oxide thin films. If growing LaNiO3 films with a pseudo‐cubic lattice constant of 3.84 Å on LaAlO3 (3.75 Å) substrates, an in‐plane compressive strain was applied; if using SrTiO3 (3.905 Å) as the substrate, an in‐plane tensile strain was applied. In addition to the strain, in some cases, the choice of the substrate is vital to obtain high‐quality films. For example, Proffit et al. reported that (1 1 0) orthorhombic CaRuO3 films grown on orthorhombic (1 1 0) NdGaO3 substrates (symmetry matched) exhibit atomically smooth surfaces, whereas films on cubic lanthanum aluminate‐strontium aluminium tantalate (LSAT) substrates (symmetry mismatched) show rather rough surfaces [21]. Another example is the film growth of LaNiO3 with polar orienta‐ tions. As polar discontinuity is suggested to induce surface reconstructions which further lead to bad quality of films [22, 23], metallic Nb‐doped SrTiO3 and iso‐polarity LaAlO3 substrates were shown to be more suitable than the common SrTiO3 in the growth of LaNiO3 films [24].

**Figure 3.** A number line showing the *a*‐axis lattice constants in angstroms of the perovskite and perovskite‐related sub‐ strates that are commercially available.

As illustrated above, the remotely controlled shutters in OMBE allow elemental fluxes to be supplied to in a continuous or a sequential way. Take the growth of perovskite ABO3 (can be viewed as alternate stacking of AO and BO2 layers along the (0 0 1) direction) as an example. As schematically shown in **Figure 4**, both shutters of A cell and B cell keeping open in the whole growth make a co‐deposition growth. If the shutters of A cell and B cell alternately turn open (finishing the growth of one AO layer, and then starting the growth of one BO2 layer), we can call this the shutter growth. In either growth, stoichiometry is the most important goal which needs to be achieved.

**Figure 4.** Schematics of co‐deposition and shutter growth in OMBE growth.

**Figure 2.** OMBE system in Shanghai Institute of Microsystem and Information Technology (SIMIT), Chinese Academy

In this section, we show generally how to grow high‐quality oxide thin film with good stoichiometry by OMBE. Starting with choosing the proper substrate, we mainly talk about two growth methods commonly used—absorption‐controlled growth and shutter growth—

Epitaxial thin films cannot be obtained without using proper substrates. The substrate not only guides the growth of thin film with the right crystalline structure but also provides the knob of strain which can essentially tune the electronic structure of the material [20]. **Figure 3** displays lattice constants of single‐crystalline perovskite substrates which are commercially available and commonly used to grow perovskite or layered perovskite oxide thin films. If growing LaNiO3 films with a pseudo‐cubic lattice constant of 3.84 Å on LaAlO3 (3.75 Å) substrates, an in‐plane compressive strain was applied; if using SrTiO3 (3.905 Å) as the substrate, an in‐plane tensile strain was applied. In addition to the strain, in some cases, the choice of the substrate is vital to obtain high‐quality films. For example, Proffit et al. reported that (1 1 0) orthorhombic CaRuO3 films grown on orthorhombic (1 1 0) NdGaO3 substrates (symmetry matched) exhibit atomically smooth surfaces, whereas films on cubic lanthanum aluminate‐strontium aluminium tantalate (LSAT) substrates (symmetry mismatched) show rather rough surfaces [21]. Another example is the film growth of LaNiO3 with polar orienta‐ tions. As polar discontinuity is suggested to induce surface reconstructions which further lead

**3. Fabricating oxide thin films with good stoichiometry by OMBE**

of Sciences.

to achieve high‐quality films.

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

A well‐known growth method to achieve stoichiometric films is the adsorption‐controlled growth, which was previously used to grow GaAs and recently has been used to fabricate several oxides such as PbTiO3 [25], BiFeO3 [26] and BiMnO3 [27]. It is used in the growth of some compounds containing volatile species which can re‐evaporate during growth while the others are less volatile. If always making this volatile species excess during the growth as well as optimizing the substrate temperature and oxygen partial pressure, stoichiometric growth can be conveniently achieved. Lee et al. reported the adsorption‐controlled growth of BiM‐ nO3 in which the bismuth oxides are volatile [27]. **Figure 5** displays the calculated Ellingham diagram and obtained RHEED patterns [27]. The Bi:Mn flux ratio was fixed to be 3:1. Besides, the substrate temperature and oxygen partial pressure were fully explored to finally expose the growth window (see shadow region II in **Figure 5**) for phase‐pure stoichiometric BiMnO3 films which were verified by the shiny diffraction spots in the RHEED pattern.

**Figure 5.** Calculated Ellingham diagram and RHEED patterns collected along the [1 1 0] azimuth of SrTiO3 during the deposition of BiMnO3 at different temperatures and oxygen partial pressures. Phase stability between Bi*x*O*y* gases and BiMnO3 + Mn3O4, BiMnO3 and BiMnO3 + Bi2O2.5 condensed phases is represented by regions I, II and III, respectively. Reproduced with permission from Ref. [27].

For most oxides, adsorption‐controlled growth unfortunately cannot be applied. To achieve good stoichiometry, generally, one has to adjust the deposition amount based on cycles of combined studies of QCM, RHEED patterns and oscillations, X‐ray diffraction (XRD) pattern fitting, Rutherford backscattering spectroscopy, and so on. In the homo‐epitaxial growth of SrTiO3, Schlom's group reported the empirical method of optimizing shuttered RHEED oscil‐ lations to successfully achieve stoichiometric SrTiO3 film within 1% composition deviation [28, 29]. In the shutter growth of SrTiO3, the intensity of diffraction spot would exhibit periodic oscillations at the pace of mechanically closing/opening shutters: in Ti doses, the intensity will decrease monotonically while in Sr doses the intensity will increase. It was shown that if oscil‐ lations exhibited smooth sinoidal shape with similar amplitude (**Figure 6(a)**), the resulted film was investigated to be stoichiometric [29]. If Sr is 10% excess, the combined feature of cusp and shoulder would show up (**Figure 6(b)**); if Sr is 10% deficient, the amplitude of individual oscillations would oscillate (**Figure 6(c)**). Thus, in growth, the real‐time performance of RHEED oscillations would infer what to do next to achieve the stoichiometry [29].

A well‐known growth method to achieve stoichiometric films is the adsorption‐controlled growth, which was previously used to grow GaAs and recently has been used to fabricate several oxides such as PbTiO3 [25], BiFeO3 [26] and BiMnO3 [27]. It is used in the growth of some compounds containing volatile species which can re‐evaporate during growth while the others are less volatile. If always making this volatile species excess during the growth as well as optimizing the substrate temperature and oxygen partial pressure, stoichiometric growth can be conveniently achieved. Lee et al. reported the adsorption‐controlled growth of BiM‐ nO3 in which the bismuth oxides are volatile [27]. **Figure 5** displays the calculated Ellingham diagram and obtained RHEED patterns [27]. The Bi:Mn flux ratio was fixed to be 3:1. Besides, the substrate temperature and oxygen partial pressure were fully explored to finally expose the growth window (see shadow region II in **Figure 5**) for phase‐pure stoichiometric BiMnO3

**Figure 5.** Calculated Ellingham diagram and RHEED patterns collected along the [1 1 0] azimuth of SrTiO3 during the deposition of BiMnO3 at different temperatures and oxygen partial pressures. Phase stability between Bi*x*O*y* gases and BiMnO3 + Mn3O4, BiMnO3 and BiMnO3 + Bi2O2.5 condensed phases is represented by regions I, II and III, respectively.

For most oxides, adsorption‐controlled growth unfortunately cannot be applied. To achieve good stoichiometry, generally, one has to adjust the deposition amount based on cycles of combined studies of QCM, RHEED patterns and oscillations, X‐ray diffraction (XRD) pattern fitting, Rutherford backscattering spectroscopy, and so on. In the homo‐epitaxial growth of SrTiO3, Schlom's group reported the empirical method of optimizing shuttered RHEED oscil‐ lations to successfully achieve stoichiometric SrTiO3 film within 1% composition deviation

Reproduced with permission from Ref. [27].

films which were verified by the shiny diffraction spots in the RHEED pattern.

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

**Figure 6.** Shuttered RHEED oscillation behaviour of Sr1 + *x*TiO3 + *<sup>δ</sup>* films with stoichiometric (*x* = 0) (a), 10% strontium excess (*x* = 0.1) (b) and 10% strontium deficiency (*x* = ‐0.1) (c). Reproduced with permission from Ref. [29].

Compared to adsorption‐controlled growth, shutter growth is a more straightforward way to control the film thickness and grow complex oxide structures such as Ruddlesden‐Popper (RP) series A*n*+1B*n*O3*<sup>n</sup>*+1 and various superlattices which display a wide range of physics. By conven‐ iently changing the shuttering sequence of A and B ions to match the layering sequence of the desired RP phase, Haeni et al. and Tian et al. reported the OMBE growth of five RP members of Sr*n*+1Ti*n*O3*<sup>n</sup>*+1 [30] and Sr*n*+1Ru*n*O3*<sup>n</sup>*+1 [31], respectively. Their structures with right‐layering sequences were verified by high‐resolution cross‐sectional TEM measurements (**Figure 7**).

**Figure 7.** High‐resolution cross‐sectional TEM images of the same five members of the (a) Sr*<sup>n</sup>* + 1Ti*n*O3*<sup>n</sup>* + 1 and (b) Sr*<sup>n</sup>* + 1Ru*n*O3*n* + 1 Ruddlesden‐Popper homologous series grown by OMBE. Reproduced with permission from Refs. [30] and [31].

#### **4. The unique** *in situ* **combo of OMBE and ARPES**

ARPES can directly visualize electronic band structures of solids, and therefore has emerged as an essential experimental technique to study various novel quantum materials such as superconductors and topological quantum materials [13, 14, 32]. It can be viewed as the 'k‐ space' microscope, and can provide the essential information about how electrons move inside the material. Based on the well‐known photoelectric effect, an electron inside the solid can absorb an incident photon with a high enough energy and then emit out of the solid. If the kinetic and momentum of the photoelectrons are detected, the band structure of the material (as a function of binding energy and momentum) can be reconstructed in the context of conversation laws and some reasonable assumptions, as shown in **Figure 8**. Generally being a surface‐sensitive probe, ARPES demands clean and well‐ordered sample surface, which is usually obtained by cleaving single crystals.

of Sr*n*+1Ti*n*O3*<sup>n</sup>*+1 [30] and Sr*n*+1Ru*n*O3*<sup>n</sup>*+1 [31], respectively. Their structures with right‐layering sequences were verified by high‐resolution cross‐sectional TEM measurements (**Figure 7**).

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

**Figure 7.** High‐resolution cross‐sectional TEM images of the same five members of the (a) Sr*<sup>n</sup>* + 1Ti*n*O3*<sup>n</sup>* + 1 and (b) Sr*<sup>n</sup>* + 1Ru*n*O3*n* + 1 Ruddlesden‐Popper homologous series grown by OMBE. Reproduced with permission from Refs. [30] and

ARPES can directly visualize electronic band structures of solids, and therefore has emerged as an essential experimental technique to study various novel quantum materials such as superconductors and topological quantum materials [13, 14, 32]. It can be viewed as the 'k‐ space' microscope, and can provide the essential information about how electrons move inside the material. Based on the well‐known photoelectric effect, an electron inside the solid can absorb an incident photon with a high enough energy and then emit out of the solid. If the kinetic and momentum of the photoelectrons are detected, the band structure of the material (as a function of binding energy and momentum) can be reconstructed in the context of

**4. The unique** *in situ* **combo of OMBE and ARPES**

[31].

**Figure 8.** Schematic of ARPES measurement and the obtained band structure. Photoelectrons are emitted when shining the sample with light whose energy is larger than the sample work function. By using an electron analyser to pick up these electrons, we could obtain their energy and momentum information.

Recently, there has been increasing awareness that the *in situ* combo of OMBE and ARPES (ARPES is *in situ* connected to OMBE) could be a powerful approach in tailoring many‐body interactions and uncovering the underlying physics of complex oxides. It is a win‐win case: OMBE can provide high‐quality thin films (especially those who cannot be cleaved properly in bulk form, e.g., perovskite oxides) and interesting superlattices, whose clean surfaces 'naturally' allow for the *in situ* ARPES studies; in return, ARPES can characterize the quality of these films which further give feedbacks to the growth, and moreover it can fully explore the band structures of films and study their intriguing physics [15–17, 33, 34].

**Figure 9** shows the photo of such an *in situ* combo system located in Shanghai Institute of Microsystem and Information Technology (SIMIT), Chinese Academy of Sciences. A transfer chamber with high vacuum (∼10 × 10‐10 Torr) is used to bridge the OMBE system and ARPES

**Figure 9.** A photo of the *in situ* combo of OMBE and ARPES system in Shanghai Institute of Microsystem and Informa‐ tion Technology (SIMIT), Chinese Academy of Sciences.

system. Once the growth was finished, the film was immediately transferred to the ARPES chamber within 5 min through the transfer chamber. In so doing, the clean sample surface is expected to be preserved for ARPES studies.

Below, we first present studies on ultrathin perovskite LaNiO3 films (*3d* system) [16], SrRuO3 films (*4d* system) [35] and SrIrO3 films (*5d* system) [36] of which the bulk form lacks natural cleaving planes, as examples to show the power of this *in situ* combo in studying the many‐ body interactions of complex oxides. Generally, when going from *3d* to *5d* elements, Coulomb interactions among electrons would increase as electrons are more and more delocalized, while the spin‐orbital coupling (SOC) would become stronger due to the heavier elements. Then, we show the studies on (LaMnO3)2n/(SrMnO3)n superlattices [17].

**Figure 10.** Metal‐insulator transition in ultra‐thin LaNiO3 films. (a) The evolution of the electronic structure across the metal‐insulator transition with decreasing the LaNiO3 film thickness. (b) Schematic of 8‐UC‐thick LaNiO3 film on LaA‐ lO3 substrate. (c) Angle‐integrated photoemission spectroscopy along (0.5π/a, 0.7π/a) of films varying from 8‐UC‐thick to 1‐UC‐thick. The inset shows the near‐*E*F angle‐integrated spectra. (d) X‐ray diffraction and atomic force microscopy results. (e) Schematic of 1‐UC‐thick LaNiO3 film on LaAlO3. Reproduced with permission from Ref. [16].

Bulk LaNiO3, though being strongly correlated with *3d7* configuration of Ni3+, remains a paramagnetic metal at low temperatures. By means of OMBE, King et al. synthesized atomi‐ cally defined layers of LaNiO3 films down to just one unit cell (UC) thickness, and observed an abrupt metal‐insulator transition in 2‐UC‐thick film by transport studies [16]. The high quality of the films with atomically flat surfaces was revealed by the Kiessig fringes in XRD patterns and atomic force microscope image (**Figure 10(d)**). Then, *in situ* ARPES studies were carried out to investigate the competing electronic phases while crossing the transition, as shown in **Figure 10(a)** and **(c)**. The bulk‐like electronic structure and Fermi liquid character‐ istics were found to remain almost unaffected by film thickness down to 3 UC (all exhibiting similar electron pocket), which is in contrast to the previous reports that an insulating state emerged in 5‐UC thickness or above and again reflects the high quality of the films. This makes 3‐UC‐thick LNO the thinnest metallic nickelate reported. Reducing the thickness by just one further UC (2 UC), however, causes the spectral weight near *E*<sup>F</sup> to be suddenly suppressed. For 1‐UC‐thick LaNiO3 film, no spectral weight was observed at *E*F, indicative of a full‐charge gap. This evolution was also observed in angle‐integrated photoemission spectra. The authors claimed that the metal‐insulator transition is driven by instability to an incipient order of the underlying quantum many‐body interactions, and demonstrated the power of artificial confinement to harness control over competing phases in complex oxides with atomic‐scale precision [16].

system. Once the growth was finished, the film was immediately transferred to the ARPES chamber within 5 min through the transfer chamber. In so doing, the clean sample surface is

Below, we first present studies on ultrathin perovskite LaNiO3 films (*3d* system) [16], SrRuO3 films (*4d* system) [35] and SrIrO3 films (*5d* system) [36] of which the bulk form lacks natural cleaving planes, as examples to show the power of this *in situ* combo in studying the many‐ body interactions of complex oxides. Generally, when going from *3d* to *5d* elements, Coulomb interactions among electrons would increase as electrons are more and more delocalized, while the spin‐orbital coupling (SOC) would become stronger due to the heavier elements. Then, we

**Figure 10.** Metal‐insulator transition in ultra‐thin LaNiO3 films. (a) The evolution of the electronic structure across the metal‐insulator transition with decreasing the LaNiO3 film thickness. (b) Schematic of 8‐UC‐thick LaNiO3 film on LaA‐ lO3 substrate. (c) Angle‐integrated photoemission spectroscopy along (0.5π/a, 0.7π/a) of films varying from 8‐UC‐thick to 1‐UC‐thick. The inset shows the near‐*E*F angle‐integrated spectra. (d) X‐ray diffraction and atomic force microscopy

Bulk LaNiO3, though being strongly correlated with *3d7* configuration of Ni3+, remains a paramagnetic metal at low temperatures. By means of OMBE, King et al. synthesized atomi‐ cally defined layers of LaNiO3 films down to just one unit cell (UC) thickness, and observed an abrupt metal‐insulator transition in 2‐UC‐thick film by transport studies [16]. The high quality of the films with atomically flat surfaces was revealed by the Kiessig fringes in XRD patterns and atomic force microscope image (**Figure 10(d)**). Then, *in situ* ARPES studies were carried out to investigate the competing electronic phases while crossing the transition, as shown in **Figure 10(a)** and **(c)**. The bulk‐like electronic structure and Fermi liquid character‐ istics were found to remain almost unaffected by film thickness down to 3 UC (all exhibiting similar electron pocket), which is in contrast to the previous reports that an insulating state

results. (e) Schematic of 1‐UC‐thick LaNiO3 film on LaAlO3. Reproduced with permission from Ref. [16].

expected to be preserved for ARPES studies.

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

show the studies on (LaMnO3)2n/(SrMnO3)n superlattices [17].

**Figure 11.** Electron‐phonon coupling origin of the kink in the dispersion of SrRuO3. (a) The evolution of the electronic structure of SrRuO3 film with reducing the thickness. The kink persists in all films down to 4‐UC thick. (b) and (c) Real part and imaginary part analyses of self‐energies. All kinks occur at around 62 meV below *E*F. (d) Electron‐phonon coupling origin of the kink in SrRuO3 films. Left axis: measured negligible thickness dependence of the kink energy. Right axis: reported thickness dependence of the Curie temperature. The black rectangular marker displays the energy scale of in‐phase‐stretching phonon mode proposed by the combined studies of Raman spectroscopy and calculation. Reproduced with permission from Ref. [35].

Perovskite SrRuO3, a prototypical conductive ferromagnetic oxide, exhibits a kink in its band dispersion signalling the unusual electron dynamics therein [34]. The kink could originate from electron‐magnon coupling or electron‐phonon coupling. Uncovering the origin of this kink would hint on the studies of kinks' origins in many other intriguing systems [37–39] including the cuprate superconductor family [40]. Yang et al. reported the systematic thickness‐ dependent electronic structure studies on SrRuO3 films with well‐controlled thicknesses by using the OMBE and ARPES system [35]. **Figure 11(a)** shows the evolution of band dispersions of SrRuO3 films with reducing the film thickness. Evidently, in all these spectra, the slope of the dispersion near *E*F is markedly smaller than that of high‐binding‐energy region, namely the kink persists even down to 4‐UC‐thick film. Empirical self energy analyses were carried out to determine the kink energy. As shown in **Figure 11(b)** and **(c)**, both real part and imaginary part analyses reveal that all films' kinks are around 62 meV below *E*F. This is in sharp contrast to the report that reducing the thickness would decrease the Curie temperature [41], which implies that electron‐magnon should not play a dominant role. On the other hand, the kink energy matches that of in‐phase‐stretching phonon mode proposed by the combined studies of Raman spectroscopy and calculation [42]. Thus, electron‐phonon coupling should mainly contribute to this kink. This work can serve as an example to study the low‐energy excitations of complex oxides.

Perovskite SrIrO3, due to the heavy element of Ir, is expected to have strong SOC which is the key ingredient in building topological quantum materials [43, 44]. Therefore, novel topological phases were proposed in artificial SrIrO3‐based structures [45]. In particular, SrIrO3 was proposed to be an exotic semimetal induced by the delicate interplay between SOC and electron correlations, in which a Dirac nodal ring near the *U* point would render a non‐trivial topological semi‐metallic state [46, 47]. Based on the OMBE‐ARPES combo, Liu et al. synthe‐ sized high‐quality SrIrO3 films, and investigated its low‐lying electronic structure [36]. **Figure 12** displays the measured band structure of SrIrO3. In addition to the semimetal state (the Fermi level simultaneously crosses the hole and electron pockets), near the *U* point, a lifted Dirac node was directly observed, which agrees well with the authors' calculations. This Dirac node lifting could be due to the selectively breaking of *n*‐glide symmetry in the hetero‐epitaxial SrIrO3 structure [36]. Iridates would continue acting as the profound platform to explore novel physics that may combine the SOC and electron correlations.

**Figure 12.** Dirac line node degeneracy lifting around the *U* point of SrIO3. (a) and (b) The second derivative images along *Z‐U* and *U‐R* high‐symmetry directions, respectively. (c) and (d) The corresponding energy‐distribution curves for the photoemission data in (a, b). (e) and (f) The calculated band dispersions along *Z‐U* and *U‐R* high‐symmetry directions, respectively. Reproduced with permission from Ref. [36].

With the powerful capability of the OMBE‐ARPES combo, one can fabricate artificial superlat‐ tices which do not exist in nature and study their intriguing emergent physics. Monkman et al. reported the comprehensive investigations on the interfacial electronic structure of (LaM‐ nO3)2*n*/(SrMnO3)*<sup>n</sup>* superlattices as a function of dimensionality [17]. Bulk LaMnO3 and SrMnO3 are anti‐ferromagnetic Mott and band insulators, respectively, and La2/3Sr1/3MnO3 is a ferro‐ magnetic metal that exhibits colossal magnetoresistance around its Curie temperature of 370 K. Monkman et al. synthesized high‐quality (LaMnO3)2*n*/(SrMnO3)*<sup>n</sup>* superlattices with *n* = 1–3. Transport studies show that *n* = 1 and *n* = 2 members show metallic behaviours at low temper‐ atures, while the *n* = 3 member exhibits a metal‐insulator crossover. **Figure 13** illustrates the evolution of the electronic structure and properties of superlattices upon different *n*. For *n* = 1 and 2 members, the Fermi surfaces are apparent and consist of two Mn *eg*‐derived states: a hole pocket around the Brillouin zone corner, and a smaller electron pocket around the zone centre. For the insulating *n* = 3 member, the spectral weight at *E*F is highly suppressed, although clear states are still observed below *E*F. The authors also investigated the near‐*E*<sup>F</sup> band dispersions. The *n* = 1 and 2 members exhibit well‐defined and dispersive bands, whereas the *n* = 3 sample shows only pseudo‐gapped intensity at *E*<sup>F</sup> which is similar to polaronic systems with strong electron‐phonon coupling. For the critical *n* = 2, there exists a dramatic difference in electronic states compared to that of *n* = 1. The states near *E*<sup>F</sup> are substantially suppressed, and a broad incoherent feature appears below *E*F, indicating the enhanced correlations. This work pro‐ vides unique insight into how many‐body interactions could be engineered at correlated ox‐ ide interfaces, which is an important prerequisite to exploiting such effects in novel electronics [17].

the kink persists even down to 4‐UC‐thick film. Empirical self energy analyses were carried out to determine the kink energy. As shown in **Figure 11(b)** and **(c)**, both real part and imaginary part analyses reveal that all films' kinks are around 62 meV below *E*F. This is in sharp contrast to the report that reducing the thickness would decrease the Curie temperature [41], which implies that electron‐magnon should not play a dominant role. On the other hand, the kink energy matches that of in‐phase‐stretching phonon mode proposed by the combined studies of Raman spectroscopy and calculation [42]. Thus, electron‐phonon coupling should mainly contribute to this kink. This work can serve as an example to study the low‐energy

Perovskite SrIrO3, due to the heavy element of Ir, is expected to have strong SOC which is the key ingredient in building topological quantum materials [43, 44]. Therefore, novel topological phases were proposed in artificial SrIrO3‐based structures [45]. In particular, SrIrO3 was proposed to be an exotic semimetal induced by the delicate interplay between SOC and electron correlations, in which a Dirac nodal ring near the *U* point would render a non‐trivial topological semi‐metallic state [46, 47]. Based on the OMBE‐ARPES combo, Liu et al. synthe‐ sized high‐quality SrIrO3 films, and investigated its low‐lying electronic structure [36]. **Figure 12** displays the measured band structure of SrIrO3. In addition to the semimetal state (the Fermi level simultaneously crosses the hole and electron pockets), near the *U* point, a lifted Dirac node was directly observed, which agrees well with the authors' calculations. This Dirac node lifting could be due to the selectively breaking of *n*‐glide symmetry in the hetero‐epitaxial SrIrO3 structure [36]. Iridates would continue acting as the profound platform to explore novel

**Figure 12.** Dirac line node degeneracy lifting around the *U* point of SrIO3. (a) and (b) The second derivative images along *Z‐U* and *U‐R* high‐symmetry directions, respectively. (c) and (d) The corresponding energy‐distribution curves for the photoemission data in (a, b). (e) and (f) The calculated band dispersions along *Z‐U* and *U‐R* high‐symmetry

physics that may combine the SOC and electron correlations.

directions, respectively. Reproduced with permission from Ref. [36].

excitations of complex oxides.

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

**Figure 13.** (a)–(c) High‐angle annular dark‐field scanning transmission electron micrograph (i), Fermi surface (ii) and band dispersion (iii) of (LaMnO3)2*n*/(SrMnO3)*<sup>n</sup>* superlattices with *n* = 1–3, respectively. Reproduced with permission from Ref. [17].

These examples reflect the powerful capability of the integrated OMBE‐ARPES system in studying many‐body interactions and resulted novel physics in complex oxides.

### **5. Conclusion and outlook**

In this chapter, we presented the brief inductions to OMBE technique, growth methods and the *in situ* combo integrating oxide MBE and APRES. We demonstrate that OMBE and the *in situ* combo formed with ARPES will continue playing a role in unveiling the intriguing many‐ body physics of correlated oxide thin films and superlattices as well as exploring novel topological materials in oxide structures.

### **Author details**

Dawei Shen1,2,3\*, Haifeng Yang1 and Zhengtai Liu1

\*Address all correspondence to: dwshen@mail.sim.ac.cn

1 State Key Laboratory of Functional Materials for Informatics, Shanghai Institute of Microsystem and Information Technology (SIMIT), Chinese Academy of Sciences, Shanghai, China

2 CAS Centre for Excellence in Superconducting Electronics (CENSE), Shanghai, China

3 CAS‐Shanghai Science Research Centre, Shanghai, China

### **References**


[6] A. Ohtomo, H. Y. Hwang. A high‐mobility electron gas at the LaAlO3/SrTiO3 heteroin‐ terface. Nature. 2004; 427 (6973): 423–426.

These examples reflect the powerful capability of the integrated OMBE‐ARPES system in

In this chapter, we presented the brief inductions to OMBE technique, growth methods and the *in situ* combo integrating oxide MBE and APRES. We demonstrate that OMBE and the *in situ* combo formed with ARPES will continue playing a role in unveiling the intriguing many‐ body physics of correlated oxide thin films and superlattices as well as exploring novel

studying many‐body interactions and resulted novel physics in complex oxides.

and Zhengtai Liu1

1 State Key Laboratory of Functional Materials for Informatics, Shanghai Institute of

2 CAS Centre for Excellence in Superconducting Electronics (CENSE), Shanghai, China

Microsystem and Information Technology (SIMIT), Chinese Academy of Sciences, Shanghai,

[1] P. A. Lee, N. Nagaosa, X. ‐G. Wen. Doping a Mott insulator: Physics of high‐temperature

[2] M. Imada, A. Fujimori, Y. Tokura. Metal‐insulator transitions. Rev. Mod. Phys. 1998; 70

[3] M. B. Salamon, M. Jaime. The physics of manganites: Structure and transport. Rev. Mod.

[4] H. Y. Hwang, Y. Iwasa, M. Kawasaki, B. Keimer, N. Nagaosa, Y. Tokura. Emergent

[5] J. Chakhalian, A. J. Millis, J. Rondinelli. Whither the oxide interface. Nat. Mater. 2012;

phenomena at oxide interfaces. Nat. Mater. 2012; 11 (2): 103–113.

**5. Conclusion and outlook**

**Author details**

China

**References**

(4): 1039–1263.

11 (2): 92–94.

Phys. 2001; 73 (3): 583–628.

Dawei Shen1,2,3\*, Haifeng Yang1

topological materials in oxide structures.

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

\*Address all correspondence to: dwshen@mail.sim.ac.cn

3 CAS‐Shanghai Science Research Centre, Shanghai, China

superconductivity. Rev. Mod. Phys. 2006; 78 (1): 17–85.


[32] Y. L. Chen. Studies on the electronic structures of three‐dimensional topological insulators by angle resolved photoemission spectroscopy. Front Phys‐Beijing. 2012; 7 (2): 175–192.

[20] B. Burganov, C. Adamo, A. Mulder, M. Uchida, P. D. C. King, J. W. Harter, D. E. Shai, A. S. Gibbs, A. P. Mackenzie, R. Uecker, M. Bruetzam, M. R. Beasley, C. J. Fennie, D. G. Schlom, K. M. Shen. Strain control of Fermiology and many‐body interactions in two‐

[21] D. L. Proffit, H. W. Jang, S. Lee, C. T. Nelson, X. Q. Pan, M. S. Rzchowski, C. B. Eom. Influence of symmetry mismatch on heteroepitaxial growth of perovskite thin films.

[22] N. Nakagawa, H. Y. Hwang, D. A. Muller. Why some interfaces cannot be sharp. Nat.

[23] J. L. Blok, X. Wan, G. Koster, D. H. A. Blank. Epitaxial oxide growth on polar (111)

[24] H. F. Yang, Z. T. Liu, C. C. Fan, Q. Yao, P. Xiang, K. L. Zhang, M. Y. Li, J. S. Liu, D. W. Shen. Avoiding polar catastrophe in the growth of polarly orientated nickel perovskite thin films by reactive oxide molecular beam epitaxy. AIP Advs. 2016; 6: 085115.

[25] C. Theis, J. Yeh, D. Schlom, M. Hawley, G. Brown. Adsorption‐controlled growth of PbTiO3 by reactive molecular beam epitaxy. Thin Solid Films. 1998; 325 (1C2): 107–114.

[26] J. F. Ihlefeld, A. Kumar, V. Gopalan, D. G. Schlom, Y.B. Chen, X.Q. Pan, T. Heeg, J. Schubert, X. Ke, P. Schiffer, J. Orenstein, L. W. Martin, Y.H. Chu, R. Ramesh. Adsorption‐ controlled molecular‐beam epitaxial growth of BiFeO3. Appl. Phys. Lett. 2007; 91 (7):

[27] J. H. Lee, X. Ke, R. Misra, J. F. Ihlefeld, X. S. Xu, Z. G. Mei, T. Heeg, M. Roeckerath, J. Schubert, Z. K. Liu, J. L. Musfeldt, P. Schiffer, D. G. Schlom. Adsorption‐controlled growth of BiMnO3 by molecular‐beam epitaxial. Appl. Phys. Lett. 2010; 96 (26): 262905.

[28] J. Haeni, C. Theis, D. Schlom. RHEED intensity oscillations for the stoichiometric growth of SrTiO3 thin films by reactive molecular beam epitaxy. J. Electroceram. 2000;

[29] C. M. Brooks, L. F. Kourkoutis, T. Heeg, J. Schubert, D. A. Muller, D. G. Schlom. Growth of homoepitaxial SrTiO3 thin films by molecular‐beam epitaxy. Appl. Phys. Lett. 2009;

[30] J. H. Haeni, C. D. Theis, D. G. Schlom, W. Tian, X. Q. Pan, H. Chang, I. Takeuchi, X. ‐D. Xiang. Epitaxial growth of the first five members of the Srn+1TinO3n+1Ruddlesden‐Popper

[31] W. Tian, J. H. Haeni, D. G. Schlom, E. Hutchinson, B. L. Sheu, M. M. Rosario, P. Schiffer, Y. Liu, M. A. Zurbuchen, X. Q. Pan. Epitaxial growth and magnetic properties of the first five members of the layered Srn+1RunO3n+1 oxide series. Appl. Phys. Lett. 2007; 90

homologous series. Appl. Phys. Lett. 2001; 78 (21): 3292–3294.

dimensional Ruthenates. Phys. Rev. Lett. 2016; 116 (19): 197003.

Appl. Phys. Letts. 2008; 93 (11): 111912.

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

surfaces. Appl. Phys. Lett. 2011; 99 (15): 151917.

Mater. 2006; 5 (3): 204–209.

071922.

4 (2–3): 385–391.

94 (16): 162905.

(2): 022507.


### **Anomalous Rashba Effect of Bi Thin Film Studied by Spin-Resolved ARPES Anomalous Rashba Effect of Bi Thin Film Studied by Spin-Resolved ARPES**

Akari Takayama Akari Takayama

[44] Z. Hasan, C. L. Kane. Colloquium: Topological insulators. Rev. Mod. Phys. 2010; 82 (4):

[45] D. Xiao, W. Zhu, Y. Ran, N. Nagaosa, S. Okamoto. Interface engineering of quantum Hall effects in digital transition‐metal oxide heterostructures. Nat. Commun. 2011; 2

[46] J. ‐M. Carter, V. V. Shankar, M. A. Zeb, H. ‐Y. Kee. Semimetal and topological insulator

[47] M. A. Zeb, H. ‐Y. Kee. Interplay between spin‐orbit coupling and Hubbard interactions in SrIrO3 and related Pbnm perovskite oxides. Phys. Rev. B. 2012; 86 (86): 085149.

in perovskite iridates. Phys. Rev. B. 2012; 85 (11): 115105.

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

3045–3067.

(6): 596.

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

http://dx.doi.org/10.5772/66278

#### **Abstract**

The Rashba effect is a momentum‐dependent splitting of spin bands in two‐dimen‐ sional systems such as surface, interface and heterostructure. The effect is caused by broken space‐inversion symmetry and spin‐orbit coupling and allows to manipulate and generate the spin by the electric fields, that is, without the magnetic field. It means that the devices applied to the Rashba effect have many advantages. Bismuth is known as a promising candidate to investigate the surface Rashba effect, and the spin struc‐ ture of Bi surface has also been intensively discussed. However, it is unclear to what extent the so far believed simple vortical spin structure is adequate. To understand the surface properties of the Rashba system is particularly important when utilizing the Rashba effect to the spintronic devices, since it is desirable to control the spin polariza‐ tion when developing new types of devices. In this chapter, we report that the surface spin states of the Bi thin film exhibit unusual characteristics unlike the conventional Rashba splitting by using a spin‐ and angle‐resolved photoemission spectroscopy measurement.

**Keywords:** Rashba effect, spin‐resolved ARPES, thin film, bismuth

### **1. Introduction**

As we know and use, spintronic devices to use a spin‐polarized electrons have actualized. The magnetic storage technology uses giant magneto‐resistance [1]. A more advanced approach is to control spin‐polarized electrons without the aid of a ferromagnetism nor to apply the magnetic field [2]. Spin‐orbit coupling (SOC) makes it possible to generate and manipulate spin‐polarized electrons only by the electric field, since the electric field acts on a moving charge carrier as an effective magnetic field. Thus, it is regarded as an

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

© 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, © 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.

essential ingredient for further development of next‐generation spintronic devices such as the spin‐field‐effect transistor [3]. In nonmagnetic solids, the electronic states with oppo‐ site spin have the same energy (Kramers degeneracy) because of the time‐reversal and the space‐inversion symmetries (TRS and SIS). In the strong SOC environment with the broken space‐inversion symmetry (typically at the surface or interface), the energy band splits in the momentum (*k*) space (Rashba effect [4]), leading to a spin helical structure of surface bands. This Rashba effect leads to the vortical spin structure of surface bands where the spin vector points parallel to the surface and perpendicular to the measured momentum. To be more specific, here we postulate the model of two‐dimensional free electron gas at the surface. In the surface, as shown in **Figure 1a**, there is asymmetry of the potential in the direction perpendicular to the two‐dimensional plane [∇V = (0, 0, *E*z )]. When an electron moves with momentum (*p*), the ∇V × *p* term acts as an effective magnetic field *B*eff, which is orthogonal to *p* (=*ħk*) and ∇V. As a result, the electron spin is quantized along the direction perpendicular to *k* in the surface plane. The energy of free electron gives the following,

$$E(k) = \frac{\hbar^2}{2m}k^2 \pm a\_\kappa k \tag{1}$$

which is well known as the Rashba effect, where *α*R is a so‐called Rashba parameter. **Figure 1b** shows the band dispersion of this model. As described above, two‐dimensional (2D) system with strong SOC has provided a useful platform for realizing novel quantum phenomena

**Figure 1.** (a) Relationship between momentum (*p*) and surface potential (∇V) at surface. (b) Rashba‐type spin‐splitting band structure in 2D free electron gas model. (c) A band structure breaking the time‐reversal symmetry by applying a magnetic field.

applicable to advanced spintronic devices [2, 5, 6]. To discuss and understand Rashba effect, spin‐ and angle‐resolved photoemission spectroscopy (spin‐resolved ARPES) is a powerful experimental technique, which can simultaneously determine all key quantum parameters of electrons in solids, that is, momentum, energy and spin. Furthermore, by high‐resolution measurements, we will be able to discuss the spin‐related physical phenomena not only quali‐ tatively but also quantitatively, which would certainly lead to the deeper understanding of the condensed matter physics. Elucidation of the electron spin is very important to under‐ stand physical property in solid state and its surface as well as the possible applications to spintronic devices.

essential ingredient for further development of next‐generation spintronic devices such as the spin‐field‐effect transistor [3]. In nonmagnetic solids, the electronic states with oppo‐ site spin have the same energy (Kramers degeneracy) because of the time‐reversal and the space‐inversion symmetries (TRS and SIS). In the strong SOC environment with the broken space‐inversion symmetry (typically at the surface or interface), the energy band splits in the momentum (*k*) space (Rashba effect [4]), leading to a spin helical structure of surface bands. This Rashba effect leads to the vortical spin structure of surface bands where the spin vector points parallel to the surface and perpendicular to the measured momentum. To be more specific, here we postulate the model of two‐dimensional free electron gas at the surface. In the surface, as shown in **Figure 1a**, there is asymmetry of the potential in the direction perpendicular to the two‐dimensional plane [∇V = (0, 0,

)]. When an electron moves with momentum (*p*), the ∇V × *p* term acts as an effective magnetic field *B*eff, which is orthogonal to *p* (=*ħk*) and ∇V. As a result, the electron spin is quantized along the direction perpendicular to *k* in the surface plane. The energy of free

which is well known as the Rashba effect, where *α*R is a so‐called Rashba parameter. **Figure 1b** shows the band dispersion of this model. As described above, two‐dimensional (2D) system with strong SOC has provided a useful platform for realizing novel quantum phenomena

**Figure 1.** (a) Relationship between momentum (*p*) and surface potential (∇V) at surface. (b) Rashba‐type spin‐splitting band structure in 2D free electron gas model. (c) A band structure breaking the time‐reversal symmetry by applying

<sup>2</sup>*<sup>m</sup> k*<sup>2</sup> ± *α<sup>R</sup> k* (1)

*E*z

a magnetic field.

electron gives the following,

*E*(*k*) = *<sup>ħ</sup>*<sup>2</sup> *\_\_\_*

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

The first observation of the surface Rashba effect by ARPES is the Au(1 1 1) surface [7]. After that, various materials such as group‐V semimetals and their alloy surfaces [8–16], as well as heavy‐atom adsorbed semiconductor surfaces [17–22] and so on, are studied. Among them, the group‐V semimetal bismuth (Bi) is a prime candidate to investigate the surface Rashba effect and many experiments and theoretical calculations studied in order to clarify the fundamental properties of the Rashba effect [8, 10–14]. However, in previous researches, although the band structure and Fermi surfaces of Bi distinctly show strong anisotropy, the spin structure of Bi was argued by assuming an isotropic two‐circular Fermi‐surface model like Au(1 1 1) [7, 20–22]. The reason is because the energy and momentum resolutions of the previous spin‐resolved ARPES machine are insufficient. To understand whole aspect of the Rashba effect, it is necessary to clarify that the spin structure of Bi consists with the conventional Rashba model or not. Moreover, in the film, electronic states originating from the bulk Bi are quantized, and they are connected to the surface bands continuously. So it would be necessary to take into account the relationship between bulk and surface states. On the other hand, intensive attempts have been made to extend the investigations on 2D Rashba systems to quasi one‐dimensional (1D) system like artificially grown nanowires and quantum wires, because of the merits in downsizing of devices.1D Rashba effect in utiliza‐ tion of vicinal surfaces such as in Au chains on vicinal Si [23] and the vicinal Bi surface is also reported [24]. In such a case, breaking the TRS by applying magnetic field or adding magnetic impurities would create an energy gap at the Kramers point, and when the chemi‐ cal potential is tuned to be located in the spin‐orbit gap, the dissipation‐less spin transport and the quantized conductance [25, 26] may be realized (**Figure 1c**). However, the Rashba effect in edge state is not known because the signal from the edge is extremely faint. To understand the proposed novel properties of a true 1D system, we may be able to apply them to advanced spintronic devices.

In this chapter, we introduce the electronic structure of Bi thin film to elucidate the details of the Rashba effect by utilizing the high‐resolution spin‐resolved ARPES spectrometer equipped with a highly efficient mini‐Mott detector. We show three novel Rashba effects of Bi thin film: (i) anisotropic Rashba effect from momentum‐dependent measurement [27], (ii) the interface Rashba effect between metal‐semiconductor from thickness‐dependent [28] and (iii) 1D Rashba effect of edge state [29]. The present finding provides a useful platform to study the Rashba effect and at the same time opens a pathway to utilize the novel properties to advanced spintronic devices.

### **2. Experimental technique and sample fabrication**

### **2.1. Spin‐ and angle‐resolved photoemission spectromete**

**Figure 2** shows a schematic diagram of the ultra‐high‐resolution spin‐resolved ARPES spec‐ trometer with a highly efficient mini‐Mott detector [30]. This spectrometer consists of mainly four parts: (i) a photoemission measurement system including a hemispherical electron energy analyzer and an ultra‐high‐vacuum measurement chamber, (ii) a spin‐detection sys‐ tem based on a mini‐Mott detector, (iii) an intense xenon/helium plasma discharge lamp and (iv) a surface chamber to prepare the thin‐film samples. We explain the detail of each part. We have improved a MBS‐A1 electron energy analyzer to achieve both spin‐resolved and regular (non‐spin‐resolved) ARPES measurement. The spectrometer has two detectors: one is a multichannel plate for ARPES measurement, and the other is mini‐Mott detector for spin‐ resolved ARPES. To determine the three‐dimensional spin polarization, an electron deflector has been placed between the analyzer and the Mott detector. The Mott detector observes the spin polarization of essentially two independent axes by using four channeltrons, enabling us to determine the in‐plane and out‐of‐plane spin component. The scattering efficiency of the Mott detector is as high as 2.3 × 10‐2 . The optical system consists of helium (He) and xenon (Xe) plasma discharge lamps and a monochromator with the gratings, in which we can select

**Figure 2.** Schematic view of high‐resolution spin‐resolved photoemission spectrometer.

photon energy if necessary. In this study, we used one of the Xe I lines (*h*ν = 8.437 eV) to excite photoelectrons. In order to fabricate high‐quality samples, a surface chamber has been constructed and is connected to the spin‐resolved ARPRS spectrometer. The spectrometer achieves the energy resolutions of 0.9 and 8 meV for non–spin‐resolved and spin‐resolved modes, respectively. **Figure 3** shows a schematic view of the surface chamber. A surface chamber contains heating systems for a semiconductor sample, a few kinds of dispensers, a quartz crystal microbalance and a low‐energy electron diffraction (LEED) system for checking the quality of the sample surface. The vacuum of the surface chamber is basically kept 1 × 10‐10 Torr to prepare a high‐quality sample surface. It allows the in situ preparation of the sample and its transfer to the spin‐resolved spectrometer. It is particularly useful for elucidating the electronic states of the samples containing a clean well‐ordered surfaces required for the accu‐ rate ARPES measurements. Thus, we can prepare the Bi thin‐film samples by the evaporation of Bi on semiconducting substrate. We expect that the performance of the spectrometer is demonstrated by the observation of a clear Rashba splitting of the surface states in Bi. The energy and momentum resolutions for the regular (spin‐integrated) ARPES were 5–20 meV and 0.3°, and for spin‐resolved ARPES measurements were 40 meV and 3°, respectively. The Sherman function value was set at 0.07. The measurement temperature was 300 and 30 K.

**Figure 3.** Schematic view of the sample preparation chamber.

**2. Experimental technique and sample fabrication**

**Figure 2** shows a schematic diagram of the ultra‐high‐resolution spin‐resolved ARPES spec‐ trometer with a highly efficient mini‐Mott detector [30]. This spectrometer consists of mainly four parts: (i) a photoemission measurement system including a hemispherical electron energy analyzer and an ultra‐high‐vacuum measurement chamber, (ii) a spin‐detection sys‐ tem based on a mini‐Mott detector, (iii) an intense xenon/helium plasma discharge lamp and (iv) a surface chamber to prepare the thin‐film samples. We explain the detail of each part. We have improved a MBS‐A1 electron energy analyzer to achieve both spin‐resolved and regular (non‐spin‐resolved) ARPES measurement. The spectrometer has two detectors: one is a multichannel plate for ARPES measurement, and the other is mini‐Mott detector for spin‐ resolved ARPES. To determine the three‐dimensional spin polarization, an electron deflector has been placed between the analyzer and the Mott detector. The Mott detector observes the spin polarization of essentially two independent axes by using four channeltrons, enabling us to determine the in‐plane and out‐of‐plane spin component. The scattering efficiency of

(Xe) plasma discharge lamps and a monochromator with the gratings, in which we can select

**Figure 2.** Schematic view of high‐resolution spin‐resolved photoemission spectrometer.

. The optical system consists of helium (He) and xenon

**2.1. Spin‐ and angle‐resolved photoemission spectromete**

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

the Mott detector is as high as 2.3 × 10‐2

### **2.2. Sample fabrication**

To get a high‐quality Bi(1 1 1) thin film, we prepare a clean surface of the Si(1 1 1) substrate. We use a commercially available Si wafer (n‐type, As‐doped: 0.001–0.005 Ω cm, Sb‐doped: 0.01– 0.02 Ω cm), and the surface of Si forms native SiO<sup>2</sup> in the air. So we must remove it by electrical heating in high vacuum. **Figure 4** displays the design of a holder for the electrical heating system (**Figure 4**). The holder is made of molybdenum, and a main holder is insulated from a subholder by the alumina. We can control a current and temperature with 5 mA and 1°C, respectively. The sample size is 13 × 3 mm, and we have selected four different kinds of crystal orientations on the basis of orientation flat as shown in **Figure 5**. The sample geometry has been confirmed by the brightness symmetry of the LEED spots. The most stable structure of the Si(1 1 1) surface is the 7 × 7 reconstructed surface, as shown in **Figure 6a** [31, 32]. Now, we describe the method to pre‐

**Figure 4.** (a) A main and subparts of sample holder. (b) A constructed holder.

**Figure 5.** Cutting the sample from the Si(1 1 1) wafer. Inset shows the picture of a sample holder compared to the sample size. The geometry of the sample was determined by the direction of "orientation flat."

**2.2. Sample fabrication**

0.02 Ω cm), and the surface of Si forms native SiO<sup>2</sup>

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

**Figure 4.** (a) A main and subparts of sample holder. (b) A constructed holder.

To get a high‐quality Bi(1 1 1) thin film, we prepare a clean surface of the Si(1 1 1) substrate. We use a commercially available Si wafer (n‐type, As‐doped: 0.001–0.005 Ω cm, Sb‐doped: 0.01–

heating in high vacuum. **Figure 4** displays the design of a holder for the electrical heating system (**Figure 4**). The holder is made of molybdenum, and a main holder is insulated from a subholder by the alumina. We can control a current and temperature with 5 mA and 1°C, respectively. The sample size is 13 × 3 mm, and we have selected four different kinds of crystal orientations on the basis of orientation flat as shown in **Figure 5**. The sample geometry has been confirmed by the brightness symmetry of the LEED spots. The most stable structure of the Si(1 1 1) surface is the 7 × 7 reconstructed surface, as shown in **Figure 6a** [31, 32]. Now, we describe the method to pre‐

**Figure 5.** Cutting the sample from the Si(1 1 1) wafer. Inset shows the picture of a sample holder compared to the sample

size. The geometry of the sample was determined by the direction of "orientation flat."

in the air. So we must remove it by electrical

**Figure 6.** (a) Atomic‐structure model (Das model) of Si(1 1 1)‐7 × 7. (b) Detailed annealing process to prepare the Si(1 1 1)‐7 × 7 clean surface.

**Figure 7.** Comparison of the LEED pattern. (a) Si(1 1 1)‐7 × 7. (b) and (c) Bi(1 1 1)‐1 × 1. These two samples have a different cutting directions from Si wafer [sample direction in (c) is rotated in 180° from (b)]. (d) A multiple‐domain sample of Bi(1 1 1)‐1 × 1.

pare a Si(1 1 1)‐7 × 7 reconstructed surface. First, the Si wafers were outgassed for more than 12 h below 750°C. After the outgassing enough, we have carried out flash annealing. **Figure 6b** shows a flash‐annealing process to get a Si(1 11)‐7 × 7 reconstructed surface; sample was (i) heated at 750°C to 1050–1200°C for a few seconds, (ii) keep temperature to maximum for 5 s, (iii) cooled down to 850°C for a few seconds and (iv) cooled to 750°C in 30 s [33]. We have repeated this cycle, and all of the above processes should be performed under the ultra‐high vacuum of ~1 × 10‐10 Torr. As shown in **Figure 7a**, we have obtained the LEED pattern of the well‐ordered 7 × 7 surface.

Next, Bi atoms are evaporated on Si substrate, which is called as a molecular beam epi‐ taxy (MBE). Bi atoms are deposited at room temperature on Si(1 1 1)‐7 × 7 reconstructed surface. Then, the Bi thin film was annealed at 150°C. The deposition rate is estimated by the quartz oscillator thickness monitor, and the film thickness was controlled by varying the deposition time with keeping the constant deposition rate. We can also estimate the film thickness from the energy position of the quantum well states (QWSs) in the ARPES spectra [28]. It is noted that 1 bilayer (BL) Bi is defined as 1.14 × 10<sup>15</sup> atoms/cm<sup>2</sup> , and the thickness is 0.39 nm [13]. In this study, we prepared several thickness sample (8–40 BL). After the deposition of Bi atoms, the LEED pattern shows the 1 × 1 surface structure as shown in **Figure 7b** and **c**, and the intensity of the LEED spot has threefold symmetry. When the sample has a multi‐domain structure, the LEED pattern shows circular features surrounding the 1 × 1 spots (**Figure 7d**). We could repeatedly use one Si substrate by a flash‐ ing. Here we describe a structure of Bi thin film on Si substrate. Details of the thin‐film growth process are shown in **Figure 8a**, which is reported in previous works such as STM and LEED [34]. Although the lattice constant of Bi (4.538 Å) is very different from Si (5.43 Å), it is possible to fabricate the Bi/Si thin film due to the existence of disordered layer called "wetting layer" between Bi and Si substrate. As shown in **Figure 8b**, the structural transi‐ tion from {0 1 2} direction to (1 1 1) direction more than 8.4 ML suddenly takes place upon Bi deposition.

pare a Si(1 1 1)‐7 × 7 reconstructed surface. First, the Si wafers were outgassed for more than 12 h below 750°C. After the outgassing enough, we have carried out flash annealing. **Figure 6b** shows a flash‐annealing process to get a Si(1 11)‐7 × 7 reconstructed surface; sample was (i) heated at 750°C to 1050–1200°C for a few seconds, (ii) keep temperature to maximum for 5 s, (iii) cooled down to 850°C for a few seconds and (iv) cooled to 750°C in 30 s [33]. We have repeated this cycle, and all of the above processes should be performed under the ultra‐high vacuum of ~1 × 10‐10 Torr. As shown in **Figure 7a**, we have obtained the LEED pattern of the well‐ordered 7 × 7

**Figure 7.** Comparison of the LEED pattern. (a) Si(1 1 1)‐7 × 7. (b) and (c) Bi(1 1 1)‐1 × 1. These two samples have a different cutting directions from Si wafer [sample direction in (c) is rotated in 180° from (b)]. (d) A multiple‐domain sample of

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

Next, Bi atoms are evaporated on Si substrate, which is called as a molecular beam epi‐ taxy (MBE). Bi atoms are deposited at room temperature on Si(1 1 1)‐7 × 7 reconstructed surface. Then, the Bi thin film was annealed at 150°C. The deposition rate is estimated by the quartz oscillator thickness monitor, and the film thickness was controlled by varying the deposition time with keeping the constant deposition rate. We can also estimate the film thickness from the energy position of the quantum well states (QWSs) in the ARPES

thickness is 0.39 nm [13]. In this study, we prepared several thickness sample (8–40 BL). After the deposition of Bi atoms, the LEED pattern shows the 1 × 1 surface structure as shown in **Figure 7b** and **c**, and the intensity of the LEED spot has threefold symmetry. When the sample has a multi‐domain structure, the LEED pattern shows circular features surrounding the 1 × 1 spots (**Figure 7d**). We could repeatedly use one Si substrate by a flash‐ ing. Here we describe a structure of Bi thin film on Si substrate. Details of the thin‐film growth process are shown in **Figure 8a**, which is reported in previous works such as STM and LEED [34]. Although the lattice constant of Bi (4.538 Å) is very different from Si (5.43 Å), it is possible to fabricate the Bi/Si thin film due to the existence of disordered layer called "wetting layer" between Bi and Si substrate. As shown in **Figure 8b**, the structural transi‐ tion from {0 1 2} direction to (1 1 1) direction more than 8.4 ML suddenly takes place upon

, and the

spectra [28]. It is noted that 1 bilayer (BL) Bi is defined as 1.14 × 10<sup>15</sup> atoms/cm<sup>2</sup>

surface.

Bi(1 1 1)‐1 × 1.

Bi deposition.

**Figure 8.** (a) Thickness dependence of spot‐profile‐analyzing‐LEED pattern (top) and STM image (bottom) [34]. (b) Schematic illustration of the transformed structure in growth process of Bi thin film [34].

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

#### **3.1. Anisotropic Rashba effect of Bi thin film**

At first, we have performed normal (non–spin‐resolved) ARPES measurement of the Bi thin film in order to check the sample quality and geometry, since even a subtle misalignment of the sample orientation would cause a significant error in determining the spin polarization. **Figure 9a** and **b** shows the band dispersion along the Γ¯ ¯ M line and Fermi surface of Bi/Si(1 1 1). Areas with hatched lines in **Figure 9a** are bulk band projection. We distinctly see that several bands cross *E*<sup>F</sup> and three kinds Fermi surfaces exist; a hexagonal Fermi surface centered at the Γ¯ point (S<sup>1</sup> ), surrounding elongated pockets (S<sup>2</sup> ) and the ellipsoidal pocket (S<sup>3</sup> ) near the ¯ M point. Judged from the band dispersion in **Figure 9a**, the S<sup>1</sup> and S<sup>3</sup> are attributed to the electron pock‐ ets, while the S<sup>2</sup> to a hole pocket, all of which arise from the spin‐split surface states [8, 10–14].

This section focuses on the spin structure of the S<sup>2</sup> band. In a conventional 2D Rashba model, the in‐plane spin has a vortical structure and isotropic magnitude as denoted by black arrows in **Figure 10a**. In **Figure 10b** and **e**, we display the near‐*E*<sup>F</sup> spin‐resolved energy distribution curves (EDCs) for the in‐plane (*y*/*x*) and out‐of‐plane (*z*) spin components measured in various *k* regions, A–J as shown in **Figure 10a**. As displayed in **Figure 10b** and **e**, the spin‐resolved

**Figure 9.** (a) Band dispersion of the ARPES spectra of Bi/Si(1 1 1) along the *<sup>Γ</sup>*¯¯ M line at *T* = 30 K. Shaded areas indicate the bulk band projection. (b) 2D ARPES intensity plot at *E*<sup>F</sup> as a function of *kx* and *ky* around the *Γ*¯ ¯ M line.

**3. Results and discussion**

bands cross *E*<sup>F</sup>

ets, while the S<sup>2</sup>

point (S<sup>1</sup>

**3.1. Anisotropic Rashba effect of Bi thin film**

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

**Figure 9a** and **b** shows the band dispersion along the Γ¯ ¯

), surrounding elongated pockets (S<sup>2</sup>

Judged from the band dispersion in **Figure 9a**, the S<sup>1</sup>

This section focuses on the spin structure of the S<sup>2</sup>

in **Figure 10a**. In **Figure 10b** and **e**, we display the near‐*E*<sup>F</sup>

**Figure 9.** (a) Band dispersion of the ARPES spectra of Bi/Si(1 1 1) along the *<sup>Γ</sup>*¯¯

bulk band projection. (b) 2D ARPES intensity plot at *E*<sup>F</sup>

At first, we have performed normal (non–spin‐resolved) ARPES measurement of the Bi thin film in order to check the sample quality and geometry, since even a subtle misalignment of the sample orientation would cause a significant error in determining the spin polarization.

Areas with hatched lines in **Figure 9a** are bulk band projection. We distinctly see that several

the in‐plane spin has a vortical structure and isotropic magnitude as denoted by black arrows

curves (EDCs) for the in‐plane (*y*/*x*) and out‐of‐plane (*z*) spin components measured in various *k* regions, A–J as shown in **Figure 10a**. As displayed in **Figure 10b** and **e**, the spin‐resolved

and three kinds Fermi surfaces exist; a hexagonal Fermi surface centered at the Γ¯

and S<sup>3</sup>

to a hole pocket, all of which arise from the spin‐split surface states [8, 10–14].

) and the ellipsoidal pocket (S<sup>3</sup>

M line and Fermi surface of Bi/Si(1 1 1).

are attributed to the electron pock‐

spin‐resolved energy distribution

M line at *T* = 30 K. Shaded areas indicate the

M line.

as a function of *kx* and *ky* around the *Γ*¯ ¯

band. In a conventional 2D Rashba model,

) near the ¯

M point.

**Figure 10.** (a) Fermi surface of Bi/Si(1 1 1) around the Γ¯ point. Red lines are guides for the Fermi surface of the surface bands. Lines A–N represent the *k* region where the spin‐resolved EDCs in (b) and (e) were obtained. Expected spin configuration from the normal Rashba spin‐obit coupling is indicated by black arrows. Spin‐resolved EDCs in regions (b) A–D, (c) E–H, (d) I, J and (e) K–N, respectively.

EDC mainly consists of two components: a slope‐like feature, which rapidly increases its intensity at the binding energy (*E*B) higher than 0.15 eV corresponding to the tail of the QWS, and a weaker broad feature at *E*<sup>F</sup> ‐0.1 eV assigned as the S<sup>2</sup> band.

First, we take a closer look at the spin polarization in regions A–D. In regions A and B, as seen in **Figure 10b**, the in‐plane spin polarization of the S<sup>2</sup> band is dominated by the up spin. On the other hand, the down spin is barely superior in regions C and D. It means that the in‐plane spin structure is qualitatively consistent with the Rashba picture [13–14]. However, it seems that the spin polarization is markedly suppressed in C and D. In fact, the magnitude of the spin polarization along the *y* direction |*P*<sup>y</sup> | is 0.5–0.7 in regions A and B, while it is 0.2–0.3 in regions C and D. This is unexpected since |*P*<sup>y</sup> | should keep the same value across the Γ¯ point in the normal Rashba picture. Then, we focus on the out‐of‐plane (*z*) spin component. We immediately notice that there exists a sizable up‐spin polarization, while theoretically predicted out‐of‐plane spin polarization for simple isotropic Rashba system is zero. More sur‐ prisingly, the magnitude of the z‐axis spin polarization |*P*<sup>z</sup> | (0.4–0.7) is as large as that of |*P*<sup>y</sup> |.

We have also observed a finite in‐plane and out‐of‐plane spin polarization in regions E–H. As shown in **Figure 10c**, the estimated maximum |*P*<sup>y</sup> | of the S<sup>2</sup> band in regions E and F (0.1–0.2) is much smaller than that in regions G and H (0.4–0.5), and it is similar to the result of regions A–D. But then, the down‐spin component dominates for out‐of‐plane spin polar‐ ization in regions E–H. This result suggests that *P*<sup>z</sup> has opposite spin direction across the Γ¯ ¯ M line: the up‐spin component is superior in the negative *k*<sup>y</sup> (regions A–D), while the down‐spin component is superior in the positive *k*<sup>y</sup> (regions E–H).

The data in regions I and J are also the same trend as shown in **Figure 10d**. The difference in |*P*<sup>y</sup> | across the Γ¯ point is also recognized by comparing the EDCs between regions I and J. In addition, the value of |*P*<sup>z</sup> |(~zero) is a good agreement in regions I and J where the up‐ and down‐spin components almost overlap with each other because of the cancellation of two opposite spins. To see if a similar trend is observed along another high‐symmetry line Γ¯ ¯ K, we demonstrate in **Figure 10e** the spin‐resolved EDCs measured in regions K–N. It is apparent that the sign of *P*<sup>z</sup> (also *P*<sup>x</sup> ) in region K (L) is the same as that in the region N (M), indicating that *P*<sup>z</sup> (*P*<sup>x</sup> ) does not switch the sign across the Γ¯ ¯ K line.

**Figure 11** shows a schematical view of the spin polarization vectors of the S<sup>2</sup> band from **Figure 10**. We symmetrized the data by taking into account the threefold crystal symmetry due to the presence of a second bismuth layer. The in‐plane spin component has a vortical structure, but the magnitude of the spin polarization is perpendicular to *k*, called here *P*θ. *P*<sup>z</sup> has a large component and switches the sign by every 60° step. These features would lead to the periodic oscillation of *P*θ and *P*<sup>z</sup> , unlike the general Rashba SOC where *P*θ = const and *P*z  = 0. Now, we discuss the origin of anomalous Rashba effect of Bi thin film. In the conven‐ tional Rashba effect, the in‐plane spin polarization of S<sup>2</sup> is symmetry with respect to the Γ¯ point. In this study, the TRS is not broken by some magnetic impurities because the Bi and Si are nonmagnetic materials. Another possibility of causing the broken TRS is the local surface conditions or the final‐state effect [35, 36]. However, this possibility might be also unlikely by the reproducibility of data and comparing with previous studies. Here it is noted that the LEED pattern of Bi/Si(1 1 1) in **Figure 7** shows the threefold symmetry due to the bilayer‐Bi

**Figure 11.** Schematic view of the spin structure of the S<sup>2</sup> hole pocket in Bi/Si(1 1 1). Size of arrows roughly scales with the observed spin polarization. Data are folded by taking into account the threefold symmetric variation of the bilayer‐ crystal structure.

crystal structure, and we conjecture that the origin of the in‐plane spin asymmetry would be related to the crystal symmetry at present. Next, we discuss the out‐of‐plane spin polar‐ ization. Unlike in‐plane spin polarization, the threefold symmetry in the out‐of‐plane spin component is reported, and sign‐switching behavior is in good agreement with the previous studies in Bi/Ag(1 1 1) and Bi1‐*<sup>x</sup>* Sb*<sup>x</sup>* [15, 37]. The spin structure of Bi surface state is also similar to the spin‐resolved ARPES experiments on the hexagonally warped Dirac‐cone Fermi sur‐ face of the topological insulator as well as the prediction of the *k·p* theory [38, 39]. On the other hand, the absolute value of the out‐of‐plane spin polarization obtained by our experi‐ ment is 40–70%, while the theoretical value of previous studies is a few percent. In order to clarify the spin structure of Bi Rashba effect, we need to evolve theoretical and experimental studies.

#### **3.2. Rashba effect at interface of a Bi Thin film on Si(1 1 1)**

EDC mainly consists of two components: a slope‐like feature, which rapidly increases its intensity at the binding energy (*E*B) higher than 0.15 eV corresponding to the tail of the QWS,

First, we take a closer look at the spin polarization in regions A–D. In regions A and B, as seen

the other hand, the down spin is barely superior in regions C and D. It means that the in‐plane spin structure is qualitatively consistent with the Rashba picture [13–14]. However, it seems that the spin polarization is markedly suppressed in C and D. In fact, the magnitude of the

point in the normal Rashba picture. Then, we focus on the out‐of‐plane (*z*) spin component. We immediately notice that there exists a sizable up‐spin polarization, while theoretically predicted out‐of‐plane spin polarization for simple isotropic Rashba system is zero. More sur‐

We have also observed a finite in‐plane and out‐of‐plane spin polarization in regions E–H.

(0.1–0.2) is much smaller than that in regions G and H (0.4–0.5), and it is similar to the result of regions A–D. But then, the down‐spin component dominates for out‐of‐plane spin polar‐

 (regions E–H). The data in regions I and J are also the same trend as shown in **Figure 10d**. The difference in


down‐spin components almost overlap with each other because of the cancellation of two opposite spins. To see if a similar trend is observed along another high‐symmetry line Γ¯ ¯

demonstrate in **Figure 10e** the spin‐resolved EDCs measured in regions K–N. It is apparent

**Figure 10**. We symmetrized the data by taking into account the threefold crystal symmetry due to the presence of a second bismuth layer. The in‐plane spin component has a vortical structure, but the magnitude of the spin polarization is perpendicular to *k*, called here *P*θ. *P*<sup>z</sup> has a large component and switches the sign by every 60° step. These features would lead

 = 0. Now, we discuss the origin of anomalous Rashba effect of Bi thin film. In the conven‐

point. In this study, the TRS is not broken by some magnetic impurities because the Bi and Si are nonmagnetic materials. Another possibility of causing the broken TRS is the local surface conditions or the final‐state effect [35, 36]. However, this possibility might be also unlikely by the reproducibility of data and comparing with previous studies. Here it is noted that the LEED pattern of Bi/Si(1 1 1) in **Figure 7** shows the threefold symmetry due to the bilayer‐Bi

**Figure 11** shows a schematical view of the spin polarization vectors of the S<sup>2</sup>

K line.

band.



) in region K (L) is the same as that in the region N (M), indicating

, unlike the general Rashba SOC where *P*θ = const and

is symmetry with respect to the Γ¯

band is dominated by the up spin. On



has opposite spin direction across the Γ¯ ¯

(regions A–D), while the down‐spin


band in regions E and F


M

K, we

band from

‐0.1 eV assigned as the S<sup>2</sup>

and a weaker broad feature at *E*<sup>F</sup>

in **Figure 10b**, the in‐plane spin polarization of the S<sup>2</sup>

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

spin polarization along the *y* direction |*P*<sup>y</sup>

in regions C and D. This is unexpected since |*P*<sup>y</sup>

prisingly, the magnitude of the z‐axis spin polarization |*P*<sup>z</sup>

As shown in **Figure 10c**, the estimated maximum |*P*<sup>y</sup>

line: the up‐spin component is superior in the negative *k*<sup>y</sup>

) does not switch the sign across the Γ¯ ¯

tional Rashba effect, the in‐plane spin polarization of S<sup>2</sup>

ization in regions E–H. This result suggests that *P*<sup>z</sup>

component is superior in the positive *k*<sup>y</sup>

(also *P*<sup>x</sup>

to the periodic oscillation of *P*θ and *P*<sup>z</sup>

addition, the value of |*P*<sup>z</sup>

that the sign of *P*<sup>z</sup>

(*P*<sup>x</sup>


that *P*<sup>z</sup>

*P*z

In this section, we focus on the spin structure around the ¯ M point of the Bi/Si(1 1 1) thin film. As shown in **Figure 12a**, the band dispersion of the QWSs dramatically alters around the ¯ M point as a function of the film thickness *d*, and we prepared five different thickness film for *d* = 8, 10, 15, 20 and 40 BL. All of the samples, the bottom energy of an electron‐like dispersion is located at ~30 meV below *E*<sup>F</sup> . This band is well separated from the hole‐like band at higher *E*<sup>B</sup> arising from the QWSs that originate from the confinement of the wave function in the direc‐ tion perpendicular to the surface (*z* direction). There is no evidence for the Kramers degeneracy of the S<sup>3</sup> band or the QWSs at the ¯ M point (*k*<sup>y</sup>  = 0). The absence of the Kramers degeneracy at the ¯ M point has also been observed in the previous research of the Bi, and it is explained by the hybridization of the surface states and the QWS [9, 12]. As clearly visible, we identify the QWSs at *E*B = 0.1–0.25 eV for *d* = 40 BL. Upon decreasing *d*, the top of the highest QWS shifts down‐ ward, and the energy separation of each QWS becomes wider, enabling us to unambiguously identify the overall energy dispersion of the individual QWS in the band dispersion plots.

In conventional Rashba picture, the in‐plane spin polarization of the S<sup>3</sup> band is dominated by the up spin independently of a thickness at the cut 1 shown in **Figure 12b**. **Figure 12c** shows the spin‐resolved EDCs and spin polarization of *P*<sup>y</sup> for various *d* values of 8–40 BL. All of the spin‐resolved EDCs consist of the tail of the QWSs at *E*B > 0.1 eV and the S<sup>3</sup> electron band at ~50 meV. As seen in **Figure 12c**, the in‐plane spin‐resolved EDCs of the S<sup>3</sup> band for 10–40 BL are dominated by the up spin, and this indicates that the in‐plane spin direction is qualitatively consistent with the normal Rashba picture. However, our result clearly demon‐ strates that the observed *P*<sup>y</sup> value strongly depends on *d*. The obtained *P*<sup>y</sup> value at *d* = 40 BL has a maximum at ~0.7, then gradually reduces at *d* = 15 BL (*P*<sup>y</sup> ~0.3) and finally reaches ~0 at *d* = 8 BL. It is noted that the much smaller spin polarization in the above‐*E*<sup>F</sup> region for *d* = 15 BL as compared to that for *d* = 40 BL might be due to much weaker peak weight in the original EDC and an intrinsic suppression of the spin polarization, which would lead to a relative enhancement of the background weight and less statistical reliability of the states above *E*<sup>F</sup> .

We discuss the physical mechanism behind the unusual thickness dependence of *P*. The cou‐ pling between the Bi film and the Si(1 1 1) substrate is fairly weak due to the presence of a dis‐ ordered wetting layer at the interface, suggesting that the Bi film is nearly freestanding [34]. Namely, the Bi/Si interface is also broken SIS, and it can be thought of as another (bottom) surface as with vacuum‐side (top) surface. These Rashba states of top and bottom surface should have opposite spin directions as illustrated in **Figure 13**. In case of bulk Bi, which is thick enough, the top and bottom Rashba states do not interfere with each other. On reducing

**Figure 12.** (a) Thickness dependence of the band dispersion near *E*<sup>F</sup> along the ¯ K¯ M ¯ K line, obtained by taking the second derivative of the EDCs at 30 K. EDC at the ¯ M point is also shown for each thickness. (b) The Fermi surface around the ¯ M point. Blue and red arrows indicate expected spin configuration of the normal Rashba effect. (c) Corresponding spin‐ resolved EDCs and their spin polarizations for cut 1 in (b).

**Figure 13.** Schematic view of the spin vectors for the Fermi surface at the top and bottom surfaces.

thickness, the wave functions of the two Rashba states overlap and hybridize, which means that the up‐ and down‐spin states merge and the *P* observed by spin‐resolved ARPES would become small. This picture is a very simple model but explains in a nice way the thickness dependence of *P*. We also found that the decay length of the surface‐state wave function is around at least 20 BL (~80 Å) since the reduction in experimental *P* is already started in 20 BL (**Figure 12**). A thin film of Bi<sup>2</sup> Se<sup>3</sup> topological insulator also exhibits a similar decay length, as inferred from the experimental fact that the hybridization gap starts to open at six quintuple layers (~60 Å) [40]. Our result also suggests that the Rashba effect of Bi/Si interface is useful for the spintronic devices because the interface is generally more stable than the surface.

### **3.3. 1D edge state with Rashba effect in Bi thin film**

hybridization of the surface states and the QWS [9, 12]. As clearly visible, we identify the QWSs at *E*B = 0.1–0.25 eV for *d* = 40 BL. Upon decreasing *d*, the top of the highest QWS shifts down‐ ward, and the energy separation of each QWS becomes wider, enabling us to unambiguously identify the overall energy dispersion of the individual QWS in the band dispersion plots.

by the up spin independently of a thickness at the cut 1 shown in **Figure 12b**. **Figure 12c**

10–40 BL are dominated by the up spin, and this indicates that the in‐plane spin direction is qualitatively consistent with the normal Rashba picture. However, our result clearly demon‐

has a maximum at ~0.7, then gradually reduces at *d* = 15 BL (*P*<sup>y</sup> ~0.3) and finally reaches ~0 at

as compared to that for *d* = 40 BL might be due to much weaker peak weight in the original EDC and an intrinsic suppression of the spin polarization, which would lead to a relative enhancement of the background weight and less statistical reliability of the states above *E*<sup>F</sup>

We discuss the physical mechanism behind the unusual thickness dependence of *P*. The cou‐ pling between the Bi film and the Si(1 1 1) substrate is fairly weak due to the presence of a dis‐ ordered wetting layer at the interface, suggesting that the Bi film is nearly freestanding [34]. Namely, the Bi/Si interface is also broken SIS, and it can be thought of as another (bottom) surface as with vacuum‐side (top) surface. These Rashba states of top and bottom surface should have opposite spin directions as illustrated in **Figure 13**. In case of bulk Bi, which is thick enough, the top and bottom Rashba states do not interfere with each other. On reducing

value strongly depends on *d*. The obtained *P*<sup>y</sup>

along the ¯

point. Blue and red arrows indicate expected spin configuration of the normal Rashba effect. (c) Corresponding spin‐

K¯ M ¯

M point is also shown for each thickness. (b) The Fermi surface around the ¯

K line, obtained by taking the second

M

All of the spin‐resolved EDCs consist of the tail of the QWSs at *E*B > 0.1 eV and the S<sup>3</sup>

band at ~50 meV. As seen in **Figure 12c**, the in‐plane spin‐resolved EDCs of the S<sup>3</sup>

band is dominated

value at *d* = 40 BL

region for *d* = 15 BL

electron

band for

.

for various *d* values of 8–40 BL.

In conventional Rashba picture, the in‐plane spin polarization of the S<sup>3</sup>

*d* = 8 BL. It is noted that the much smaller spin polarization in the above‐*E*<sup>F</sup>

shows the spin‐resolved EDCs and spin polarization of *P*<sup>y</sup>

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

**Figure 12.** (a) Thickness dependence of the band dispersion near *E*<sup>F</sup>

resolved EDCs and their spin polarizations for cut 1 in (b).

derivative of the EDCs at 30 K. EDC at the ¯

strates that the observed *P*<sup>y</sup>

As observed by the atomic force microscopy of our Bi thin film (**Figure 14**), triangular‐shaped Bi BL islands with typically ~0.1 μm edge length are formed on the top surface of the Bi thin film as reported previously [34], and the edge of each island is perpendicular to the Γ¯ <sup>M</sup>¯ direc‐ tion considering sample geometry; namely, the edge runs along the Γ¯K¯ direction in the *k* space. According to previous experimental and theoretical studies, there are no bulk and surface states near the K¯ point, and we might observe some electronic state if edge state exists along the Γ¯K¯ line.

We demonstrate the band dispersion of Bi thin film along the Γ¯¯ K line for *d* = 15 BL in **Figure <sup>15</sup>**. With a careful look at the region between Γ¯¯ K line (yellow rectangle in **Figure 15a**), one finds unexpected faint intensity displaying a finite energy dispersion, as better illustrated with enhanced color contrast in the inset. To establish the energy dispersion of this unex‐ pected feature, we have measured the ARPES data along several cuts in the surface Brillouin zone (**Figure 15b**). A careful band searching with the enhanced intensity scale (**Figure 15c**) shows the band dispersion with a characteristic *x*‐shape along cuts 1–3. The intersection of

**Figure 14.** The atomic force microscopy image of a Bi thin film (*d* = 15 BL) at *T* = 300 K.

the x‐shaped band is at *k*<sup>y</sup> ~ 0.7 Å‐1 not at the high‐symmetry points of the surface and bulk Brillouin zone. Intriguingly, this band is robust against the change in the *k*<sup>x</sup> location of cut (cuts 1–3). This demonstrates that the *x*‐shaped dispersion has a 1D character along the *k*<sup>y</sup> direction. Indeed, the observed band structure along cut 4 (perpendicular to cuts 1–3) shows no obvious dispersion, confirming the 1D nature. To experimentally clarify the spin‐split nature of the edge band, we have performed a spin‐resolved ARPES experiment. As shown in **Figure 15d** which plots the spin‐resolved EDCs at a representative *k*<sup>y</sup> point (marked by a pink line in **Figure 15c**), we clearly find a difference between the up‐ and down‐spin spectra in both the in‐plane and out‐of‐plane components.

As for the origin of the unexpected 1D state, we have taken into account various possibilities such as the mixture of domains with different film thickness, surface reconstruction, slight isolation of the topmost Bi bilayer and surface stacking faults. A most natural and convincing explanation is that it originates from the edge states of Bi bilayer. To further strengthen our conclusion, we have carried out first‐principles electronic band structure calculations for a specific crystal structure (**Figure 16**). Electronic band structure calculations were carried out by means of a first‐principles density functional theory approach with the all‐electron full‐poten‐ tial linearized augmented‐plane‐wave method in the scalar relativistic scheme. The spin‐orbit coupling was included as the second variation in the self‐consistent‐field iterations. Thin‐film systems were simulated by adopting periodic slab models with sufficiently thick vacuum layer. In this model, Bi atoms in 1D allay are removed from the topmost Bi 1BL (**Figure 16a**) so as to reproduce the infinitely long edge structure along the *y* direction. Assumption of such an idealized model crystal is turned out to be sufficient for reasonably simulating the edge band structure. As shown in **Figure 16c**, the Brillouin zone for this model crystal structure has a rectangular shape. The vertical length of Brillouin zone is the same as the ¯ <sup>M</sup> ¯ <sup>M</sup> interval since the size of unit cell along *y*‐axis is the same for the edge structure and the Bi thin film. On the other hand, the horizontal Brillouin zone length has no important physical role because the unit‐cell length for *x*‐axis was simply chosen for the sake of calculations. Thus,

Anomalous Rashba Effect of Bi Thin Film Studied by Spin-Resolved ARPES http://dx.doi.org/10.5772/66278 93

the x‐shaped band is at *k*<sup>y</sup> ~ 0.7 Å‐1

the in‐plane and out‐of‐plane components.

Brillouin zone. Intriguingly, this band is robust against the change in the *k*<sup>x</sup>

**Figure 15d** which plots the spin‐resolved EDCs at a representative *k*<sup>y</sup>

**Figure 14.** The atomic force microscopy image of a Bi thin film (*d* = 15 BL) at *T* = 300 K.

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

(cuts 1–3). This demonstrates that the *x*‐shaped dispersion has a 1D character along the *k*<sup>y</sup> direction. Indeed, the observed band structure along cut 4 (perpendicular to cuts 1–3) shows no obvious dispersion, confirming the 1D nature. To experimentally clarify the spin‐split nature of the edge band, we have performed a spin‐resolved ARPES experiment. As shown in

line in **Figure 15c**), we clearly find a difference between the up‐ and down‐spin spectra in both

As for the origin of the unexpected 1D state, we have taken into account various possibilities such as the mixture of domains with different film thickness, surface reconstruction, slight isolation of the topmost Bi bilayer and surface stacking faults. A most natural and convincing explanation is that it originates from the edge states of Bi bilayer. To further strengthen our conclusion, we have carried out first‐principles electronic band structure calculations for a specific crystal structure (**Figure 16**). Electronic band structure calculations were carried out by means of a first‐principles density functional theory approach with the all‐electron full‐poten‐ tial linearized augmented‐plane‐wave method in the scalar relativistic scheme. The spin‐orbit coupling was included as the second variation in the self‐consistent‐field iterations. Thin‐film systems were simulated by adopting periodic slab models with sufficiently thick vacuum layer. In this model, Bi atoms in 1D allay are removed from the topmost Bi 1BL (**Figure 16a**) so as to reproduce the infinitely long edge structure along the *y* direction. Assumption of such an idealized model crystal is turned out to be sufficient for reasonably simulating the edge band structure. As shown in **Figure 16c**, the Brillouin zone for this model crystal structure has a rectangular shape. The vertical length of Brillouin zone is the same as the ¯ <sup>M</sup> ¯ <sup>M</sup>

interval since the size of unit cell along *y*‐axis is the same for the edge structure and the Bi thin film. On the other hand, the horizontal Brillouin zone length has no important physical role because the unit‐cell length for *x*‐axis was simply chosen for the sake of calculations. Thus,

not at the high‐symmetry points of the surface and bulk

location of cut

point (marked by a pink

**Figure 15.** (a) Band structure of a Bi thin film along the Γ¯ ¯ K high‐symmetry lines. To better trace the weak signal around the ¯ K points, an enhanced color scale image is shown in the inset. (b) ARPES intensity plot at *E*<sup>F</sup> around the Γ¯ ¯ K line. (c) Band dispersion near EF along cuts 1–4 in (b), respectively. (d) Spin‐resolved EDCs for the in‐plane and out‐of‐plane spin components at the *k* point indicated by the pink line in (c), together with the corresponding energy dependence of the spin polarization.

the high‐symmetry point ¯Y now becomes the time‐reversal‐invariant momentum and is located exactly at the horizontal projection of the ¯ M point (*k*<sup>y</sup>  = 0.69 Å‐1 ), which actually coincides with the intersection of the *x*‐shaped band. **Figure 16c** displays the calculated band dispersion along the Γ ¯ ¯Y direction of Bi thin film for *d* = 10 BL. We identify two prominent dispersive bands, which cross *E*<sup>F</sup> and have the edge‐state origin. Moreover, the edge bands show the Rashba spin splitting due to the strong spin‐orbit coupling at the edge, as evidenced by the degeneracy at the ¯Y point. We found that the overall spin‐vector direction in the experiment, that is, sign of spin polarization for each component, is also consistent with the

**Figure 16.** (a) Side and top views of the artificially constructed model crystal structure used to calculate the energy band structure of the edge state. Area enclosed by gray and red solid lines shows the unit cell. Crystal structure in the unit cell contains 1 BL Bi ribbon on 3 BL Bi where hydrogen atoms terminate one side of the edge in Bi ribbon. (b) The 1D Brillouin zone (red line) of the model crystal structure shown in (a), compared to the 2D hexagonal surface Brillouin zone (black line). (c) Band calculations for the model crystal structure along the Γ¯ ¯ Y line. The wide of red or blue curves indicates the concentration of the up‐ and down‐spin electrons along the axis specified in the figure.

calculation, while the absolute magnitude of the spin polarization in the experiment (~0.2) is much smaller than that of the calculation (~0.26 to 0.64) in both the in‐plane and out‐of‐plane components, likely due to a finite contribution from the angle‐integrated‐type background in the ARPES spectra and the spin‐orbit entanglement effect [41]. Finally, we have estimated the Rashba parameter αR by numerical simulation for the 1D parabolic band with the SOC and obtained αR = 0.80 ± 0.05 eVÅ (**Figure 17**). This value is much larger than that for the 2D sur‐ face state (0.56 eVÅ) [11, 15], and the difference could be explained in terms of the presence of an in‐plane potential gradient [15, 42] at the edge in addition to the out‐of‐plane component, which already exists in the 2D film, as supported by observation of the out‐of‐plane spin polarization as large as the in‐plane counterpart. Recently, STM study on the edge states of Bi crystal and thin film [43, 44]. The STM study reported that one of two types of edges, where Bi atoms are terminated at close to vacuum, which corresponds to our calculation, has a 1D character. The band dispersion of the edge state observed in ARPES shows a good agreement with the calculated bands, suggesting that both ARPES and STM observe the same edge state. It is true that 1D electronic state that observed our ARPES measurement is localized at edge.

**Figure 17.** Experimental band dispersion for cuts 1–3 extracted from the peak position of the EDCs in Figure 15c, compared with the numerical simulation for a simple 1D Rashba splitting (dashed black curve).

The edge state in Bi would be also useful in realizing novel physical properties and new spintronic devices.

### **4. Conclusions**

**Figure 16.** (a) Side and top views of the artificially constructed model crystal structure used to calculate the energy band structure of the edge state. Area enclosed by gray and red solid lines shows the unit cell. Crystal structure in the unit cell contains 1 BL Bi ribbon on 3 BL Bi where hydrogen atoms terminate one side of the edge in Bi ribbon. (b) The 1D Brillouin zone (red line) of the model crystal structure shown in (a), compared to the 2D hexagonal surface Brillouin zone (black

calculation, while the absolute magnitude of the spin polarization in the experiment (~0.2) is much smaller than that of the calculation (~0.26 to 0.64) in both the in‐plane and out‐of‐plane components, likely due to a finite contribution from the angle‐integrated‐type background in the ARPES spectra and the spin‐orbit entanglement effect [41]. Finally, we have estimated the Rashba parameter αR by numerical simulation for the 1D parabolic band with the SOC and obtained αR = 0.80 ± 0.05 eVÅ (**Figure 17**). This value is much larger than that for the 2D sur‐ face state (0.56 eVÅ) [11, 15], and the difference could be explained in terms of the presence of an in‐plane potential gradient [15, 42] at the edge in addition to the out‐of‐plane component, which already exists in the 2D film, as supported by observation of the out‐of‐plane spin polarization as large as the in‐plane counterpart. Recently, STM study on the edge states of Bi crystal and thin film [43, 44]. The STM study reported that one of two types of edges, where Bi atoms are terminated at close to vacuum, which corresponds to our calculation, has a 1D character. The band dispersion of the edge state observed in ARPES shows a good agreement with the calculated bands, suggesting that both ARPES and STM observe the same edge state. It is true that 1D electronic state that observed our ARPES measurement is localized at edge.

Y line. The wide of red or blue curves indicates the

line). (c) Band calculations for the model crystal structure along the Γ¯ ¯

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

concentration of the up‐ and down‐spin electrons along the axis specified in the figure.

We have demonstrated anomalous Rashba effect of Bi thin film on Si(1 1 1) by using spin‐ resolved ARPES. Major findings of the present work are the following three features: (i) the surface Rashba states of Bi exhibit the asymmetric in‐plane spin polarization and the giant out‐of‐plane spin polarization, (ii) the spin polarization of the surface states is reduced on decreasing thickness and (iii) 1D band dispersion from the edge state of Bi islands on Si(1 1 1) exhibits large Rashba effect. These observed peculiar spin states are not explained in terms of the conventional Rashba effect, and these results open a pathway for realizing exotic physical properties at the strong SOC systems.

### **Acknowledgements**

This study is a collaborative research with Prof. Takahashi, Associate Prof. Sato and Associate Prof. Souma in Tohoku University and Prof. Oguchi in Osaka University. We thank K. Sugawara and K. Kosaka for his assistance in the ARPES experiment. This work was sup‐ ported by JSPS, MEXT of Japan and the Mitsubishi foundation.

### **Author details**

Akari Takayama

Address all correspondence to: a.takayama@surface.phys.s.u‐tokyo.ac.jp

The University of Tokyo, Tokyo, Japan

### **References**


[9] Sugawara K, Sato T, Souma S, Takahashi T, Arai M, and Sasaki T. Fermi surface and anisotropic spin‐orbit coupling of Sb(1 1 1) studied by angle‐resolved photoemission spectroscopy. Phys. Rev. Lett. 2006;**96**:046411. doi:10.1103/PhysRevLett.96.046411

the conventional Rashba effect, and these results open a pathway for realizing exotic physical

This study is a collaborative research with Prof. Takahashi, Associate Prof. Sato and Associate Prof. Souma in Tohoku University and Prof. Oguchi in Osaka University. We thank K. Sugawara and K. Kosaka for his assistance in the ARPES experiment. This work was sup‐

[1] Baibich MN, Broto JM, Fert A, Nguyen Van Dau F, Petroff F, Eitenne P, Creuzet G, Friedrich A, and Chazelas J. Giant magnetoresistance of (001)Fe/(001)Cr magnetic super‐

[2] Awschalom D and Samarth N. Trend: Spintronics without magnetism. Phyics 2009;**2**:50.

[3] Datta S and Das B. Electronic analog of the electro‐optic modulator. Appl. Phys. Lett.

[4] Bychkov YA and Rashba EI. Properties of a 2D electron gas with lifted spectral degen‐

[5] Hasan MZ and Kane CL. Colloquium: topological insulators. Rev. Mod. Phys.

[6] Qi XL and Zhang SC. Topological insulators and superconductors. Rev. Mod. Phys.

[7] LaShell S, McDougall BA, and Jensen E. Spin splitting of an Au(1 1 1) surface state band observed with angle resolved photoelectron spectroscopy. Phys. Rev. Lett. 1996;**77**:3419.

[8] Hofmann P. The surfaces of bismuth: structural and electronic properties. Prog. Surf. Sci.

lattices. Phys. Rev. Lett. 1988;**61**:2472. doi:10.1103/PhysRevLett.61.2472

ported by JSPS, MEXT of Japan and the Mitsubishi foundation.

Address all correspondence to: a.takayama@surface.phys.s.u‐tokyo.ac.jp

properties at the strong SOC systems.

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

The University of Tokyo, Tokyo, Japan

doi:10.1103/Physics.2.50

1990;**56**:665–667. DOI:10.1063/1.102730

2010;**82**:3045. doi:10.1103/RevModPhys.82.3045

2011;**83**:1057. doi:10.1103/RevModPhys.83.1057

2006;**81**:191. doi:10.1016/j.progsurf.2006.03.001

eracy. JETP Lett. 1984;**39**:78–81.

doi:10.1103/PhysRevLett.77.3419

**Acknowledgements**

**Author details**

Akari Takayama

**References**


the direction perpendicular to the surface. Phys. Rev. Lett. 2009;**102**:096805. doi:10.1103/ PhysRevLett.102.096805


[33] Viernow J, Lin JL, Petrovykh DY, Leibsle FM, Men FK, and Himpsel FJ. Regular step arrays on silicon. Appl. Phys. Lett. 1998;**72**:948. doi:10.1063/1.120882

the direction perpendicular to the surface. Phys. Rev. Lett. 2009;**102**:096805. doi:10.1103/

[21] Yaji K, Ohtsubo Y, Hatta S, Okuyama H, Miyamoto K, Okuda T, Kimura A, Namatame T, Taniguchi M, and Aruga T. Large Rashba spin splitting of a metallic surface‐state band on a semiconductor surface. Nature Commun. 2010;**1**:17. doi:10.1038/ncomms1016 [22] Matetskiy AV, Ichinokura S, Bondarenko LV, Tupchaya AY, Gruznev DV, Zotov AV, Saranin AA, Hobara R, Takayama A, and Hasegawa S. Two‐dimensional superconduc‐ tor with a giant Rashba Effect: one‐atom‐layer Tl‐Pb compound on Si(1 1 1). Phys. Rev.

[23] Barke I, Zheng F, Rughheimer TK, and Himpsel FJ. Experimental evidence for spin‐split bands in a one‐dimensional chain structure. Phys. Rev. Lett. 2006;**97**:226405. doi:10.1103/

[24] Wells JW, Dil JH, Meier F, Lobo‐Checa J, Petrov VN, Osterwalder J, Ugeda MM, Fernandez‐ Torrente I, Pascual JI, Rienks EDL, Jensen MF, and Hofmann Ph. Nondegenerate metallic states on Bi(1 1 4): a one‐dimensional topological metal. Phys. Rev. Lett. 2009;**102**:096802.

[25] Quay CHL, Hughes TL, Sulpizio JL, Pfeiffer LN, Baldwin KW, West KW, Goldhaber‐ Gordon D, and de Picciotto R. Observation of a one‐dimensional spin–orbit gap in a

[26] Pershin YV, Nesteroff JA, and Privman V. Effect of spin‐orbit interaction and in‐plane magnetic field on the conductance of a quasi‐one‐dimensional system. Phys. Rev. B.

[27] Takayama A, Sato T, Souma S, and Takahashi T. Giant out‐of‐plane spin component and the asymmetry of spin polarization in surface Rashba States of Bismuth thin film. Phys.

[28] Takayama A, Sato T, Souma S, Oguchi T, and Takahashi T. Tunable spin polarization in bismuth ultrathin film on Si(1 1 1). Nano Lett. 2012;**12**:1776. doi:10.1021/nl2035018 [29] Takayama A, Sato T, Souma S, Oguchi T, and Takahashi T. One‐dimensional edge states with giant spin splitting in a bismuth thin film. Phys. Rev. Lett. 2015;**114**:066402.

[30] Souma S, Takayama A, Sugawara K, Sato T, and Takahashi T. Ultrahigh‐resolution spin‐resolved photoemission spectrometer with a mini Mott detector. Rev. Sci. Instrum.

[31] Schlier E and Farnsworth HE. Structure and adsorption characteristics of clean surfaces of germanium and silicon. J. Chem. Phys. 1959;**30**:917. doi:10.1063/1.1730126

[32] Takayanagi K, Tanishiro Y, Takahashi M, and Takahashi S. Structural analysis of Si(1 1 1)‐7×7 by UHV‐transmission electron diffraction and microscopy. J. Vac. Sci.

Lett. 2015;**115**:147003. doi:10.1103/PhysRevLett.115.147003

quantum wire. Nat. Phys. 2010;**6**:336. doi:10.1038/nphys1626

Rev. Lett. 2011;**106**:166401. doi:10.1103/PhysRevLett.106.166401

2004;**69**:121306(R). doi:10.1103/PhysRevB.69.121306

PhysRevLett.102.096805

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

PhysRevLett.97.226405.

doi:10.1103/PhysRevLett.102.096802

doi:10.1103/PhysRevLett.114.066402

2010;**81**:095101. doi:10.1063/1.3480542

Technol. A. 1985;**3**:1502. doi:10.1116/1.573160


**Provisional chapter**

### **Electrochemical Deposition of P3AT Films Used as a Probe of Optical Properties in Polymeric System Electrochemical Deposition of P3AT Films Used as a Probe of Optical Properties in Polymeric System**

Sankler Soares de Sá, Fernando Costa Basílio, Henrique de Santana, Alexandre Marletta and Eralci Moreira Therézio Henrique de Santana, Alexandre Marletta and Eralci Moreira Therézio Additional information is available at the end of the chapter

Sankler Soares de Sá, Fernando Costa Basílio,

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/66921

#### **Abstract**

Poly(3‐alkylthiophene) (P3ATs) have been extensively used in photovoltaic devices such as a p‐type organic Semiconductors. However, several electronic properties of P3ATs present energy transfer inter‐ and intra‐chains that have direct consequences on the performance of optoelectronic devices. Traditionally electrochemical techniques, such as cyclic voltammetry, chronoamperometry and chronocoulometry, have been applied to process polymer thin films and unconventional spectroscopy techniques are used to characterize the electronic properties. In the present work, we used an innovative tech‐ nique called ellipsometry emission to investigate the optical properties of P3AT films. We propose a new approach to study the electrochemical synthesize and unintentional doping processes of polymeric systems. We showed a strong correlation between the electrochemical synthesis and the optical properties controlling the film growth condi‐ tions for P3ATs. The results obtained in the present study can be potentially utilized for applications in organic devices, mainly in photovoltaic cells when the film deposition and the optical properties control are relevant.

**Keywords:** poly(3‐alkylthiophene), electrochemical synthesis, optical properties, energy transfer, emission ellipsometry

### **1. Introduction**

Over the last decades, semiconductor polymers have attracted considerable interest, partic‐ ularly for the production of organic electroluminescent diodes (OLEDs) and organic photo‐ voltaic cells (OPVs), in which they present high emission efficiency in the visible region and UV‐Vis absorption in the broad spectral window [1, 2]. Devices using conductive polymers

© 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.

exhibit some advantages over inorganic semiconductors. They are easily deposited on thin solid films by low‐cost techniques such as spin coating, casting or electrochemical [2, 3]. Among the wide variety of conductive polymers, poly(3‐*alkylthiophene*) (P3AT) has been studied due to its various physical‐chemistry characteristics, e.g., good chemical stability, solubility (making it an easy deposition material on substrates) and has electrochromic and thermochromic charac‐ teristics [4]. Besides these properties, the luminescence efficiency of this polymer has increased significantly in the function of the *alkyl* chain length [5, 6]. Basically, P3ATs are the derivatives of polythiophene (PT) which are obtained from the polymerization of thiophene (monomer), a sulfur heterocyclic ring [7]. The precursor monomer of P3ATs, 3‐*alkylthiophene* is also composed of thiophene ring and alkyl groups, in compliance with the following combination [5]:

$$\mathbf{C}\_n \mathbf{H}\_{2n+1} \tag{1}$$

where C is the carbon chemical element, H is the hydrogen chemical element and *n* is the number of carbons that compose the molecule.

The P3ATs chemically synthesized presented an energy *gap* of around 1.93 eV (640 nm) [5, 6, 8–12]. Interestingly, this *energy gap* independent of the size of the *alkyl* lateral chain because it is not conjugated. Therefore, the recombination of excited carriers occurs only in the main conjugated polymer chain. However, the intensity of the emission band is directly related to the alkyl chain [6]. Another important observation about the emission band intensity is the anomalous temperature dependence [5].

Ohmori et al. [6] have observed luminescence intensity dependence in function of the length of the *alkyl* chain. They used three P3ATs with different sizes of the *alkyl* chain. The P3ATs have tra‐ ditionally been prepared by chemical synthesis from 3‐AT monomers with FeCl3 as a catalyst. Chemical synthesis of P3AT polymer, using a standard way in the literature, was first obtained by Yoshino et al. in 1984 [13]. In addition, Yoshino et al. [5] noted that the photoluminescence intensity (PL) of P3AT films increases in the function of the sample temperature and decreases after the melting point. This result has been discussed in terms of the effective conjugation length, since the dynamics of the excited species are influenced by the occurrence of a twist between the vicinity of the thiophene rings together with the interchain interaction. In the last decade, the interest in the organic electronic devices has increased significantly; however, some effects on their operation are not fully understood, in particular the interface effects of the sub‐ strate/polymer and energy transfer of excited carriers [9, 14–16]. Since the physical‐chemistry properties and investigation of organic active layers, such as P3ATs thin solid films, can eluci‐ date the development of new optoelectronic devices [16–18]. Interface effects cause significant quenching of excited carriers and it is commonly investigated by conventional spectroscopic techniques [15, 19], such as ultraviolet‐visible absorption (UV‐Vis), photoluminescence (PL), photoluminescence excitation (PLE), vibrational spectroscopy (FT‐IR and RAMAN) [8, 12, 20, 21] and the morphological technique of atomic force microscopy (AFM) [22–24]. In the case of energy transfer processes of excited carriers, the analysis of polarization of emitted light can be directly correlated with polymeric chain position parallel to the direction of the transition dipole moment [25–27]. Moreover, we need to consider the effects of the deposition method.

In present work, we used the electrochemical synthesis for deposition of P3AT thin solid film correlating with optical properties. We demonstrated an easy and efficient alternative method to control the processing of organic optoelectronic devices. In that context, Therézio et al. in [8] show two distinct structures of the polymer chain morphology shifting the band gap to higher energies, using an optical analysis of P3ATs films electrochemically prepared. Moreover, refer‐ ence [9] also indicated that it is possible to analyze the polarization of the emitted light by the supporting electrolyte effects on the emission properties of P3ATs films. The results conclude quantitatively that the best‐supporting electrolyte concentration for the P3ATs film's production is 0.100 mol L‐1 . It is in absolute agreement of the electrolyte concentration used in the litera‐ ture to P3ATs synthesis. Recently, Santana and coworkers [10, 12, 28, 29] have shown that the P3ATs synthesis is possible using different supporting electrolytes, solvents and thicknesses. Therefore, it was possible to correlate the growth conditions of the electrochemical synthesis and optical properties to produce polymer films with possible application in optoelectronic devices. We present the systematic study of two sets of poly(3‐dodecylthiophene) (P3DDT) films grown electrochemically with two different electrolytes. All thin films were deposited on a transparent electrode FTO (fluorine‐doped tin oxide) and cycles ranging from 1 to 10 cycles. In addition, it was possible to obtain the emission optical characteristics in the function of the amount of polymer deposited or the polymeric film thickness. We used, in the present investiga‐ tion, UV‐Vis absorption, photoluminescence and ellipsometry emission. As a result, we assign the use of an alternative optical characterization to probe the organic semiconductors obtained via electrochemical techniques [20, 23].

### **2. Electrochemical deposition in poly(3‐alkylthiophenes) films**

P3DDT samples were deposited on the FTO substrate by the electrochemical synthesis of 3‐ dodecylthiophene (C16H28S) monomer in an electrolyte solution containing a solvent, salt and monomer. In stoke solution, we used acetonitrile (CH3 CN), monomer 3‐dodecylthiophene and lithium perchlorate (LiClO4 ) or tetraethylammonium tetrafluoroborate ((C<sup>2</sup> H5 ) 4 NBF4 or Et4 NBF4 ) salts [9]. The concentrations used are 0.100 mol L‐1 for the support electrolyte (SE), 0.050 mol L‐1 for the monomer and 0.040 mol L‐1 for Et4 NBF4 or LiClO4 electrolyte. The options for the previous concentrations is based on the literature for poly(3‐methylthiophene) and poly(3‐ octylthiophene) [9, 10, 30–32]. The films grown in the present study are labeled in **Table 1**.

### **2.1. Cyclic voltammetry**

exhibit some advantages over inorganic semiconductors. They are easily deposited on thin solid films by low‐cost techniques such as spin coating, casting or electrochemical [2, 3]. Among the wide variety of conductive polymers, poly(3‐*alkylthiophene*) (P3AT) has been studied due to its various physical‐chemistry characteristics, e.g., good chemical stability, solubility (making it an easy deposition material on substrates) and has electrochromic and thermochromic charac‐ teristics [4]. Besides these properties, the luminescence efficiency of this polymer has increased significantly in the function of the *alkyl* chain length [5, 6]. Basically, P3ATs are the derivatives of polythiophene (PT) which are obtained from the polymerization of thiophene (monomer), a sulfur heterocyclic ring [7]. The precursor monomer of P3ATs, 3‐*alkylthiophene* is also composed

of thiophene ring and alkyl groups, in compliance with the following combination [5]:

number of carbons that compose the molecule.

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

anomalous temperature dependence [5].

 C*<sup>n</sup>* H2*<sup>n</sup>*+1 (1) where C is the carbon chemical element, H is the hydrogen chemical element and *n* is the

The P3ATs chemically synthesized presented an energy *gap* of around 1.93 eV (640 nm) [5, 6, 8–12]. Interestingly, this *energy gap* independent of the size of the *alkyl* lateral chain because it is not conjugated. Therefore, the recombination of excited carriers occurs only in the main conjugated polymer chain. However, the intensity of the emission band is directly related to the alkyl chain [6]. Another important observation about the emission band intensity is the

Ohmori et al. [6] have observed luminescence intensity dependence in function of the length of the *alkyl* chain. They used three P3ATs with different sizes of the *alkyl* chain. The P3ATs have tra‐

Chemical synthesis of P3AT polymer, using a standard way in the literature, was first obtained by Yoshino et al. in 1984 [13]. In addition, Yoshino et al. [5] noted that the photoluminescence intensity (PL) of P3AT films increases in the function of the sample temperature and decreases after the melting point. This result has been discussed in terms of the effective conjugation length, since the dynamics of the excited species are influenced by the occurrence of a twist between the vicinity of the thiophene rings together with the interchain interaction. In the last decade, the interest in the organic electronic devices has increased significantly; however, some effects on their operation are not fully understood, in particular the interface effects of the sub‐ strate/polymer and energy transfer of excited carriers [9, 14–16]. Since the physical‐chemistry properties and investigation of organic active layers, such as P3ATs thin solid films, can eluci‐ date the development of new optoelectronic devices [16–18]. Interface effects cause significant quenching of excited carriers and it is commonly investigated by conventional spectroscopic techniques [15, 19], such as ultraviolet‐visible absorption (UV‐Vis), photoluminescence (PL), photoluminescence excitation (PLE), vibrational spectroscopy (FT‐IR and RAMAN) [8, 12, 20, 21] and the morphological technique of atomic force microscopy (AFM) [22–24]. In the case of energy transfer processes of excited carriers, the analysis of polarization of emitted light can be directly correlated with polymeric chain position parallel to the direction of the transition dipole moment [25–27]. Moreover, we need to consider the effects of the deposition method.

In present work, we used the electrochemical synthesis for deposition of P3AT thin solid film correlating with optical properties. We demonstrated an easy and efficient alternative method

as a catalyst.

ditionally been prepared by chemical synthesis from 3‐AT monomers with FeCl3

The P3DDT films were electropolymerized and deposited on FTO substrates by cycles of the voltammetry method (CV) using the IVIUM COMPACTSTAT potentiostat/galvanostat. Alternatively, reference [9] presented additional deposition techniques of chronoamperometry and chronocoulometry used to synthesize P3AT films [9]. Polymerization was accomplished by continuous cycling the potential of the FTO electrode between +2.200 and ‐0.000 V for Et4 NBF4 and +2.900 and ‐1.000 V for LiClO4 . Electropolymerization scan rate was determined at 0.050 V/s. Different sample thicknesses were obtained by increasing the number of cycles in CV ranging from 1 to 10 for each electrolyte. In the electropolymerization CV technique, we used standard‐three‐electrodes‐cells: platinum auxiliary electrode, reference electrode con‐ taining saturated demonized water solution of potassium chloride (KCl) and FTO‐working electrode. Electrodes were immersed in an electrolyte solution containing acetonitrile, mono‐ mer 3‐dodecylthiophene and inert salt under a controlled atmosphere using argon gas.


**Table 1.** Label of P3DDT film in function of the cycles number and electrolyte.

Initially, the voltammetry cyclic synthesis, oxidate or reducte the polymer monomers bond‐ ing its covalently increasing the main chain length and deposit the polymer via physical‐ chemistry interaction with the working electrode. The polymer chain length formed in this process occurs until the saturation limit of the chain is achieved. Then, in each CV cycle new polymer chains are deposited on the previously deposited polymer layer. **Figure 1** shows the cyclic voltammogram for EtNBF10 (**Figure 1a**) and LiClO10 (**Figure 1b**) films. We observe that between the first and the last cycles, there are different maximum of reduction poten‐ tials 0.56 and 0.32 V for EtNBF10 and 0.36 and 0.48 V for LiClO10, respectively. These ddp differences indicate the film thickness and material deposited amount is increasing between consecutive CV cycles [33].

The increase in number of electropolymerization cycles leads to the formation of various lay‐ ers until the saturation point of the film. Assuming that in each cycle, the maximum voltage of oxidation or reduction represents the cathode and anodic ionizing potential, respectively, the difference in ionization potentials allows us to infer the *energy gap* (*E*<sup>g</sup> ) of the material. It is important to observe that the CV curves need to display only one oxidation or reduction process, as in the case of P3DDT films, see **Figure 1**. This method was first introduced by Eckhardt et al. [34] for organic semiconductor materials. **Table 2** shows the oxidation and reduction maximum values for all P3DDT films. Calculated *E*<sup>g</sup> values increase from the first to the last cycle regardless of the electrolyte since the resistance of the carriers is greater for the additional P3DDT layer after each cycle, thus increasing film thickness [30, 33].

By comparing the values, shown in **Table 2** for the *energy gap*, it can be observed that the films with the Et4 NBF4 electrolyte require less energy for their formation. The increase in voltage required for the formation of films and of energy is due to the presence of BF<sup>4</sup> ‐ and ClO4 ‐ anions in the electropolymerization, which has great influence on the morphology, structure and electrochemical polymer properties [7–9, 19, 20, 31]. In addition, great *E*<sup>g</sup> can be correlated with the high doping of P3DDT chains producing higher polymer quinone form. It was showed by Therézio et al. [8, 19] studied that the energy gap in doped P3AT films increases. Therefore, in our study, we also recommend the P3DDT films with lesser deposition cycles and lesser doped or lower *gap energy* which displays major polymer chains in the pristine form.

**Figure 1.** Cyclic voltammetry of (a) EtNBF10 and (b) LiClO10 films.

Initially, the voltammetry cyclic synthesis, oxidate or reducte the polymer monomers bond‐ ing its covalently increasing the main chain length and deposit the polymer via physical‐ chemistry interaction with the working electrode. The polymer chain length formed in this process occurs until the saturation limit of the chain is achieved. Then, in each CV cycle new polymer chains are deposited on the previously deposited polymer layer. **Figure 1** shows the cyclic voltammogram for EtNBF10 (**Figure 1a**) and LiClO10 (**Figure 1b**) films. We observe that between the first and the last cycles, there are different maximum of reduction poten‐ tials 0.56 and 0.32 V for EtNBF10 and 0.36 and 0.48 V for LiClO10, respectively. These ddp differences indicate the film thickness and material deposited amount is increasing between

**Electrolyte Number of cycles Nomenclature used**

2 EtNBF02 4 EtNBF04 6 EtNBF06 8 EtNBF08 10 EtNBF10

2 LiClO02 4 LiClO04 6 LiClO06 8 LiClO08 10 LiClO10

NBF4 1 EtNBF01

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

LiClO4 1 LiClO01

**Table 1.** Label of P3DDT film in function of the cycles number and electrolyte.

The increase in number of electropolymerization cycles leads to the formation of various lay‐ ers until the saturation point of the film. Assuming that in each cycle, the maximum voltage of oxidation or reduction represents the cathode and anodic ionizing potential, respectively,

is important to observe that the CV curves need to display only one oxidation or reduction process, as in the case of P3DDT films, see **Figure 1**. This method was first introduced by Eckhardt et al. [34] for organic semiconductor materials. **Table 2** shows the oxidation and reduction maximum values for all P3DDT films. Calculated *E*<sup>g</sup> values increase from the first to the last cycle regardless of the electrolyte since the resistance of the carriers is greater for the

By comparing the values, shown in **Table 2** for the *energy gap*, it can be observed that the films

in the electropolymerization, which has great influence on the morphology, structure and

electrolyte require less energy for their formation. The increase in voltage

) of the material. It

‐

and ClO4

‐ anions

the difference in ionization potentials allows us to infer the *energy gap* (*E*<sup>g</sup>

additional P3DDT layer after each cycle, thus increasing film thickness [30, 33].

required for the formation of films and of energy is due to the presence of BF<sup>4</sup>

consecutive CV cycles [33].

Et4

with the Et4

NBF4


**Table 2.** Cathodic and anodic ionization potential for P3DDT films.

### **3. Optical characterizations**

#### **3.1. UV‐Vis absorption1**

In the UV‐Vis spectral range with maximum absorption centered at ~450 nm. For the EtNBF02 film, a well‐resolved absorbance band at ~775 nm is also observed. This band is the result of the interaction between the BF4 ‐ anion and P3DDT polymer chains [8, 9, 20]. Similar results are observed in different P3AT polymers [35–37]. To higher electropolymerization cycle number, the band at ~775 nm is further evident (not shown), in which it is possible to correlate the cycle number and UV‐Vis absorption intensity to follow the polymer‐grown deposition. **Figure 2b** shows the absorption spectra for the LiClO01 and LiClO02 films in the UV‐Vis spectral range. It observes in **Figure 2b** that the maximum absorption is approximated at ~400 nm for the LiClO02 film, but it cannot be confirmed by the exact spectral position because of the absorp‐ tion of the FTO substrate. The blue shift of the absorption maximum position in comparison of absorbance spectra of the Et4 NBF4 films should be considered due to the presence of two different P3DDT molecules morphologies or the diminish of the length of the polymer chains. As a result, the decrease of the conjugation and the increase of the gap energy of the mate‐ rial occur. That effect was recently reported by Therézio et al. [8, 12] to the P3AT derivative where the maximum position and intensity of absorption change in function of the electrolyte. Other castellation is the exposition of the film on the atmosphere environment that induces relatively quick (days) polymer films degradation. Lower intensity and poorly resolved absor‐ bance band at ~650 nm is observed in **Figure 2b** for LiClO02 and thick films (not shown) due to the interaction between the ClO4 ‐ anions and the P3DDT molecules [8, 9, 12, 20, 31].

**Figure 2.** UV‐Vis spectra for P3DDT films with one and two electropolymerization cycles: (a) EtNBF01 and EtNBF02 films and (b) LiClO01 and LiClO02 films.

1 UV‐Vis measurements were conducted using a spectrophotometer FEMTO XI 800, operating in the 190‐900 nm range.

#### **3.2. Photoluminescence (PL)2**

**3. Optical characterizations**

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

the interaction between the BF4

of absorbance spectra of the Et4

the interaction between the ClO4

films and (b) LiClO01 and LiClO02 films.

1

‐

NBF4

‐

In the UV‐Vis spectral range with maximum absorption centered at ~450 nm. For the EtNBF02 film, a well‐resolved absorbance band at ~775 nm is also observed. This band is the result of

observed in different P3AT polymers [35–37]. To higher electropolymerization cycle number, the band at ~775 nm is further evident (not shown), in which it is possible to correlate the cycle number and UV‐Vis absorption intensity to follow the polymer‐grown deposition. **Figure 2b** shows the absorption spectra for the LiClO01 and LiClO02 films in the UV‐Vis spectral range. It observes in **Figure 2b** that the maximum absorption is approximated at ~400 nm for the LiClO02 film, but it cannot be confirmed by the exact spectral position because of the absorp‐ tion of the FTO substrate. The blue shift of the absorption maximum position in comparison

different P3DDT molecules morphologies or the diminish of the length of the polymer chains. As a result, the decrease of the conjugation and the increase of the gap energy of the mate‐ rial occur. That effect was recently reported by Therézio et al. [8, 12] to the P3AT derivative where the maximum position and intensity of absorption change in function of the electrolyte. Other castellation is the exposition of the film on the atmosphere environment that induces relatively quick (days) polymer films degradation. Lower intensity and poorly resolved absor‐ bance band at ~650 nm is observed in **Figure 2b** for LiClO02 and thick films (not shown) due to

UV‐Vis measurements were conducted using a spectrophotometer FEMTO XI 800, operating in the 190‐900 nm range.

**Figure 2.** UV‐Vis spectra for P3DDT films with one and two electropolymerization cycles: (a) EtNBF01 and EtNBF02

anion and P3DDT polymer chains [8, 9, 20]. Similar results are

films should be considered due to the presence of two

anions and the P3DDT molecules [8, 9, 12, 20, 31].

**3.1. UV‐Vis absorption1**

We consider several radiative contributions associated with different interactions to simulate the emission spectra due to the presence of the anion in the electrolyte solution and polymer chain [8, 9, 12, 31]. Basically, for P3ATs emission spectra we may approximate the line shape considering quinone or oligomer structures (high energy) and pristine (low energy) structures. The main structures present in these polymers such as the quinone and pristine structures are shown in **Figure 3**. The result is the maximum shift or relative intensity change due to the interaction of electron‐vibrational modes of quinone or pristine structures [8, 19] or different lengths of polymer chains [38].

**Figure 3.** Scheme for P3DDT (a) pristine and (b) quinone structure.

**Figure 4a** and **b** shows the PL spectra of P3DDT film synthesized using both Et<sup>4</sup> NBF4 or LiClO4 electrolytes, respectively. Note that the spectra are broad in the UV‐Vis spectral range. **Figure 4a** shows normalized PL spectra for P3DDT films grown using the Et<sup>4</sup> NBF4 electrolyte, where the maximum of the emission redshift increased the number of cycles. This is due to the presence of oligomers with smaller conjugation lengths [38] and quinone chains [8, 19]. In addition for EtNBF01–EtNBF04 films, it is possible the occurrence of polymer chains with lower molecular weight and interface substrate/polymer, polymer/polymer and polymer/ electrolyte effects [14, 22]. By increasing the film thickness (>4 cycles), PL spectra shifted to high wavelengths (*redshift*), it according the rise of in residence time in the electropolymeriza‐ tion process in the presence of higher conjugated polymer chains and pristine structures [8, 9, 12]. **Figure 4b** shows the PL spectra of P3DDT films synthesized with LiClO<sup>4</sup> , presenting the similar line shape characteristics observed in **Figure 4(a)**. Film emission spectra utilizing that electrolyte also display different radiative processes. In this case, we consider two different configurations of P3DDT molecules, i.e., pristine or quinone structures [8, 9, 12]. The emission line shape is practically identical to LiClO01–LiClO08 films. However, the PL spectra for the thicker LiClO10 film is red shifted due to the new polymer structures created by the interac‐ tion ClO4 ‐ ion present in the solution [8, 12]. It is possible in the synthesis the presence of high

<sup>2</sup> PL measurements were obtained by exciting the samples with the 405 nm line of a diode laser at 4.0 mW (Laser Line‐ iZi), vertically polarized in relation to the laboratory reference. The emission was detected and analyzed by a USB 2000 ocean optics spectrophotometer.

polymer conjugation length or pristine structures [8, 9, 12]. Results are coherent with recent observation of P3ATs when different electrolytes tend to influence the molecular polymer structure [7, 33].

**Figure 4.** PL spectra for P3DDT films formed with electrolyte (a) Et4NBF4 and (b) LiClO4. PL simulation for (c) EtNBF10 and (d) LiClO10 films.

**Figure 4(c)** and **(d)** shows the spectra simulation for EtNBF10 and LiClO10 films, respectively, using the multi‐Gaussian function, following the procedure utilized for P3AT polymers [8, 9, 12, 13]. First, we consider the polymer chains without the presence of dopants and without structural changes. Second, we introduce the contribution to the emission line shape of dop‐ ants and possible new structural changes (**Figure 3**) [8], as established by the P3AT family [35– 37] in the energy range of polymeric chains or oligomers [38, 39]. Moreover, it was also added bands due to the interaction of the electron‐vibrational modes. The typical optically active vibration mode is 1450 cm‐1 for the C=C group [37]. In **Figure 4(c)**, the PL spectra centered at ~516 nm can be assigned to high conjugation polymer chains. The second band at ~584 nm is due to quinone structures, which is the result of interaction between the salt and the polymer chains. The band at ~638 nm is attributed to pristine without the salt‐polymer interaction. Finally, the last one at ~678 nm is normally reported as the vibrational replica [35–37]. We perform a similar procedure to simulate the spectra of P3DDT films processed with LiClO<sup>4</sup> and the results are presented in **Figure 4(d)**. Note in this case, the spectrum may be simulated with only Gaussians curves [8, 12]. The emission band at ~572 nm is assigned to the interac‐ tion between the electrolyte and the polymer chain forming mainly quinone structures and a band at ~610 nm is due to the pristine structure. That result is agreement with the reports on the P3ATs family [8, 9, 12, 19]. A similar simulation was performed for all samples analyzed and the results are shown in **Table 3**. And it observed that the films synthesized with Et<sup>4</sup> NBF4 have vibrational replicas at 680 nm. The presence of the bands assigned to the quinone and pristine structures is also evident for both electrolytes. As a result, it is possible to correlate the maximum emission spectral position to the amount of these structures in the polymeric film.

polymer conjugation length or pristine structures [8, 9, 12]. Results are coherent with recent observation of P3ATs when different electrolytes tend to influence the molecular polymer

**Figure 4(c)** and **(d)** shows the spectra simulation for EtNBF10 and LiClO10 films, respectively, using the multi‐Gaussian function, following the procedure utilized for P3AT polymers [8, 9, 12, 13]. First, we consider the polymer chains without the presence of dopants and without structural changes. Second, we introduce the contribution to the emission line shape of dop‐ ants and possible new structural changes (**Figure 3**) [8], as established by the P3AT family [35– 37] in the energy range of polymeric chains or oligomers [38, 39]. Moreover, it was also added bands due to the interaction of the electron‐vibrational modes. The typical optically active

**Figure 4.** PL spectra for P3DDT films formed with electrolyte (a) Et4NBF4 and (b) LiClO4. PL simulation for (c) EtNBF10

~516 nm can be assigned to high conjugation polymer chains. The second band at ~584 nm is due to quinone structures, which is the result of interaction between the salt and the polymer chains. The band at ~638 nm is attributed to pristine without the salt‐polymer interaction.

for the C=C group [37]. In **Figure 4(c)**, the PL spectra centered at

structure [7, 33].

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

vibration mode is 1450 cm‐1

and (d) LiClO10 films.

**Table 2** shows that the P3DDT *band gap* should change, increasing the number of the deposi‐ tion cycles. However, data in **Table 3** demonstrate that the emission of quinone or pristine chains does not change the spectral position, respectively, at 569 ± 7 and 631 ± 8 nm for films grown in Et4 NBF4 and 538 ± 9 and 588 ± 7 nm for the films grown in LiClO<sup>4</sup> . The redshift in the emission spectra is due to the emission bands of the pristine or quinone species. In addition, the present result shows how the use of electropolymerization is able to synthesize regular polymer chains, in which it is an important point when the reproducibility of polymeric lay‐ ers is important, mainly to applied in an organic device area.


**Table 3.** Maximum of curves of the deconvolution in the PL spectra for EtNBFxx and LiClOxx films.

#### **3.3. Emission ellipsometry (EE)3**

Recently, we demonstrated the correlation between polarized emission light and P3DDT films grown using the electrochemical using emission ellipsometry technique [25, 26, 40]. Photoluminescence polarization reveal important properties of the material structure, e.g., anisotropy, which has immediate application in industry [41]. By using the ellipsometry tech‐ nique, the polarization state of the emitted light can be determined by calculating the *S*<sup>0</sup> *, S*<sup>1</sup> *, S*2 and *S*<sup>3</sup> Stokes parameters. *S*<sup>0</sup> is associated with the total light emitted amount, *S*<sup>1</sup> describes the linearly polarized amount of light in the vertical or horizontal direction, *S*<sup>2</sup> describes the linear polarization amount rotated by +45° or ‐45° and *S*<sup>3</sup> describes the circularly polarized light to the right or left. These parameters are obtained by adjusting the intensity *I* of the equa‐ tion [25, 40]:

$$I(\theta) = \frac{1}{2} [A + B \cdot \sin(2\theta) + C \cdot \cos(4\theta) + D \cdot \sin(4\theta)] \tag{2}$$

where *I* is the electric field intensity, *θ* is the angle between the axes of the quarter‐wave plate and of the polarizer, *<sup>A</sup>* <sup>=</sup> *<sup>S</sup>*<sup>0</sup> <sup>−</sup> *S* \_\_1 <sup>2</sup> , *<sup>B</sup>* <sup>=</sup> *<sup>S</sup>*<sup>3</sup> , *C* = − *S* \_\_1 <sup>2</sup> and *<sup>D</sup>* <sup>=</sup> <sup>−</sup> *S* \_\_2 <sup>2</sup> , where *S*<sup>0</sup> *, S*<sup>1</sup> *, S*<sup>2</sup> and *S*<sup>3</sup> are the Stokes parameters. In practice, the quarter‐wave plate is rotated by discrete angles *θ<sup>j</sup>* such that:

$$\begin{aligned} A &= \frac{2}{N} \sum\_{\nu=1}^{N} I\left(n \,\,\Theta\_{\rangle}\right) \\ B &= \frac{4}{N} \sum\_{n=1}^{N} I\left(n \,\,\Theta\_{\rangle}\right) \sin\left(2n \,\,\Theta\_{\rangle}\right) \\ C &= \frac{4}{N} \sum\_{n=1}^{N} I\left(n \,\,\Theta\_{\rangle}\right) \cos\left(4n \,\,\Theta\_{\rangle}\right) \\ D &= \frac{4}{N} \sum\_{n=1}^{N} I\left(n \,\,\Theta\_{\rangle}\right) \sin\left(4n \,\,\Theta\_{\rangle}\right) \end{aligned} \tag{3}$$

where *N* is the number of steps of the quarter‐wave plate. Eq. [2] can be solved considering the eight possible combinations of the harmonic functions (sine and cosine) and the total intensity (parameter *S*<sup>0</sup> ). In other words, the minimum number of points for solving Eq. (2) is *<sup>N</sup>* <sup>=</sup> 9 or Δ*<sup>θ</sup>* <sup>=</sup> 40° and, from the experimental point of view, the symmetry *I*(*θ*) <sup>=</sup> *<sup>I</sup>*(*<sup>θ</sup>* <sup>+</sup> <sup>2</sup>*<sup>π</sup>*). The new method to solve Eq. (2) was introduced by Basílio [42]. Stokes parameters are associated with the degree of polarization (*P*) of the emitted light by [25, 40]:.

$$P = \frac{\left(S\_1^2 + S\_2^2 + S\_3^2\right)^{\frac{1}{2}}}{S\_0}.\tag{4}$$

Moreover, it is also possible to obtain the dissymmetry factor *g* the circularly polarized light emission degree and the anisotropy factor *r*, it is associated with molecular ordering of the

<sup>3</sup> The emission ellipsometry experiment was performed using the setup described by Alliprandini et al. [25–27]. The samples were excited by a laser in 405 nm, and the emitted light was collected by a set of lenses and directed through an achromatic quarter‐wave‐plate (Newport 10RP54‐1), as a compensator and an achromatic polarizer (Newport 10LP‐ VIS‐B). The emission was detected and analyzed by an USB 2000 Ocean Optics spectrophotometer. The experiment was performed by rotating the compensator in its own plane from 0 rad (0°) to ~6.28 rad (360°), with steps of ~0.17 rad (10°). All measurements were performed at room temperature (~20°C) and under 10‐4 Torr vacuum.

polymer chains [22, 43]. These factors are obtained from the Stokes parameters in Eqs. (5) and (6). The use of the dissymmetry factor *g* equation is conditioned to the referential adopted: the vertical direction (*y*‐axis) as a positive sign and the horizontal direction (*x*‐axis) as a negative sign.

$$\mathcal{g} = \begin{array}{c} \pm 2 \frac{S\_3}{S\_0} \end{array} \tag{5}$$

$$r = \frac{-2\frac{S\_1}{S\_0}}{3 + \frac{S\_1}{S\_0}}\tag{6}$$

#### *3.3.1. Emission ellipsometry in P3DDT*

**3.3. Emission ellipsometry (EE)3**

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

Stokes parameters. *S*<sup>0</sup>

*I*(*θ*) = \_\_1

plate and of the polarizer, *<sup>A</sup>* <sup>=</sup> *<sup>S</sup>*<sup>0</sup> <sup>−</sup>

intensity (parameter *S*<sup>0</sup>

linear polarization amount rotated by +45° or ‐45° and *S*<sup>3</sup>

*S*2 and *S*<sup>3</sup>

that:

3

tion [25, 40]:

Recently, we demonstrated the correlation between polarized emission light and P3DDT films grown using the electrochemical using emission ellipsometry technique [25, 26, 40]. Photoluminescence polarization reveal important properties of the material structure, e.g., anisotropy, which has immediate application in industry [41]. By using the ellipsometry tech‐ nique, the polarization state of the emitted light can be determined by calculating the *S*<sup>0</sup>

light to the right or left. These parameters are obtained by adjusting the intensity *I* of the equa‐

where *I* is the electric field intensity, *θ* is the angle between the axes of the quarter‐wave

, *C* = − *S* \_\_1

Stokes parameters. In practice, the quarter‐wave plate is rotated by discrete angles *θ<sup>j</sup>*

*A* = \_\_2 *<sup>N</sup>* ∑ *n*=1 *N*

the linearly polarized amount of light in the vertical or horizontal direction, *S*<sup>2</sup>

*S* \_\_1 <sup>2</sup> , *<sup>B</sup>* <sup>=</sup> *<sup>S</sup>*<sup>3</sup>

*B* = \_\_4 *<sup>N</sup>* ∑ *n*=1 *N*

*C* = \_\_4 *<sup>N</sup>* ∑ *n*=1 *N*

*D* = \_\_4 *<sup>N</sup>* ∑ *n*=1 *N*

with the degree of polarization (*P*) of the emitted light by [25, 40]:.

All measurements were performed at room temperature (~20°C) and under 10‐4 Torr vacuum.

*<sup>P</sup>* <sup>=</sup> (*S*<sup>1</sup>

is associated with the total light emitted amount, *S*<sup>1</sup>

<sup>2</sup> and *<sup>D</sup>* <sup>=</sup> <sup>−</sup>

). In other words, the minimum number of points for solving Eq. (2) is

*I*(*n θj*)

*I*(*n θj*)sin(2*n θj*)

*I*(*n θj*)cos(4*n θj*)

*I*(*n θj*)sin(4*n θj*)

where *N* is the number of steps of the quarter‐wave plate. Eq. [2] can be solved considering the eight possible combinations of the harmonic functions (sine and cosine) and the total

*<sup>N</sup>* <sup>=</sup> 9 or Δ*<sup>θ</sup>* <sup>=</sup> 40° and, from the experimental point of view, the symmetry *I*(*θ*) <sup>=</sup> *<sup>I</sup>*(*<sup>θ</sup>* <sup>+</sup> <sup>2</sup>*<sup>π</sup>*). The new method to solve Eq. (2) was introduced by Basílio [42]. Stokes parameters are associated

> <sup>2</sup> + *S*<sup>2</sup> <sup>2</sup> + *S*<sup>3</sup> 2 ) \_\_1 2

Moreover, it is also possible to obtain the dissymmetry factor *g* the circularly polarized light emission degree and the anisotropy factor *r*, it is associated with molecular ordering of the

The emission ellipsometry experiment was performed using the setup described by Alliprandini et al. [25–27]. The samples were excited by a laser in 405 nm, and the emitted light was collected by a set of lenses and directed through an achromatic quarter‐wave‐plate (Newport 10RP54‐1), as a compensator and an achromatic polarizer (Newport 10LP‐ VIS‐B). The emission was detected and analyzed by an USB 2000 Ocean Optics spectrophotometer. The experiment was performed by rotating the compensator in its own plane from 0 rad (0°) to ~6.28 rad (360°), with steps of ~0.17 rad (10°).

\_\_\_\_\_\_\_\_\_ *S*0

<sup>2</sup>[*<sup>A</sup>* <sup>+</sup> *<sup>B</sup>* . sin(2*θ*) <sup>+</sup> *<sup>C</sup>* . cos(4*θ*) <sup>+</sup> *<sup>D</sup>* . sin(4*<sup>θ</sup>* )] (2)

*S* \_\_2

<sup>2</sup> , where *S*<sup>0</sup>

*, S*<sup>1</sup> *,* 

describes

are the

such

(3)

describes the

describes the circularly polarized

*, S*<sup>1</sup> *, S*<sup>2</sup>

. (4)

and *S*<sup>3</sup>

The results presented in this section are dedicated to study the photophysical effects of P3AT films. In addition, it is possible to correlate the energy transfer mechanisms [44, 45] and polar‐ ization light states [22, 25–27]. Note that, in general, they are intrinsic characteristics semicon‐ ductor polymers. Stokes parameters are directly related to the light polarization states and, consequently, with factors related to orientation of the polymer chains in the films [25, 26], in which it provides information about the samples molecular ordering along the polymer films. In principle, the electrochemically synthesis by cyclic voltammetry does not show a molecular order [9].

**Figure 5** shows the parameters *S*<sup>1</sup> */S*0 , *S*<sup>2</sup> */S*0 and *S*<sup>3</sup> */S*0 in the spectral range of P3DDT emission obtained from the EE data for EtNBF01 and EtNBF10 films. The Stokes parameter values are virtually null in **Figure 5a**. It indicates that the light emitted by EtNBF01 films has random polarizing directions, i.e., depolarized with high probability of energy transfer from the pho‐ toexcited carriers in all directions of the polymeric film plane. It is important to remember that the excitation polarization is linear in the vertical direction (laboratory referential) and only chronophers with transition dipole in the parallel direction are excited. However, for EtNBF10 film thickness, we observe, in **Figure 4b** significant variation for the *S*<sup>1</sup> /*S*0 and *S*<sup>2</sup> /*S*0 parameters. In principle, the films grown using electrochemical techniques have not molecu‐ lar order [40], but they may have partially polarized emission when the excitation light is linearly polarized [9, 24, 40, 43]. Another important observation is that the inversion signal of the *S*<sup>1</sup> /*S*0 curve occurs simultaneously with the reversal of *S*<sup>2</sup> /*S*0 curve. This may be explained, according to Foster's energy transfer process mechanism [44] by the excitation low molecular weight chains or oligomers that have their transition electric dipole or part thereof aligned with the excitation source, energy absorbing (405 nm). Thus, part of the absorbed light is transferred via Förster processes to another polymer with larger conjugation length, pristine and quinone structures, in a random direction, depolarized and decreasing the values of the *S*1 /*S*0 factor close to zero above 525 nm. However, some excited oligomers and low molecular weight chains may not energy transfer increasing the emission in the parallel direction of exci‐ tation polarized light [9, 26]. It is observed in the *S*<sup>1</sup> /*S*0 factor from 475 to 525 nm spectral range due to the effects of the substrate‐polymer or polymer‐polymer interface [14, 22]. Finally, the *S*3 /*S*0 parameter indicates that the light emitted does not have significant right or left circular polarization.

**Figure 5.** EE curves for films (a) EtNBF01 and (b) EtNBF10.

Förster energy transfer occurs when two conditions are met [44, 45]. First (i), the distance between the donor and the acceptor chromophores is up to ~ 10 nm, and (ii) second, paral‐ lelism condition, i. e, the electric dipole moment of the donor and the acceptor is parallel (aligned) or component of electric dipole moment of the acceptor is in the parallel direction of electric dipole moment of the donor [44]. Thus, when conjugated polymer segments absorb energy partially or totally is transferred to another, generally, higher conjugation polymer degree due to the electronic‐vibrational relaxation mechanism. During the transfer process, the energy undergoes depolarization or rotational changes due to the misalignment of the dipoles in the same direction of the polymer main chains. As a result, when there is an emis‐ sion in another part of the polymer or different polymer chains, a different polarization com‐ pared to the initial direction of the excitation polarization is obtained. Thus, the molecules, whose dipole or parts of its components are aligned with the excitation source, transfer energy to the polymers with greater conjugation lengths, causing an emission at lower energies. For the thick film in **Figure 5b**, the oligomers show great probability to transfer energy to other chains with greater conjugation lengths (quinone and pristine chains), parameter *S*<sup>1</sup> /*S*0 ~0, see spectral range assigned to oligomers between 475 and 525 nm. Similarly, **Figure 6** shows the EE curves for LiClO01 (**Figure 6a**) and LiClO10 (**Figure 6b**) films. Note that the emission of these films has horizontal and linear polarization, *S*<sup>1</sup> /*S*0 > 0 and polarization rotation to + 45° in relation to the polarization of the excitation light *S*<sup>2</sup> /*S*0 < 0. Circular polarization emission is not observed at *S*<sup>3</sup> /*S*0 ~0 to 500–675 nm. All samples processed using LiClO4 have the same EE curves characteristics (**Figure 6**). An explanation for the absence of reverse bias directions and rotation observed for the films processed using the Et<sup>4</sup> NBF4 electrolyte is the preferential formation of polymer chains with smaller conjugation lengths. It is compatible with the lower energy transfer at a higher spectral range <550 nm.

**Figure 6.** EE curves for films (a) LiClO01 and (b) LiClO10.

Förster energy transfer occurs when two conditions are met [44, 45]. First (i), the distance between the donor and the acceptor chromophores is up to ~ 10 nm, and (ii) second, paral‐ lelism condition, i. e, the electric dipole moment of the donor and the acceptor is parallel (aligned) or component of electric dipole moment of the acceptor is in the parallel direction of electric dipole moment of the donor [44]. Thus, when conjugated polymer segments absorb energy partially or totally is transferred to another, generally, higher conjugation polymer degree due to the electronic‐vibrational relaxation mechanism. During the transfer process, the energy undergoes depolarization or rotational changes due to the misalignment of the dipoles in the same direction of the polymer main chains. As a result, when there is an emis‐ sion in another part of the polymer or different polymer chains, a different polarization com‐ pared to the initial direction of the excitation polarization is obtained. Thus, the molecules, whose dipole or parts of its components are aligned with the excitation source, transfer energy to the polymers with greater conjugation lengths, causing an emission at lower energies. For the thick film in **Figure 5b**, the oligomers show great probability to transfer energy to other

chains with greater conjugation lengths (quinone and pristine chains), parameter *S*<sup>1</sup>

these films has horizontal and linear polarization, *S*<sup>1</sup>

/*S*0

**Figure 5.** EE curves for films (a) EtNBF01 and (b) EtNBF10.

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

is not observed at *S*<sup>3</sup>

in relation to the polarization of the excitation light *S*<sup>2</sup>

and rotation observed for the films processed using the Et<sup>4</sup>

spectral range assigned to oligomers between 475 and 525 nm. Similarly, **Figure 6** shows the EE curves for LiClO01 (**Figure 6a**) and LiClO10 (**Figure 6b**) films. Note that the emission of

EE curves characteristics (**Figure 6**). An explanation for the absence of reverse bias directions

/*S*0

~0 to 500–675 nm. All samples processed using LiClO4

/*S*0

NBF4

/*S*0

have the same

> 0 and polarization rotation to + 45°

< 0. Circular polarization emission

electrolyte is the preferential

~0, see

The polarization degree *P* (Eq. (4)) indicates the amount of light that is polarized, without dif‐ ferentiating between linearly or circularly polarized lights [25, 27, 40]. There are other two important parameter to quantify the polarization: the anisotropy factor *r* (Eq. (6)) that may be correlated with the direction of the polymer chains and the dissymmetry factor *g* (Eq. (5)) asso‐ ciated with the emission of the circularly polarized light [22, 40, 43]. The parameters *P*, *r* and *g* are obtained directly from Stokes parameter values [22, 43]. **Figures 5a** and **7a** show the polar‐ ization degree parameters for the EtNBF10 film in the function of the emission wavelength, presenting maximum at 495 and 630 nm. At 495 nm, we observe the signal inversion of *S*<sup>2</sup> /*S*0 coinciding with the minimal value for *S*<sup>1</sup> /*S*0 curves. The dissymmetry factor *g* in the emission spectral region is around ~5% or below, indicating that there is no emission of circularly polar‐ ized light [44]. **Figure 5b** shows the polarization degree *P*, anisotropy factors *r* and dissymme‐ try factor *g* for the LiClO10 film. We can observe that the polarization of the emitted light is high (~28%) at a low wavelength region (<550 nm) addressed with oligomers decreasing monotonic above 550 nm. The parameters *r* and *g* followed the spectral dependence of *p*. However, the parameter *g* does not display significant values. On the other hand, *r* values in the spectral range of oligomers emission are related to the decrease of energy transfer between adjacent polymer chain, according to reference [26]. The dissymmetry factor *g* has no significant intensity.

**Figure 7.** Polarization degree *P*, anisotropy factor *r* and dissymmetry factor *g* obtained from the Stokes parameters for the (a) EtNBF10 film and (b) LiClO10 film.

### **4. Conclusions**

We showed the strong correlation between the optical properties of a P3AT polymer and the con‐ ditions of the polymer electrochemical growth, demonstrating that it is possible to control the optical properties of polymeric films by controlling the growth conditions. For this, we use dif‐ ferent electrolytes during the synthesis of polymeric films, which were linked to the number of growing cycles, according to the cyclic voltammetry electrochemical technique. Electrochemical synthesis is shown in efficient growth polymeric films when the concentrations of the reagents are equilibrated, e.g., the concentration for electropolymerization of the P3DDT occurs homoge‐ neously with 0.050 mol L‐1 monomer and 0.100 mol L‐1 of LiClO4 . Furthermore, it is possible to measure the energy gap *E*<sup>g</sup> for organic semiconductor directly from the cyclic voltammogram.

Through the UV‐Vis results, it is possible to conclude that the electrolytes reacted differently in this material, shifting the spectrum to other regions of absorption as occurs to electrolyte change. PL results showed the existence of several contributions in each spectrum, in which the highest intensity contributions, quinone and pristine are the result of two structures that the polymer chains can take through the interactions between each electrolyte/polymer. These contributions are able to shift the maximum emission in each spectrum when the film thickness increases, and more effectively for the films containing Et<sup>4</sup> NBF4 . There are also contributions resulting from the oligomers emission and the electron‐phonon interactions. The EE demon‐ strated energy transfer processes by the Förster mechanism, where the emission polarization is observed and this has gradually changed with increasing emission wavelength. However, it is an isotropic material when obtained for CV, shown by the anisotropy factor, *r*. These analyses show that the Förster *energy transfer process* occurs in this material and is responsible for the emission throughout the spectral window. Furthermore, more accentuated emissions polarization can be related to the oligomer emission.

Results show that the P3ATs deposited electrochemically has great potential for application in optoelectronic organic devices since P3ATs optical properties can be easily adjusted by con‐ trolling the deposition. In addition, results also showed the great light absorption capacity and a broad spectral window for the P3DDT emission. Moreover, it is also observed the presence of electron‐phonon combination, which can contribute to the occurrence of energy transfer or charge transfer significantly, enhancing the use of P3DDT in optoelectronic devices, which makes this promising material to form the active layer of multiple devices, such as organic light emitting diodes (OLEDs), photovoltaics, photodetectors and mobile devices screens (displays), among others.

## **Acknowledgements**

The authors are grateful to the following Brazilian Agencies: FAPEMIG, CAPES, CNPQ and FUFMT.

### **Author details**

**4. Conclusions**

the (a) EtNBF10 film and (b) LiClO10 film.

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

neously with 0.050 mol L‐1

measure the energy gap *E*<sup>g</sup>

We showed the strong correlation between the optical properties of a P3AT polymer and the con‐ ditions of the polymer electrochemical growth, demonstrating that it is possible to control the optical properties of polymeric films by controlling the growth conditions. For this, we use dif‐ ferent electrolytes during the synthesis of polymeric films, which were linked to the number of growing cycles, according to the cyclic voltammetry electrochemical technique. Electrochemical synthesis is shown in efficient growth polymeric films when the concentrations of the reagents are equilibrated, e.g., the concentration for electropolymerization of the P3DDT occurs homoge‐

**Figure 7.** Polarization degree *P*, anisotropy factor *r* and dissymmetry factor *g* obtained from the Stokes parameters for

Through the UV‐Vis results, it is possible to conclude that the electrolytes reacted differently in this material, shifting the spectrum to other regions of absorption as occurs to electrolyte change. PL results showed the existence of several contributions in each spectrum, in which the highest intensity contributions, quinone and pristine are the result of two structures that the polymer chains can take through the interactions between each electrolyte/polymer. These contributions are able to shift the maximum emission in each spectrum when the film thickness

resulting from the oligomers emission and the electron‐phonon interactions. The EE demon‐ strated energy transfer processes by the Förster mechanism, where the emission polarization is observed and this has gradually changed with increasing emission wavelength. However, it is an isotropic material when obtained for CV, shown by the anisotropy factor, *r*. These

of LiClO4

for organic semiconductor directly from the cyclic voltammogram.

NBF4

. Furthermore, it is possible to

. There are also contributions

monomer and 0.100 mol L‐1

increases, and more effectively for the films containing Et<sup>4</sup>

Sankler Soares de Sá1 , Fernando Costa Basílio2 , Henrique de Santana3 , Alexandre Marletta<sup>2</sup> and Eralci Moreira Therézio4 \*


### **References**


[15] Marletta A. Optical properties of organic semiconductors based on Light Emitting Polymers [thesis]. São Carlos: Universidade de São Paulo; 2001. DOI:10.11606/T.76.2001. tde‐10022002‐091803

[3] Valaski R, Moreira LM, Micaroni L, Hümmelgen IA. The electronic behavior of poly(3‐ octylthiophene) electrochemically synthesized onto Au substrate. Brazilian Journal of

[4] Wang G, Yuan C, Lu Z, Wei Y. Enhancement of organic electroluminescent intensity by charge transfer from guest to host. Journal of Luminescence. 1996;**68**(1):49–54. DOI:

[5] Yoshino K, Manda Y, Sawada K, Onoda M, Sugimoto Ri. Anomalous dependences of luminescence of poly(3‐alkylthiophene) on temperature and alkyl chain length. Solid

[6] Ohmori Y, Uchida M, Muro K, Yoshino K. Visible‐light electroluminescent diodes uti‐ lizing poly(3‐alkylthiophene). Japanese Journal of Applied Physics. 1991;**30**(11B):L1938.

[7] Roncali J. Conjugated poly(thiophenes): synthesis, functionalization and applications.

[8] Therézio EM, Duarte JL, Laureto E, Di Mauro E, Dias IL, Marletta A, et al. Analysis of the optical properties of poly(3‐octylthiophene) partially dedoped. Journal of Physical

[9] Therézio EM, Franchello F, Dias IFL, Laureto E, Foschini M, Bottecchia OL, et al. Emission ellipsometry as a tool for optimizing the electrosynthesis of conjugated polymers thin

[10] Cervantes TNM, Bento DC, Maia ECR, Fernandes RV, Laureto E, Moore GJ, et al. The influence of different electrolytes on the electrical and optical properties of polymer films electrochemically synthesized from 3‐alkylthiophenes. Journal of Materials Science:

[11] Cervantes TNM, Bento DC, Maia ECR, Zaia DAM, Laureto E, da Silva MAT, et al. In situ and ex situ spectroscopic study of poly(3‐hexylthiophene) electrochemically synthe‐ sized. Journal of Materials Science: Materials in Electronics. 2012;**23**(10):1916–1921. DOI:

[12] Maia ECR, Bento DC, Laureto E, Zaia DAM, Therézio EM, Moore JG, et al. Spectroscopic analysis of the structure and stability of two electrochemically synthesized poly(3‐alkyl‐ thiophene)s. Journal of the Serbian Chemical Society. 2013;**78**(4):507–521. DOI: 10.2298/

[13] Yoshino K, Hayashi S, Sugimoto R. Preparation and properties of conducting het‐ erocyclic polymer films by chemical method. Japanese Journal of Applied Physics.

[14] Therézio EM, Piovesan E, Anni M, Silva RA, Oliveira ON, Marletta A. Substrate/semi‐ conductor interface effects on the emission efficiency of luminescent polymers. Journal

Materials in Electronics. 2014;**25**(4):1703–1715. DOI: 10.1007/s10854‐014‐1787‐4

State Communications. 1989;**69**(2):143–146. DOI: 10.1016/0038‐1098(89)90379‐7

Physics. 2003;**33**:392–397. DOI: 10.1590/S0103‐97332003000200043

Chemical Reviews. 1992;**92**(4):711–738. DOI: 10.1021/cr00012a009

Organic Chemistry. 2011;**24**(8):640–645. DOI: 10.1002/poc.1802

films. Thin Solid Films. 2013;**527**:255–260. DOI:10.1016/j.tsf.2012.11.093

10.1016/0022‐2313(95)00092‐5

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

DOI: 10.1143/JJAP.30.L1938

10.1007/s10854‐012‐0880‐9

1984;**23**(12A):L899. DOI: 10.1143/JJAP.23.L899

of Applied Physics. 2011;**110**(4):044504. DOI: 10.1063/1.3622143

JSC120327111R


[37] Kanemoto K, Sudo T, Akai I, Hashimoto H, Karasawa T, Aso Y, et al. Intrachain photolu‐ minescence properties of conjugated polymers as revealed by long oligothiophenes and polythiophenes diluted in an inactive solid matrix. Physical Review B. 2006;**73**(23):235203. DOI: 10.1103/PhysRevB.73.235203

[26] Alliprandini‐Filho P, da Silva RA, Barbosa Neto NM, Marletta A. Partially polarized fluorescence emitted by MEHPPV in solution. Chemical Physics Letters. 2009;**469**

[27] Alliprandini‐Filho P, Silva RA, Silva GB, Barbosa Neto NM, Cury LA, Moreira RL, et al. Measurement of the emitted light polarization state in oriented and non‐ori‐ ented PPV films. Macromolecular Symposia. 2006;**245–246**(1):406–409. DOI: 10.1002/

[28] Bento DC, Louarn G, de Santana H. Structural stability and improved properties of poly(3‐alkylthiophenes) synthesized in an acid medium. Journal of Materials Science:

[29] de Santana H, Maia ECR, Bento DC, Cervantes TNM, Moore GJ. Spectroscopic study of poly(3‐alkylthiophenes) electrochemically synthesized in different conditions. Journal of Materials Science: Materials in Electronics. 2013;**24**(9):3352–3358. DOI: 10.1007/

[30] Silva TH, Barreira SVP, Moura C, Silva F. Electrochemical characterization of a self‐ assembled polyelectrolyte film. Portugaliae Electrochimica Acta. 2003;**21**(3):281–292. [31] Bento DC, Maia ECR, Cervantes TNM, Olivati CA, Louarn G, de Santana H. C. Complementary study on the electrical and structural properties of poly(3‐alkylthio‐ phene) and its copolymers synthesized on ITO by electrochemical impedance and Raman spectroscopy. Journal of Materials Science: Materials in Electronics. 2015;**26**(1):149–161.

[32] Skompska M, Szkurłat A. The influence of the structural defects and microscopic aggre‐ gation of poly(3‐alkylthiophenes) on electrochemical and optical properties of the polymer films: discussion of an origin of redox peaks in the cyclic voltammograms. Electrochimica Acta. 2001;**46**(26–27):4007–4015. DOI: 10.1016/S0013‐4686(01)00710‐1 [33] Obaid AY, El‐Mossalamy EH, Al‐Thabaiti SA, El‐Hallag IS, Hermas AA, Asiri AM. Electrodeposition and characterization of polyaniline on stainless steel surface via cyclic, convolutive voltammetry and SEM in aqueous acidic solutions. International Journal of

[34] Eckhardt H, Shacklette LW, Jen KY, Elsenbaumer RL. The electronic and electrochemical properties of poly(phenylene vinylenes) and poly(thienylene vinylenes): an experimen‐ tal and theoretical study. The Journal of Chemical Physics. 1989;**91**(2):1303–1315. DOI:

[35] Österbacka R, An CP, Jiang XM, Vardeny ZV. Two‐dimensional electronic excitations in self‐assembled conjugated polymer nanocrystals. Science. 2000;**287**(5454):839. DOI:

[36] Kobayashi T, Hamazaki J‐i, Kunugita H, Ema K, Endo T, Rikukawa M, et al. Coexistence of photoluminescence from two intrachain states in polythiophene films. Physical

Review B. 2003;**67**(20):205214. DOI: 10.1103/PhysRevB.67.205214

Materials in Electronics. 2016;**27**(5):5371–5382. DOI: 10.1007/s10854‐016‐4437‐1

(1–3):94–98. DOI: 10.1016/j.cplett.2008.12.057

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

masy.200651357

s10854‐013‐1254‐7

10.1063/1.457153

10.1126/science.287.5454.839

DOI: 10.1007/s10854‐014‐2377‐1

Electrochemical Science. 2014;**9**(2):1003–1015.


#### **Magnetic Properties of Hausmannite Thin Films** Provisional chapter

Magnetic Properties of Hausmannite Thin Films

Petya Petkova

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

http://dx.doi.org/10.5772/66533 Additional information is available at the end of the chapter

#### Abstract

The magnetic properties of hausmannite thin films are investigated in this chapter. The Verdet constant and angle of Faraday rotation are determined. The magnetic anisotropy of Mn3O4 is explained by the measurement of the zero-field cooled (ZFC) and field cooled (FC) curves. This experiment is connected with the presentation of the ferromagnetic to superparamagnetic transition of the hausmannite.

Keywords: hausmannite Mn3O4, Faraday effect, verdet constant, ferromagnet, superparamagnet

### 1. Introduction

The hausmannite Mn3O4 can be fabricated by many methods, but the spray pyrolysis method can give it the highest quality. This material is very interesting because it is a transition metal oxide and has application in semiconductor devices [1]. This oxide has two valance states on manganese—Mn2+ and Mn3+. Thus, spinel Mn3O4 occurs in nature as the mineral hausmannite [Mn2+Mn2 3+O4 ]. The Mn2+ cations occupy the tetrahedral sites and Mn3+cations occupy the octahedral sites [2]. The nanoparticles of Mn3O4 thin film behave as single-domain ferromagnets. However, above the blocking temperature, the particles behave as paramagnets due to the dominance of thermal fluctuations over the magnetocrystalline anisotropy energy. These nanoparticles have much higher magnetic moments than other paramagnets and are called superparamagnet.

The detailed investigation of magnetic properties of hausmannite thin film is presented in this chapter.

© 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.

© 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 eproduction in any medium, provided the original work is properly cited.

## 2. Method of preparation and characterization techniques

Several techniques have been used to prepare thin films of these types of transparent and conductive materials to meet the requirements of search and industries such as MOCVD (organometallic chemical vapor deposition) [3], chemical vapor transport (CVT) [4], sputtering [5] and laser ablation [6, 7], which are generally either sophisticated or expensive and hence the need for a simple, easy to meter out and less expensive technique. In addition to these techniques, spray pyrolysis [8–11] has received a little bit of extra attention because of its simplicity and cost-effectiveness as it does not require sophisticated vacuum apparatus. Furthermore, this method can be selected for film production of large area with size grain controllable by controlling the doping concentration. Also, this technique leads to a large production area and it permits also the formation of thin films with possible control of oxygen vacancy by means of the use of both appropriate precursors and postannealing treatments in air [12–15].

Thin films of Mn3O4 were grown at 350°C on 1 × 2 cm2 glass substrate using the spray pyrolysis technique. The substrate temperature was fixed using a digital temperature controller with a k-type thermocouple. The aqueous solution with a flow of about 4 ml/min contains magnesium chloride (MnCl2.6H2O) 0.1 M as precursor. The distance between the nozzle and the substrate was about 27 cm. Spray solutions quantity (75 ml) was kept fixed during the growth. The filtered compressed nitrogen air was used as gas carrier at a flow of 4 l/min. The total deposition time was maintained at 20 min. After deposition, the coated substrates were allowed to cool down naturally to room temperature (Figure 1).

Figure 1. The experimental set up for the spray pyrolysis technique.

The crystalline structure was analyzed by X-ray diffraction, using a Siemens D500 diffractometer with monochromatic CuKα radiation (λ = 1.54 Å) [16]. The surface morphology of the Mn3O4 thin film is further analyzed using atomic force microscope (AFM) using a Veeco digital instrument 3 A microscope. The sample was probed in a tapping mode with a nanometer scale.

### 3. Magnetic study

2. Method of preparation and characterization techniques

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

allowed to cool down naturally to room temperature (Figure 1).

Figure 1. The experimental set up for the spray pyrolysis technique.

Several techniques have been used to prepare thin films of these types of transparent and conductive materials to meet the requirements of search and industries such as MOCVD (organometallic chemical vapor deposition) [3], chemical vapor transport (CVT) [4], sputtering [5] and laser ablation [6, 7], which are generally either sophisticated or expensive and hence the need for a simple, easy to meter out and less expensive technique. In addition to these techniques, spray pyrolysis [8–11] has received a little bit of extra attention because of its simplicity and cost-effectiveness as it does not require sophisticated vacuum apparatus. Furthermore, this method can be selected for film production of large area with size grain controllable by controlling the doping concentration. Also, this technique leads to a large production area and it permits also the formation of thin films with possible control of oxygen vacancy by means of the use of both appropriate precursors and postannealing treatments in air [12–15]. Thin films of Mn3O4 were grown at 350°C on 1 × 2 cm2 glass substrate using the spray pyrolysis technique. The substrate temperature was fixed using a digital temperature controller with a k-type thermocouple. The aqueous solution with a flow of about 4 ml/min contains magnesium chloride (MnCl2.6H2O) 0.1 M as precursor. The distance between the nozzle and the substrate was about 27 cm. Spray solutions quantity (75 ml) was kept fixed during the growth. The filtered compressed nitrogen air was used as gas carrier at a flow of 4 l/min. The total deposition time was maintained at 20 min. After deposition, the coated substrates were

The crystalline structure was analyzed by X-ray diffraction, using a Siemens D500 diffractometer with monochromatic CuKα radiation (λ = 1.54 Å) [16]. The surface morphology of the Mn3O4 thin film is further analyzed using atomic force microscope (AFM) using a Veeco digital The magneto-optic Faraday effect presents the connection between optics, magnetism and atomic physics. Faraday rotation manifests itself as a rotation of the polarization plane of the light passing through the sample in the presence of a magnetic field and is characterized by the Verdet constant (V) of the investigated sample (Figure 2). The rotation angle ϕ can be expressed by the formula [17]:

$$
\varphi(\lambda) = A/(\lambda^2 - \lambda\_0^{-2}),
\tag{1}
$$

where A is a constant determined from the matrix elements of the interband transitions, λ is the wavelength and λ<sup>0</sup> is the wavelength related to the interband transitions and corresponding to the natural frequency ω<sup>0</sup> = 2π/λ<sup>0</sup> of an effective harmonic oscillator. The relationship between the rotation angle and the Verdet constant is ϕ = VBl, where B is the magnetic induction of the field and l is the sample thickness (Figure 3). The magneto-optic anomaly factor γ (Figure 4) can be taken as a measure of the degree of covalency that exists in the bonds connecting the ions and atoms [18]:

$$\gamma = \frac{\varphi}{\frac{c}{2mc^2} \lambda D} \tag{2}$$

Figure 2. The Verdet constant for Mn3O4 in the spectral region 500–2500 nm.

In the paramagnetic materials, the anomaly factors γ can vary with the wavelength of the incident light, even if there is only one absorption frequency contributing to dispersion. The dispersivity of the investigated crystal can be presented by the following equation:

Figure 3. The Faraday rotation angle as a function of the wave length (500–2500 nm) for Mn3O4 thin film.

Figure 4. The magneto-optic anomaly factor γ for hausmannite in the spectral region 500–2500 nm.

$$D = -\frac{\gamma}{c} \frac{d^2 n}{d\lambda^2}.\tag{3}$$

$$\mathcal{V} = -\frac{VBl}{\frac{e\lambda^2}{2mc^2}\frac{d^2n}{d\lambda^2}}.\tag{4}$$

The spectral dependence of the spin-spin exchange interaction constant K in the case of Mn3O4 thin film is presented in Figure 5. It can be calculated by the following formula [19]:

$$K(\lambda) = \frac{V\left[\left(\left(\frac{\lambda}{\lambda\_t}\right)^2 - 1\right)\right]^{3/2}}{\lambda^{\lambda}},\tag{5}$$

where V is the Verdet constant, χ is the magnetic susceptibility of the sample (Figure 6) and λ<sup>g</sup> represents the band gap of the material. For the investigated vanadium doped crystal λ<sup>g</sup> = 556 nm [16]. When the spin-spin exchange interaction constant has negative values, the spins align antiparallel to each other so that the net magnetization is zero. Therefore, the material is antiferromagnet and it is modeled to be made up of two sublattices [20].

Figure 5. The constant of spin-spin exchange interaction K(λ) for Mn2+ (a) and Mn3+ (b) in Mn3O4 thin film (500–2500 nm).

Figure 6. The dependence χ(H)=dM/dH for Mn3O4 in the magnetic field from −20,000 to 20,000 Oe.

D ¼ − γ c d2 n

Figure 4. The magneto-optic anomaly factor γ for hausmannite in the spectral region 500–2500 nm.

Figure 3. The Faraday rotation angle as a function of the wave length (500–2500 nm) for Mn3O4 thin film.

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

<sup>γ</sup> <sup>¼</sup> <sup>−</sup> VBl eλ<sup>2</sup> 2mc<sup>2</sup> d2n dλ<sup>2</sup>

The spectral dependence of the spin-spin exchange interaction constant K in the case of Mn3O4

V <sup>λ</sup> λg <sup>2</sup> −1 <sup>3</sup>=<sup>2</sup>

where V is the Verdet constant, χ is the magnetic susceptibility of the sample (Figure 6) and λ<sup>g</sup> represents the band gap of the material. For the investigated vanadium doped crystal

thin film is presented in Figure 5. It can be calculated by the following formula [19]:

KðλÞ ¼

<sup>d</sup>λ<sup>2</sup> : (3)

χλ , (5)

: (4)

The exchange interaction energy leads to the alignment of neighboring atomic moments and this forms magnetic domains. The magnetostatic interaction energy tries to break them into smaller domains oriented antiparallel to each other. The domain size depends on the relative counterbalance between both energies. The system is composed of a single domain, when the magnetostatic energy does not allow the breaking of domains into smaller parts. This condition is connected with the critical value rc of the radius of a spherical particle. If the rotation of the atomic magnetic moments is coherent (the structure is a single-domain one), then the particle can be characterized by its total magnetic supermoment jμ ! <sup>p</sup>j ¼ MSV, where V is the particle volume and MS is saturation magnetization. The ferromagnetism and super paramagnetism are observed, respectively, below and above the blocking temperature TB. Its origin is connected with magnetic anisotropy within particles. This anisotropy tends to orientate the particle supermoment along some preferential direction.

The spin-orbit coupling and dipolar interaction dictate preferential orientation directions of the magnetic moments because of the finite size of the particles. The magnetic anisotropy energy EA of the particles can be described by a simple model. This model includes two main contributions: crystalline and shape, which are connected with the core and surface atoms, respectively. When the particles are spherical and the anisotropy is uniaxial crystalline, the considered situation is the

simplest [21]. If the magnetic anisotropy is proportional to the particle volume, then K ! eff ¼ KVn^, where K is the effective uniaxial anisotropy constant (per unit volume) and n^ is the unitary vector describing the easy-magnetization anisotropy. The energy term for the i particle can be written as:

$$E\_A^{(i)} = -\mathbf{K}\_i V\_i \left(\frac{\overrightarrow{\mu}\_i \cdot \hat{n}\_i}{|\overrightarrow{\mu}\_i|}\right)^2 = -\mathbf{K}\_i V\_i \cos^2 \theta,\tag{6}$$

where θ<sup>i</sup> is the angle between the magnetic supermoment of the particle and the easy anisotropy axis (Figure 7). The moment of the particle has therefore two preferred orientations, energetically equivalent, along the easy-magnetization anisotropy axis direction. Both directions are separated by an energy barrier EB of height KiVi.

Figure 7. Schematic drawing of the ideal simplest model of noninteracting and parallel aligned easy axes along the applied field.

If the particles are magneto anisotropic, the calculation of equilibrium magnetization is complicated. The special role for nanoparticles having superficial anisotropy is the violation of local symmetry surroundings and crystal field change that acts on the magnetic ions from the surface. The simplest type of magnetic anisotropy is the easy anisotropy axis.

When external magnetic field is applied over the nanoparticles, it tries to orientate their magnetic moments in the direction of its action. Therefore, if the magnetic field is applied perpendicular of anisotropy axis and the orientation of magnetic moment of the particle is labeled with i, the next equation is fulfilled:

#### Magnetic Properties of Hausmannite Thin Films http://dx.doi.org/10.5772/66533 127

$$E^i = E\_A^{(i)} + E\_Z^{(i)} = -\mathcal{K}\_i V\_i \left(\frac{\overrightarrow{\mu}\_i \hat{n}\_i}{|\overrightarrow{\mu}\_i|}\right)^2 - \overrightarrow{\mu}\_i \overrightarrow{H},\tag{7}$$

where EA and EZ are Zeeman energies.

connected with magnetic anisotropy within particles. This anisotropy tends to orientate the

The spin-orbit coupling and dipolar interaction dictate preferential orientation directions of the magnetic moments because of the finite size of the particles. The magnetic anisotropy energy EA of the particles can be described by a simple model. This model includes two main contributions: crystalline and shape, which are connected with the core and surface atoms, respectively. When the particles are spherical and the anisotropy is uniaxial crystalline, the considered situation is the

where K is the effective uniaxial anisotropy constant (per unit volume) and n^ is the unitary vector describing the easy-magnetization anisotropy. The energy term for the i particle can be written as:

where θ<sup>i</sup> is the angle between the magnetic supermoment of the particle and the easy anisotropy axis (Figure 7). The moment of the particle has therefore two preferred orientations, energetically equivalent, along the easy-magnetization anisotropy axis direction. Both direc-

If the particles are magneto anisotropic, the calculation of equilibrium magnetization is complicated. The special role for nanoparticles having superficial anisotropy is the violation of local symmetry surroundings and crystal field change that acts on the magnetic ions from the

Figure 7. Schematic drawing of the ideal simplest model of noninteracting and parallel aligned easy axes along the

When external magnetic field is applied over the nanoparticles, it tries to orientate their magnetic moments in the direction of its action. Therefore, if the magnetic field is applied perpendicular of anisotropy axis and the orientation of magnetic moment of the particle is

surface. The simplest type of magnetic anisotropy is the easy anisotropy axis.

labeled with i, the next equation is fulfilled:

applied field.

<sup>¼</sup> <sup>−</sup>KiVicos<sup>2</sup>

!

θ, (6)

eff ¼ KVn^,

simplest [21]. If the magnetic anisotropy is proportional to the particle volume, then K

μ ! <sup>i</sup> � n^<sup>i</sup> jμ ! i j

!<sup>2</sup>

particle supermoment along some preferential direction.

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

E<sup>ð</sup>i<sup>Þ</sup>

tions are separated by an energy barrier EB of height KiVi.

<sup>A</sup> ¼ −KiVi

The influence of external magnetic field in the orientation of magnetic supermoments is known as Stoner-Wohlfart model [22]. They assume that the coherent rotation of atomic magnetic moments exists and the magnetic field is applied at a certain angle θ<sup>0</sup> with respect to the easy anisotropy axis. When the temperature effects are ignored the problem can be solved with minimal number of energetic arguments. The situation is very interesting, when we can describe the change of magnetic moments in dependence of anisotropic energy barrier and the temperature TB. This is the reason for the study of a simple case when the field is applied parallel to the easy anisotropy axis. It should also be noted that the particles are identical and do not interact with each other. The application of the field leads only to their arrangement in the direction of the easy anisotropy axis. Thus, the following equation can be written (Figure 8):

Figure 8. The dependence of anisotropic energy barrier from the external magnetic field (a) and the angle between the applied field and easy anisotropy axis (b).

$$E = -KV\cos^2\theta \text{--} M\_S V H \cos\theta \tag{8}$$

when H < 2K=MS, Eq. (8) gives two local minima (the directions of easy magnetization) at θ = 0, π with values Emin ¼ −KV � MSVH (Figure 9) and one maximum (the direction of hard magnetization) at θ ¼ arccosðHMS=2KÞ, with value Emax ¼ KVðHMS=2KÞ 2 . The direction of hard magnetization is perpendicular of the anisotropy axis in the case, when H = 0. The value θ = 0 is valid, when the moment of particles is oriented parallel to the magnetic field (↑↑). The equation θ = π is fulfilled when the orientation of the moment of the particles is antiparallel to the magnetic field (↑↓). The difference between the shapes of the energy wells is connected with the different energy barriers. These barriers depend on the orientation of the moments of particles to the applied field which can be written as E↑↓ <sup>B</sup> and <sup>E</sup>↑↑ <sup>B</sup> for the cases of antiparallelism and parallelism, respectively. where EB=kBT ≥ 1.

The anisotropy field of the particles is introduced as HA ¼ 2K=MS. The energy barriers can be calculated as the difference between the minimal and maximal energies: E↑↓ <sup>B</sup> <sup>¼</sup> KV H2 A ðH−HAÞ <sup>2</sup> and E↑↑ <sup>B</sup> <sup>¼</sup> KV H2 A ðH þ HAÞ <sup>2</sup> (Figure 10). The difference between the heights of energy barriers also shows a change in the characteristic time for relaxation of the particles, since it depends on the relative orientation of the magnetic dipoles to the field: antiparallel oriented particles have smaller energy barrier in comparison with the particles which are oriented along the easy anisotropy axis. They have also small heat energy sufficient to overcome the barrier. The parallel oriented particles are limited by deeper anisotropy well and the jump of their magnetic moments requires higher heat energy. When the particle has to rotate its magnetic moment, the energy of jump beyond the energy barrier is EB ≈ KV. The characteristic time of heat fluctuations of the magnetic moments can be presented by the formula [23]:

$$
\pi = \tau\_0 \exp(E\_\mathcal{B}/k\_\mathcal{B}T),
\tag{9}
$$

Figure 9. The local minimal energy � Emin(H) in the directions of easy magnetization for the hausmannite Mn3O4.

Figure 10. The dependence of the energy barriers EB of: applied magnetic field and angle θ in the case of antiparallelism.

The multiplier τ<sup>0</sup> depends on many parameters such as temperature, gyromagnetic ratio, saturation magnetization, anisotropy constant and the size of the energy barrier. It is of the order of 10−<sup>9</sup> – 10−<sup>13</sup> s [24]. The formula (9) determines the characteristic time of establishing heat equilibrium in the system of noninteracting single-domain magnetic particles. At high temperatures, the following inequality is fulfilled: EB=kBT≪1. Therefore, the time of transition of the system in the state with minimal energy is small in comparison with the characteristic time of measurement τm. In this case, the system should not appear magnetic hysteresis. When EB=kBT ≫1, the time of transition of a system in the equilibrium state depends on the size of the particles. If τ<sup>m</sup> > τ, the system is in the super paramagnetic state and it quickly reaches equilibrium magnetization, when the temperature or the external field change. In the opposite case, when the external magnetic field changes, the system does not fail to relax in the new equilibrium state for the time τ<sup>m</sup> and its magnetization does not change. The case when τ<sup>m</sup> = τ is connected with the blocking temperature:

$$T\_B = \frac{KV}{25k\_B}.\tag{10}$$

Formula (10) presents the temperature TB, when the magnetic field is zero. This temperature decreases with the increasing of the external magnetic field by the law:

$$T\_B(H) = T\_B(0)\left(1 - \frac{H}{H\_c}\right)^k,\tag{11}$$

where <sup>k</sup> = 2 for small fields and <sup>k</sup> = 2/3 for big fields and Hc <sup>¼</sup> <sup>2</sup><sup>k</sup> MS .

The anisotropy field of the particles is introduced as HA ¼ 2K=MS. The energy barriers can be

shows a change in the characteristic time for relaxation of the particles, since it depends on the relative orientation of the magnetic dipoles to the field: antiparallel oriented particles have smaller energy barrier in comparison with the particles which are oriented along the easy anisotropy axis. They have also small heat energy sufficient to overcome the barrier. The parallel oriented particles are limited by deeper anisotropy well and the jump of their magnetic moments requires higher heat energy. When the particle has to rotate its magnetic moment, the energy of jump beyond the energy barrier is EB ≈ KV. The characteristic time of heat fluctuations of the

Figure 9. The local minimal energy � Emin(H) in the directions of easy magnetization for the hausmannite Mn3O4.

Figure 10. The dependence of the energy barriers EB of: applied magnetic field and angle θ in the case of antiparallelism.

<sup>2</sup> (Figure 10). The difference between the heights of energy barriers also

τ ¼ τ0expðEB=kBTÞ, (9)

<sup>B</sup> <sup>¼</sup> KV H2 A ðH−HAÞ

<sup>2</sup> and

calculated as the difference between the minimal and maximal energies: E↑↓

magnetic moments can be presented by the formula [23]:

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

E↑↑ <sup>B</sup> <sup>¼</sup> KV H2 A

ðH þ HAÞ

where EB=kBT ≥ 1.

The magnetization curve increasing to reaching saturation magnetization is measured in the study of the magnetic properties of the hausmannite Mn3O4 which containing nano-objects. To determine the temperature dependence of the magnetic moment Mare carried out two types of measurements—cooling in zero magnetic field (zero-field-cooling, ZFC) and cooling in a nonzero field (field-cooling, FC). The sample is cooled (to liquid helium temperature) during the method of ZFC in the absence of a magnetic field and then a small field (2–5 kOe) is included. The temperature values begin slowly to increase and the magnetic moment (MZFC) values can be registered. The technique FC differs from ZFC only by the fact that a sample is cooled in a nonzero magnetic field. The curves MZFC(T) and MFC(T) for the magnetic nanoobjects coincide at sufficiently high temperatures, but they begin to vary below a temperature TH (irreversibility temperature). The curve MZFC(T) has a maximum at a certain temperature Tmax and it increases monotonically down to very low temperatures (Figure 11). The dependence of the magnetization from the applied field at two various temperatures is shown in Figure 12(a) and (b). For an idealized system containing similar nanoparticles with uniaxial anisotropy and random orientation of easy magnetization axis, the difference of the temperature dependence MZFC(T) and MFC(T) at a qualitative level follows from Eq. (8). In the case of zero field during the cooling below the blocking temperature, the magnetic moments of the particles are oriented along their axes of easy magnetization (θ = 0 in Eq. (8)). The total magnetic moment of the system is zero in the beginning and in the end of the cooling process. The magnetic moments for which θ < 90° (see Eq. (8)) it is not necessary to overcome the energy barrier, when the external field H is included. Therefore, they turn to a position with minimum energy, creating a nonzero magnetization of the system. In contrast, the magnetic moments for which the external field is included (θ > 90°) are separated from the minimum energy of the potential barrier. They can overcome only this barrier for a very long time (see Eq. (9)). Therefore, in the case of ZFC measurements (T < TB), the system is in a metastable state with a small total magnetic moment <sup>M</sup><sup>2</sup> SH <sup>3</sup>kV , which does not depend on the temperature.

Figure 11. FC and ZFC induced magnetization as a function of temperature measured in a 2500 Oe field.

Figure 12. The magnetization M(H) of hausmannite thin film for two temperatures T = 10 K (a) and T = 35 K (b).

At T = TB, the system jumps into a stable superparamagnetic state with magnetic moment

$$M\_{\rm ZFC} \approx \frac{M\_{\rm S}^2 VH}{\mathfrak{B}k\_B T} \tag{12}$$

When MSVH≪kBT and random orientation of the easy magnetization axes of particles, formula (12) is also valid for T > TB. The sample is cooled in a nonzero magnetic field during FC measurement and the magnetization at temperatures above TB is determined by formula (12). At T < TB, the system cannot change its magnetization during the measurement. Therefore, the magnetic moment which is determined by the FS method for T < TB is

energy barrier, when the external field H is included. Therefore, they turn to a position with minimum energy, creating a nonzero magnetization of the system. In contrast, the magnetic moments for which the external field is included (θ > 90°) are separated from the minimum energy of the potential barrier. They can overcome only this barrier for a very long time (see Eq. (9)). Therefore, in the case of ZFC measurements (T < TB), the system is in a metastable state

At T = TB, the system jumps into a stable superparamagnetic state with magnetic moment

Figure 12. The magnetization M(H) of hausmannite thin film for two temperatures T = 10 K (a) and T = 35 K (b).

MZFC <sup>≈</sup> <sup>M</sup><sup>2</sup>

When MSVH≪kBT and random orientation of the easy magnetization axes of particles, formula (12) is also valid for T > TB. The sample is cooled in a nonzero magnetic field during FC

<sup>S</sup>VH

<sup>3</sup>kBT (12)

<sup>3</sup>kV , which does not depend on the temperature.

SH

Figure 11. FC and ZFC induced magnetization as a function of temperature measured in a 2500 Oe field.

with a small total magnetic moment <sup>M</sup><sup>2</sup>

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

$$M\_{\rm ZFC} \approx \frac{M\_{\rm S}^2 VH}{3k\_B T} = \text{const.}\tag{13}$$

The curves MZFC(T) and MFC(T) are not separated at T = TB for a system consisting of a single-domain nanoparticles with dispersion of the size, the shape, etc. Their separation realizes at a higher temperature TH > TB, where TH is called the irreversibility point. Another characteristic point on the curve MZFC(T) is the temperature Tmax which is often equated with the average blocking temperature of the system <TB>. At temperatures below <TB> we can observe the increase of MFC(T) that replaced section "saturation" and sometimes we can observe a maximum [25]. The value of TH can be identified with the blocking temperature for the particles with a maximum size and the temperature Tmax corresponds to a blocking temperature for particles with minimum size. However, all of these characteristic temperatures (as well as their relationship with the particle size distribution by volume) may depend on the cooling rate and the subsequent heating of the sample. The intensity of the interaction between the particles also influences on the characteristics temperatures. If the heating rate of the sample is much smaller than its cooling rate, a maximum may be formed on the curve MFC(T) at T < <TB> [26].

Note that the difference between the curves MZFC(T) and MFC(T) is not observed only in systems of magnetic nano-objects, but also in macroscopic magnets with disorder elements (frustration of exchange bonds, topological disorder, structural defects) and even in ordered ferromagnets with a large magnetic anisotropy. The difficulties of theoretical research of the magnetic hysteresis in nano-objects consist in the fact that the phenomenon is nonlinear, nonequilibrium and nonlocal and it is caused by the existence of energy minima (due to the magnetic anisotropy) and separated barriers which have complicated dependence of the external magnetic field. The results of theoretical studies of simple models rarely give an acceptable description for real magnetic nanomaterials since their microstructure do not account, in particular, the influence of boundaries and defects on the local magnetization is not taken into account. The ferromagnetic-to-superparamagnetic transition of the hausmannite particles in the zero-field cooled (ZFC) and field cooled (FC) curves is presented in Figure 8.

The orientation of electron spin in the manganese ions is very interesting for study. One of the electrons of the inner shell is responsible for the magnetism and its spin is oriented upwards. If the conductivity electrons move in the same region, where there is the motion of "magnetic" electrons than their spins rotate in the opposite direction. Thus, the conductivity electrons can rotate the electron spins of the other ions. This double interaction is equivalent of the interaction between two "magnetic" electrons which are oriented in one direction. This means that the neighboring spins have to be parallel, which is a result from the action of intermediate environment. This mechanism does not require all electrons to be oriented upwards. It is sufficient that conductivity electrons can be slightly oriented downwards. Thus, the possibility for the rotation of "magnetic" electrons upwards increases.

The energy of electron spin can be presented as (Figure 13):

$$\mathbf{x} = |\mu| \left( H + \frac{\lambda M}{\varepsilon\_0 c^2} \right), \tag{14}$$

where μ = 2.8363\*10−<sup>4</sup> eVT−<sup>1</sup> , λ = 5700 m−<sup>1</sup> , ε<sup>0</sup> = 8.8542\*10−<sup>12</sup> Fm−<sup>1</sup> and c = 8\*108 m/s (Figure 13).

Figure 13. The energy of electron spin x as a function of the intensity of magnetic field H for Mn3+ ions at temperatures T = 10 K and T = 35 K.

On the other hand, we can write that

$$\mathbf{x} = |m|(H + dM/\varepsilon\_0 c^2)/kT,\tag{15}$$

where d = 2.2971 × 10−<sup>10</sup> m.

The magnetic moment of the electron is

$$|m| = \frac{|\mu| \left( H + \frac{\lambda M}{\varepsilon\_0 c^2} \right)}{(H + dM/\varepsilon\_0 c^2)/kT} = 2.444110^{-7},\tag{16}$$

where k = 8.6173 × 10−<sup>5</sup> eVK−<sup>1</sup> :

$$th\mathbf{x} = th[|m|(H + dM/\varepsilon\_0 \varepsilon^2)/kT] \tag{17}$$

The energy of interaction between two electrons is expressed by the next equation (Figure 14):

Magnetic Properties of Hausmannite Thin Films http://dx.doi.org/10.5772/66533 133

$$
\langle \mathrm{II} \rangle = -\mathrm{N} |\mu| \left( H + \frac{M}{2\varepsilon\_0 c^2} \right) \hbar \mathrm{lx} \tag{18}
$$

Figure 14. The average value of the energy of interaction between two electrons in Mn3+ ions as a function of the intensity of magnetic field H: (a) T = 10 K and (b) T = 35 K.

#### 4. Conclusions

The energy of electron spin can be presented as (Figure 13):

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

, λ = 5700 m−<sup>1</sup>

where μ = 2.8363\*10−<sup>4</sup> eVT−<sup>1</sup>

On the other hand, we can write that

The magnetic moment of the electron is

jmj ¼

:

where d = 2.2971 × 10−<sup>10</sup> m.

10 K and T = 35 K.

where k = 8.6173 × 10−<sup>5</sup> eVK−<sup>1</sup>

x ¼ jμj H þ

x ¼ jmjðH þ dM=ε0c

Figure 13. The energy of electron spin x as a function of the intensity of magnetic field H for Mn3+ ions at temperatures T =

<sup>j</sup>μ<sup>j</sup> <sup>H</sup> <sup>þ</sup> <sup>λ</sup><sup>M</sup> ε0c<sup>2</sup> 

thx ¼ th½jmjðH þ dM=ε0c

The energy of interaction between two electrons is expressed by the next equation (Figure 14):

<sup>ð</sup><sup>H</sup> <sup>þ</sup> dM=ε0c<sup>2</sup>Þ=kT <sup>¼</sup> <sup>2</sup>:444110<sup>−</sup><sup>7</sup>

2

2

λM ε0c<sup>2</sup> 

, (14)

Þ=kT, (15)

, (16)

Þ=kT� (17)

, ε<sup>0</sup> = 8.8542\*10−<sup>12</sup> Fm−<sup>1</sup> and c = 8\*108 m/s (Figure 13).

The magneto-optic anomaly factor γ for the hausmannite thin films decreases with the increasing of the wave length. The spin-spin exchange interaction constant K decreases to λ<sup>1</sup> = 2172 nm (for Mn2+ ions) and λ<sup>2</sup> = 2180 nm (for Mn3+ ions) and after that it begins to increase. The magnetic susceptibility of Mn3O4 has maximal value, when the intensity of applied magnetic field is 5868 Oe. The anisotropic energy barrier decreases with the increasing of H and θ. The energy barrier increases quadratic with the increasing of H and θ. The energy of electron spin x has bigger values for Mn3+ ions in the case when T = 10 K. The values of energy of interaction between two electrons are bigger when T = 35 K.

### Acknowledgements

The author would like to express here gratefulness to Pr. Dr. Ing. Karem Boubaker, Unité de physique des dispositifs à semi-conducteurs, Tunis EL MANAR University, Tunisia, for providing hausmannite thin films for investigation. The author would like to thank for the financial support of the project RD-08-109/08.02.2016 from Shumen University "Konstantin Preslavsky."

### Author details

Petya Petkova

Address all correspondence to: petya232@abv.bg

Department of Experimental Physics, Faculty of Natural Sciences, Konstantin Preslavsky University of Shumen, Shumen, Bulgaria

### References


[11] Kerli S, Alver U, Yaykaşlı H, Appl. Surf. Sci. 318 (2014) 164

field is 5868 Oe. The anisotropic energy barrier decreases with the increasing of H and θ. The energy barrier increases quadratic with the increasing of H and θ. The energy of electron spin x has bigger values for Mn3+ ions in the case when T = 10 K. The values of energy of interaction

The author would like to express here gratefulness to Pr. Dr. Ing. Karem Boubaker, Unité de physique des dispositifs à semi-conducteurs, Tunis EL MANAR University, Tunisia, for providing hausmannite thin films for investigation. The author would like to thank for the financial support of the project RD-08-109/08.02.2016 from Shumen University "Konstantin

Department of Experimental Physics, Faculty of Natural Sciences, Konstantin Preslavsky

[2] Chen W Z, Jiao Z, Wu H M, Shek H C, Wu L M C, Lai L K J, Progr. Mater. Sci. 56 (2011)

[8] Boukhachem A, Boughalmi R, Karyaoui M, Mhamdi A, Chtourou R, Boubaker K,

[10] Sharmaa R, Acharya D A, Shrivastava B S, Shripathi T, Ganesan V, Optik 125 (2014) 6751

[3] Sanga B, Nagoyab Y, Kushiyab K, Yamase O, Sol. Energ. Mat. Sol. Cells 75 (2003) 179

[4] Fu Q, Hu L, Yu D, Sun J, Zhang H, Huo B, Zhao Z, Mater. Lett. 63 (2009) 316

[6] Henley J S, Ashfold R N M, Cherns D, Surf. Coat. Technol. 177–178 (2004) 271

[5] Shimomura T, Kim D, Nakayama M, J. Lumin. 112 (2005) 191

Amlouk M, Mater. Sci. Eng. B 188 (2014) 72

[7] Sasi B, Gopchandran G K, Sol. Energ. Mat. Sol. Cells 91 (2007) 1505

[9] Krunks M, Soon J, Unt T, Mere A, Mikli V, Vacuum 107 (2014) 242

[1] Xiao J, Yang S, Wan I, Xiao F, Wang S, J. Power Sources 245 (2014) 1027

between two electrons are bigger when T = 35 K.

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

Address all correspondence to: petya232@abv.bg

University of Shumen, Shumen, Bulgaria

Acknowledgements

Preslavsky."

Author details

Petya Petkova

References

901


### **Advance Deposition Techniques for Thin Film and Coating Advance Deposition Techniques for Thin Film and Coating**

Asim Jilani , Mohamed Shaaban Abdel-wahab and Ahmed Hosny Hammad Asim Jilani, Mohamed Shaaban Abdel-wahab and Ahmed Hosny Hammad

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

http://dx.doi.org/10.5772/65702

#### **Abstract**

Thin films have a great impact on the modern era of technology. Thin films are considered as backbone for advanced applications in the various fields such as optical devices, environmental applications, telecommunications devices, energy storage devices, and so on . The crucial issue for all applications of thin films depends on their morphology and the stability. The morphology of the thin films strongly hinges on deposition techniques. Thin films can be deposited by the physical and chemical routes. In this chapter, we discuss some advance techniques and principles of thin-film depositions. The vacuum thermal evaporation technique, electron beam evaporation, pulsed-layer deposition, direct current/radio frequency magnetron sputtering, and chemical route deposition systems will be discussed in detail.

**Keywords:** thin films, coatings, physical deposition, sol-gel, chemical bath deposition, chemical route

#### **1. Introduction**

Nowadays, most of the technologies are used for minimizing the materials into nano-size as well as nano-thickness leading to the emergence of new and unique behaviors of such materials in optical, electrical, optoelectronic, dielectric applications, and so on. Hence, a new branch of science/materials science is called thin films or coatings. Thin film can be defined as a thin layer of material, where the thickness is varied from several nanometers to few micrometers. Like all materials, the structure of thin films is divided into amorphous and polycrystalline structure depending on the preparation conditions as well as the material nature. Thin films comprise

© 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.

two parts: the layer and the substrate where the films are deposited on it. Also, thin films can be composed of different layers such as thin-film solar cells, electrochromic cells, and so on.

In order to obtain thin films with good quality, there are two common deposition techniques: physical and chemical depositions. It can be summarized as shown in **Table 1**.


**Table 1.** Methods of thin films deposition.

This chapter describes some common deposition techniques for thin films in detail to give some confidential and important points of view for readers on how thin films can be formed.

### **2. Physical deposition techniques**

### **2.1. Evaporation techniques**

Evaporation methods are considered as the common deposition of materials in the form of thin-layer films. The general mechanism of these methods is obtained by changing the phase of the material from solid phase to vapor phase and converting again to solid phase on the specific substrate. It takes place under vacuum or controlled atmospheric condition.

### *2.1.1. Vacuum thermal evaporation technique*

Vacuum evaporation technique is the simplest technique used to prepare amorphous thin films especially chalcogenide films such as CdSSe [1], MnS [2], Ge-Te-Ga [3], and so on. In general, chalcogenide materials can be used for memory-switching applications [4, 5], phase-change materials [6, 7], and solar applications [8].

The technique of thermal evaporation is strongly dependent on two parameters: thermally vaporized material and applying a potential difference to the substrate under medium- or higher-vacuum level ranging from 10−5 to 10−9 mbar. The schematic diagram for thermal evaporation is shown in **Figure 1** taken from elsewhere [9].

**Figure 1.** Schematic of thermal evaporation system with substrate holder on a planetary rotation system and directly above the evaporating source.

#### *2.1.2. Electron beam evaporation*

two parts: the layer and the substrate where the films are deposited on it. Also, thin films can be composed of different layers such as thin-film solar cells, electrochromic cells, and so on.

In order to obtain thin films with good quality, there are two common deposition techniques:

This chapter describes some common deposition techniques for thin films in detail to give some confidential and important points of view for readers on how thin films can be formed.

Evaporation methods are considered as the common deposition of materials in the form of thin-layer films. The general mechanism of these methods is obtained by changing the phase of the material from solid phase to vapor phase and converting again to solid phase on the

Vacuum evaporation technique is the simplest technique used to prepare amorphous thin films especially chalcogenide films such as CdSSe [1], MnS [2], Ge-Te-Ga [3], and so on. In general, chalcogenide materials can be used for memory-switching applications [4, 5], phase-change

specific substrate. It takes place under vacuum or controlled atmospheric condition.

**1. Sol-gel technique**

**4. Plating**

**2. Chemical bath deposition 3. Spray pyrolysis technique**

**a.** Electroplating technique. **b.** Electroless deposition. **5. Chemical vapor deposition (CVD) a.** Low pressure (LPCVD)

**b.** Plasma enhanced (PECVD)

**c.** Atomic layer deposition (ALD)

physical and chemical depositions. It can be summarized as shown in **Table 1**.

**Physical deposition Chemical deposition**

**1. Evaporation techniques**

**a.** Vacuum thermal evaporation. **b.** Electron beam evaporation.

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

**c.** Laser beam evaporation.

**e.** Molecular beam epitaxy. **f.** Ion plating evaporation.

**Table 1.** Methods of thin films deposition.

**2.1. Evaporation techniques**

**a.** Direct current sputtering (DC sputtering). **b.** Radio frequency sputtering (RF sputtering).

**2. Physical deposition techniques**

*2.1.1. Vacuum thermal evaporation technique*

materials [6, 7], and solar applications [8].

**d.** Arc evaporation.

**2. Sputtering techniques**

This type of evaporation is another method of physical deposition where the intensive beam of electrons is generated from a filament and steered through both electric and magnetic fields to hit the target and vaporize it under vacuum environment as shown in **Figure 2**. Thin films prepared by electron beam evaporation are of good quality and purity [10].

**Figure 2.** Schematic diagram of electron beam evaporation.

Large categories of materials can be prepared by electron beam evaporation technique [11] such as amorphous and crystalline semiconductors [12], metals [13], oxides [14], and molecular materials [15].

#### *2.1.3. Laser beam evaporation (pulsed-laser deposition)*

Pulsed-laser deposition (PLD) is another physical deposition technique to deposit the thinfilm-coating system [16]. During the thin-film deposition process, the laser beam is used to ablate the material for depositing the thin films inside a vacuum chamber as shown in **Figure 3**.

**Figure 3.** Schematic of pulsed-laser deposition taken from Ref. [17].

Different kinds of laser sources are being used to ablate the target. The most common sources are Nd-YAG laser, KrF (248 nm), and XeCl (308 nm). When the laser beam strikes the target material, it produces the plume which could deposit on the various substrates. The created plume may contain neural- and ground-state atoms and ionized species. In the case of metal oxide thin films, oxygen is used to deposit the oxides of metals [18]. The thin-film quality from the PLD depends on the various parameters such as wavelength of the laser, energy, ambient gas pressure, pulsed duration, and the distance of the target to the substrate [19]. The ablation process during the deposition may control and monitor by using laser-induced fluorescence [20], laser ablation molecular isotopic spectroscopy [21], and optical emission spectroscopy [22]. The morphology of the deposited thin films is also affected by the substrate temperature. The coating of thin films through PLD follows three modes: Frank–-van der Merwe, Stranski–- Krastanov, and Volmer–-Weber [23, 24]. PLD has some advantages over other physical deposition systems because of its fast deposition time and its compatibility to oxygen and other inert gases.

#### **2.2. Sputtering technique**

Sputtering technique is mostly used for depositing metal and oxide films by controlling the crystalline structure and surface roughness [11, 25]. The simple form of the sputtering system consists of an evacuated chamber containing metallic anode and cathode [25] in order to obtain a glow discharge in the residual gas in the chamber. Also, an applied voltage in the order of several KeV with pressure more than 0.01 mbar is sufficient for film deposition. The sputtering process depends on the bombardment of the ions released from the discharge to the molecules in the cathode leading to the liberation of the molecules from the cathode with higher kinetic energy. The atomic weight of the bombarding ions should be nearly to that of the target material in order to maximize the momentum transfer. These molecules move in straight lines and strike on the anode or on the substrate to form a dense thin film [25]. The diagram of the sputtering system is shown in **Figure 4**.

**Figure 4.** Sputtering system diagram.

Large categories of materials can be prepared by electron beam evaporation technique [11] such as amorphous and crystalline semiconductors [12], metals [13], oxides [14], and molecular

Pulsed-laser deposition (PLD) is another physical deposition technique to deposit the thinfilm-coating system [16]. During the thin-film deposition process, the laser beam is used to ablate the material for depositing the thin films inside a vacuum chamber as shown in **Figure 3**.

Different kinds of laser sources are being used to ablate the target. The most common sources are Nd-YAG laser, KrF (248 nm), and XeCl (308 nm). When the laser beam strikes the target material, it produces the plume which could deposit on the various substrates. The created plume may contain neural- and ground-state atoms and ionized species. In the case of metal oxide thin films, oxygen is used to deposit the oxides of metals [18]. The thin-film quality from the PLD depends on the various parameters such as wavelength of the laser, energy, ambient gas pressure, pulsed duration, and the distance of the target to the substrate [19]. The ablation process during the deposition may control and monitor by using laser-induced fluorescence [20], laser ablation molecular isotopic spectroscopy [21], and optical emission spectroscopy [22]. The morphology of the deposited thin films is also affected by the substrate temperature. The coating of thin films through PLD follows three modes: Frank–-van der Merwe, Stranski–- Krastanov, and Volmer–-Weber [23, 24]. PLD has some advantages over other physical deposition systems because of its fast deposition time and its compatibility to oxygen and other

Sputtering technique is mostly used for depositing metal and oxide films by controlling the crystalline structure and surface roughness [11, 25]. The simple form of the sputtering system

materials [15].

inert gases.

**2.2. Sputtering technique**

*2.1.3. Laser beam evaporation (pulsed-laser deposition)*

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

**Figure 3.** Schematic of pulsed-laser deposition taken from Ref. [17].

The process of sputtering has several advantages. High-melting point materials can be easily formed by sputtering. The deposited films have composition similar to the composition of the starting materials. Sputtering technique is available to use for ultrahigh vacuum applications. The sputtering sources are compatible with reactive gases such as oxygen. Contrarily,' thick coatings cannot be obtained and there is a difficulty to deposit uniformly on complex shapes.

There are two common types of sputtering process: direct current (DC) and radio frequency (RF) sputtering. The first one depends on DC power, which is generally used with electrically conductive target materials. It is easy to control with low-cost option. The RF sputtering uses RF power for most dielectric materials. A common example for sputtered films is aluminum nitride films. These films were prepared by both DC- and RF-sputtering technique, and their structure and optical properties were compared [26, 27].

### **3. Chemical deposition techniques**

Although the production of thin films via physical methods as previously described gives good quality and functionalizes properties, it is highly expensive and perhaps requires a large amount of material target. Since the need to produce good-quality thin films with low economical cost is necessary, chemical deposition techniques are widely used globally. These techniques are cheap producing good-quality films. Most of them do not require expensive equipment. The chemical deposition is strongly dependent on the chemistry of solutions, pH value, viscosity, and so on. The most common chemical deposition has been obtained via solgel route, chemical bath deposition, electrodeposition, chemical vapor deposition (CVD), and spray pyrolysis technique. This section is concerned only on sol-gel and chemical bath deposition techniques because they can form good film quality with low equipment requirement.

#### **3.1. Sol-gel technique**

The sol-gel technique is broadly used for the synthesis of oxide materials [28]. Sol-gel process is one of the famous wet-chemical methods. It works under lower-temperature processing and gives better homogeneity for multicomponent materials. The word "'sol"' means the formation of a colloidal suspension and 'gel' means the conversion of 'sol' to viscous gels or solid materials. Two routes are used to prepare transition metal oxides (TMOs) as follows:


In this section, we are concerned on the famous route "the metal alkoxide precursor solution by an alcoholic solution."

### *3.1.1. Alkoxide precursors in organic solvents*

The sol-gel technique is based on the polycondensation of metal alkoxides **M (OR)***z* in which **R** represents an alkyl group (R = CH3, C2H5, …) and **z** is the oxidation state of the metal atom **M***<sup>z</sup>***<sup>+</sup>** [29]. It can be synthesized via the reaction of metal salt (chloride, acetate, nitrate, etc.) with alcohol as follows:

$$\text{(CH}\_3\text{COO)}\_x\text{M} + \text{zROH} \rightarrow \text{M(OR)}\_x + \text{zCH}\_3\text{COOH} \tag{1}$$

After this process, two important steps should be involved:

**1.** *Hydrolysis:* this step is aimed to form reactive **M-OH** groups [30]:

$$\text{M} \cdot \text{OR} + \text{H} \text{O} \rightarrow \text{M} \cdot \text{OH} + \text{ROH} \tag{2}$$


**•** *Oxolation:* oxolation is a reaction in which an oxo bridge (—O—) is created between two metal centers. When the metal is coordinately unsaturated, oxolation with rapid kinetics leads to edge- or face-shared polyhedral as shown in **Figure 6**.

**Figure 5.** Several types of OH bridges can be formed by olation condensation process.

**Figure 6.** Formation of oxo-bridging links between two metal centers.

Hence, olation process occurs mainly for lower oxidation states of cations **(***z* **< 4),** whereas oxolation is mainly observed with cations of high oxidation state **(***z* **> 4)** [29, 31].

The previous description provides the preparation of the precursor solution. In order to make thin film from the precursor solution, there are two processes for the production of the films, that is, dip-coating and spin-coating techniques.

#### *3.1.2. Dip-coating technique*

amount of material target. Since the need to produce good-quality thin films with low economical cost is necessary, chemical deposition techniques are widely used globally. These techniques are cheap producing good-quality films. Most of them do not require expensive equipment. The chemical deposition is strongly dependent on the chemistry of solutions, pH value, viscosity, and so on. The most common chemical deposition has been obtained via solgel route, chemical bath deposition, electrodeposition, chemical vapor deposition (CVD), and spray pyrolysis technique. This section is concerned only on sol-gel and chemical bath deposition techniques because they can form good film quality with low equipment require-

The sol-gel technique is broadly used for the synthesis of oxide materials [28]. Sol-gel process is one of the famous wet-chemical methods. It works under lower-temperature processing and gives better homogeneity for multicomponent materials. The word "'sol"' means the formation of a colloidal suspension and 'gel' means the conversion of 'sol' to viscous gels or solid

materials. Two routes are used to prepare transition metal oxides (TMOs) as follows:

b. Preparing of metal alkoxide precursors via metal alkoxides in nonaqueous solvents.

In this section, we are concerned on the famous route "the metal alkoxide precursor solution

The sol-gel technique is based on the polycondensation of metal alkoxides **M (OR)***z* in which **R** represents an alkyl group (R = CH3, C2H5, …) and **z** is the oxidation state of the metal atom **M***<sup>z</sup>***<sup>+</sup>** [29]. It can be synthesized via the reaction of metal salt (chloride, acetate, nitrate, etc.) with

**2.** *Condensation*: condensation is the second step after hydrolysis leading to the departure of a water molecule. The process of condensation can be either olation process or oxolation

**•** *Olation:* a hydroxyl bridge ("ol" bridge) is formed between two metal centers as shown in

( 3 3 ) ( ) z z CH COO M + zROH M OR + zCH COOH ® (1)

(2)

a. Preparing of inorganic precursors via inorganic salts in aqueous solution.

ment.

**3.1. Sol-gel technique**

by an alcoholic solution."

alcohol as follows:

process.

**Figure 5**.

*3.1.1. Alkoxide precursors in organic solvents*

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

After this process, two important steps should be involved:

**1.** *Hydrolysis:* this step is aimed to form reactive **M-OH** groups [30]:

Dip-coating technique is almost used to fabricate transparent layers of oxides on a transparent substrate with a high degree of planarity and surface quality [32]. Other substrates are also possible to use. Well-defined film thicknesses up to 1 μm can be deposited. Several additive layers can be superimposed.

**Figure 7.** Dip-coating process levels.

Scriven [33] described the dip-coating process in five stages: immersion, start-up, deposition, drainage, and evaporation. Hence, the evaporation normally accompanies the start-up, deposition, and drainage steps as shown in **Figure 7**.

### *3.1.3. Spin-coating technique*

Another technique is also available for usage after the precursor solution is prepared known as spin coating or spinning. The solution is dripped onto a spinning substrate and spreads evenly. The spinning process is most suitable for the coating of small disks or lenses but is not very economical. The process of spinning film can be described as shown in **Figure 8**.

**Figure 8.** Spin-coating process.

#### **3.2. Chemical bath deposition technique**

possible to use. Well-defined film thicknesses up to 1 μm can be deposited. Several additive

Scriven [33] described the dip-coating process in five stages: immersion, start-up, deposition, drainage, and evaporation. Hence, the evaporation normally accompanies the start-up,

Another technique is also available for usage after the precursor solution is prepared known as spin coating or spinning. The solution is dripped onto a spinning substrate and spreads evenly. The spinning process is most suitable for the coating of small disks or lenses but is not

very economical. The process of spinning film can be described as shown in **Figure 8**.

layers can be superimposed.

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

**Figure 7.** Dip-coating process levels.

*3.1.3. Spin-coating technique*

**Figure 8.** Spin-coating process.

deposition, and drainage steps as shown in **Figure 7**.

Chemical bath deposition method is also known as solution growth technique or controlled precipitations [34]. It is the oldest method to deposit films on a substrate. Solution growth technique is mostly used to prepare chalcogenide films as well as metal oxide films. Also, the deposition can be performed at lower temperatures. In the solution growth method, the precursor solution of metal ions must be complexed by ligands. The complex solution is almost obtained with ammonia solution, triethanol amine, ethylene-diamine-tetraacetic acid (EDTA), citric acid, and so on. When the complexation is completed, the addition of the anions should take place. These anions come from the thiourea, thioacetamide, thiosulfate, and sodium Sulfide solutions [34] as sources of sulfur anions or selenourea and sodium selenosalfate for selenium anions to deposit the chalcogenides. Substrates are put in vertical, horizontal, or specific position inside the solution and left until the desired film thickness is obtained. The deposition of oxide films is quite different than chalcogenides. After making the complexation by controlling the pH value, the substrate is immersed in the solution under the desired temperature varied in the range 60°–100 °C to deposit in most cases the metal hydroxide films. The hydroxide film can then be transferred to oxide by the annealing process. **Figure 9** represents the simple chemical bath deposition method taken from Ref. [34]. Indeed, much

**Figure 9.** Home-made chemical bath deposition technique.

reviews and literature, which describes the chemical bath deposition for both chalcogenide and oxide films, are found elsewhere [34–36].

### **Author details**

Asim Jilani1\*, Mohamed Shaaban Abdel-wahab1 and Ahmed Hosny Hammad1,2

\*Address all correspondence to: asim.jilane@gmail.com

1 Center of Nanotechnology, King Abdulaziz University, Jeddah, Saudi Arabia

2 Electron Microscope and Thin Films Department, Physics Division, National Research Centre, Dokki, Giza, Egypt

### **References**


[8] Salomé PMP, Alvarez HR, Sadewasser S. Incorporation of alkali metals in chalcogenide solar cells. Sol. Ener. Mater. Sol. Cells. 2015;143:9–20. DOI: 10.1016/j.solmat.2015.06.011

reviews and literature, which describes the chemical bath deposition for both chalcogenide

2 Electron Microscope and Thin Films Department, Physics Division, National Research

[1] Hassanien AS, Akl AA. Effect of Se addition on optical and electrical properties of chalcogenide CdSSe thin films. Superlattices Microstruct. 2016;89:153–169. DOI:

[2] Hannachi A, Segura A, Meherzi HM. Growth of manganese sulfide (α-MnS) thin films by thermal vacuum evaporation: Structural, morphological, optical properties. Mater.

[3] Wang G, Nie Q, Shen X, Chen F, Li J, Zhang W, Xu T, Dai S. Phase change and optical band gap behaviour of Ge-Te-Ga thin films prepared by thermal evaporation. Vacuum.

[4] Malligavathy M, Kumar RTA, Das C, Asokan S, Padiyan DP. Growth and characteristics of amorphous Sb2Se3 thin films of various thicknesses for memory switching applications. J. Non-Cryst. Solids. 2015;429:93–97. DOI: 10.1016/j.jnoncrysol.2015.08.038

[5] Kumar RTA, Das C, Lekha PC, Asokan S, Sanjeeviraja C, Padiyan DP. Enhancement in threshold voltage with thickness in memory switch fabricated using GeSe1.5S0.5 thin

[6] Rafea MA, Farid H. Phase change and optical band gap behaviour of Se0.8S0.2 chalcogenide glass films. Mater. Chem. Phys. 2009;113(2–3):868–872. DOI: 10.1016/j.match-

[7] Sangeetha BG, Joseph CM, Suresh K. Preparation and characterization of Ge1Sb2Te4 thin films for phase change memory applications. Microelect. Eng. 2014;127:77–80. DOI:

films. J. Alloys Compd. 2014;615:629–635. DOI: 10.1016/j.jallcom.2014.07.068

Chem. Phys. 2016;181:326–332. DOI: 10.1016/j.matchemphys.2016.06.066

2012;86(10):1572–1575. DOI: 10.1016/j.vacuum.2012.03.036

and Ahmed Hosny Hammad1,2

and oxide films, are found elsewhere [34–36].

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

Asim Jilani1\*, Mohamed Shaaban Abdel-wahab1

\*Address all correspondence to: asim.jilane@gmail.com

1 Center of Nanotechnology, King Abdulaziz University, Jeddah, Saudi Arabia

**Author details**

Centre, Dokki, Giza, Egypt

10.1016/j.spmi.2015.10.044

emphys.2008.08.045

10.1016/j.mee.2014.04.032

**References**


[35] Pawar SM, Pawar BS, Kim JH, Joo OS, Lokhande CD. Recent status of chemical bath deposited metal chalcogenide and metal oxide thin films. Curr. Appl. Phys. 2011;11(2): 117–161. DOI: 10.1016/S0254-0584(00)00217-0

[21] Russo RE, Mao X, Gonzalez JJ, Zorba V, Yoo J . Laser ablation in analytical chemistry.

[22] Geyer TJ, Weimer WA. Parametric effects on plasma emission produced during excimer laser ablation of YBa2Cu3O7-x. Appl. Spectros. 1990;44(10):1659–1664. DOI: 10.1366/

[23] Karl H, Stritzker B. Reflection high-energy electron diffraction oscillations modulated by laser pulse deposition YBa2Cu3O7-x. Phys. Rev. Lett. 1992;69(20):2939–2942. DOI:

[24] Lippmaa M, Nakagawa N, Kawasaki M, Ohashi S, Inaguma, Itoh M, Koinuma H. Stepflow growth of SrTiO3 thin films with a dielectric constant exceeding 104. Appl. Phys.

[25] Angusmacleod H. Recent developments in deposition techniques for optical thin films and coatings. In: Piegari A, Flory F, editors. Optical Thin Films and Coatings from Materials to Applications. Oxford: Woodhead Publishing Series; 2013. p. 3–-25. DOI:

[26] Morosanu C, Dumitru V, Cimpoiasu E, Nenu C. Comparison between DC and RF magnetron sputtered aluminum nitride films. In: Prelas MA, Benedictus A, Lin LTS, Popovici G, Gielisse P, editors. Diamond Based Composites and Related Materials. 1st ed. Petersburg, Russia: Springer Science+Business Media Dardrecht; 1997. p. 127–132.

[27] Dumitru V, Morosanu C, Sandu V, Stoica A. Optical and structural differences between RF and DC AlxNy magnetron sputtered films. Thin Solid Films. 2000;359:17–20. DOI:

[28] Livage J, Sanchez C, Henry M, Doeuff S. The chemistry of sol-gel process. Solid State

[29] Livage J, Ganguli D. Sol-gel electrochromic coatings and devices: A review. Sol. Ener.

[30] Tjong SC, Chen H. Nanocrystalline materials and coatings. Mater. Sci. Eng. R.

[32] Klein LC. Sol-Gel Technology for Thin Films, Fiber, Preform, Electronics and Specialty Shapes. Park Ridge, NJ, USA: Noyes Publications; 1987. DOI: 10.1002/pi.4980210420

[33] Scriven LE. Physics and applications of dip coating and spin coating. In: Brinker CJ, Clark DE, Ulrich DR, editors. Better Ceramics Through Chemistry. 3rd ed. Pittsburgh,

[34] Mane RS, Lokhande CD. Chemical deposition method for metal chalcogenide thin films. Mater. Chem. Phys. 2000;(1):1–31. DOI: 10.1016/S0254-0584(00)00217-0

Ionics. 1989;32–33(2):633–638. DOI: 10.1016/0167-2738(89)90338-X

2004;45(1–2):1–88. DOI: 10.1016/j.mser.2004.07.001

PA: Materials Research Society; 1988. p. 712–729.

Mater. Sol. Cells. 2001;68:365–381. DOI: 10.1016/S0927-0248(00)00369-X

[31] Brinker CJ, Scherer GW. Sol-Gel Science. San Diego: Academic Press; 1990.

Anal. Chem. 2013;85(13):6162–6177. DOI: 10.1021/ac4005327

Lett. 1999; 74(23):3543–3545. DOI: 10.1063/1.124155

0003702904417454

10.1103/PhysRevLett.69.2939

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

10.1533/9780857097316.1.3

DOI: 10.1007/978-94-011-5592-2\_9

10.1016/S0040-6090(99)00726-9

[36] Hammad AH, Elmandouh ZS, Elmeleegi HA. Structure and some physical properties of chemically deposited nickel sulfide thin films. In: Proceedings of the 4th International Congress APMAS 2014; 24–27 April 2014; Fethiye, Turkey. Acta Phys. Polonica A: Polish Academy of Sciences Institute of Physics; 2015. p. 901–903. DOI: 10.12693/APhysPolA. 127.901

### **Vanadium Oxide Thin Films Obtained by Thermal Annealing of Layers Deposited by RF Magnetron Sputtering at Room Temperature Vanadium Oxide Thin Films Obtained by Thermal Annealing of Layers Deposited by RF Magnetron Sputtering at Room Temperature**

Hernan M. R. Giannetta, Carlos Calaza, Liliana Fraigi and Luis Fonseca Hernan M. R. Giannetta, Carlos Calaza, Liliana Fraigi and Luis Fonseca

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

http://dx.doi.org/10.5772/67054

#### **Abstract**

This chapter describes a new deposition method proposed to achieve Vanadium Oxide VOx /V2 O5 thin films with high temperature coefficient of resistance (TCR), intended to be used as functional material in IR microsensors (bolometers). The main aim of the work is to attain a deposition method compatible with the lift-off microstructuring technique in order to avoid the use of a reactive-ion etching (RIE) process step to selectively remove the VOx /V2 O5 deposited layer in the course of the definition of the bolometer geometry, preventing the harmful effects linked to the spatial variability and the lack of selectivity of the RIE process. The proposed technique makes use of a two-stage process to produce the well-controlled VO<sup>x</sup> or V2 O5 thin films by applying a suitable thermal annealing to a previously deposited layer, which was obtained before at room temperature by RF magnetron sputtering and patterned by lift-off. A set of measurements has been carried out with thin films attained in order to check the quality and properties of the materials achieved with this method. The results reached with V<sup>2</sup> O5 pure phase films are consistent with a charge transport model based on the small polarons hopping derived from Mott's model under the Schnakenberg form.

**Keywords:** VOx thin film, V<sup>2</sup> O5 thin film, lift-off compatible, RF magnetron sputtering, thermal annealing, Meyer-Neldel rule, small polaron hopping

### **1. Introduction**

Thin films of materials with high temperature coefficient of resistance (TCR) values are widely used as thermoresistive transducers in uncooled infrared imaging sensors. Mixed Vanadium

© 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.

Oxide (VO<sup>x</sup> ) thin films were among the first functional materials chosen for this application due to its simple integration with MEMS technology, which led to the development of the first IR image sensors based on focal plane arrays (FPAs) with thermoresistive transducers in the 1970s–1980s [1]. The main reasons leading to the election of the VO<sup>x</sup> thin films as bolometer functional material were its high TCR value (for most commercial devices TCRs are in the range of −2 to –3%/K [2]) and its low 1/*f* noise. First microbolometers were achieved using polycrystalline VO<sup>x</sup> mixed oxide thin films, formed by a blend of VO<sup>2</sup> , V<sup>2</sup> O3 , and V<sup>2</sup> O5 oxide phases, with a thickness of 500–1000 Å, a resistivity of 20kΩ/sq. and a TCR (measured at room temperature) of −2%/K [3]. Since then, numerous works have been published presenting methods to improve the performance of these films and the associated devices. The best balance between the resistivity and the TCR reported for a commercial device was achieved using VO<sup>x</sup> thin films with a mixing V−O proportion equal to (*x* = 1.8) [4]. The TCR value of a specific material can be obtained experimentally by measuring the slope of the variation of the film resistivity with the temperature, which is described by the expression TCR = 1/R(dR/dT) [1]. But, as will be shown later, it can be also linked to a material parameter known as the material activation energy Δ*W*, given by TCR = −Δ*W*/*k*T2 [5], where *k* is the Boltzmann constant and *T* the temperature.

Two main methodologies able to control the mixing ratio x of the oxide phases present in the sample have been described in the literature for the experimental synthesis of VO<sup>x</sup> thin films. The first one achieves the control of the mixing proportion by using multilayers of pure phase materials. Among the published works based on this method, we can highlight Ref. [6] in which authors were able to control the mixing proportion with a sequence of successive depositions of two known phases, VO<sup>2</sup> and V2 O5 , or Ref. [7] in which authors controlled this ratio with a multilayer of V<sup>2</sup> O5 /V/V2 O5 . However, the most extended method in literature controls the film-mixing ratio by managing the oxidation rate of the material taken from a pure metallic precursor target (pure vanadium material, 99.9%) used for thin film growth. The deposition process parameters are tuned during the growth cycle to adjust the oxidation rate and, consequently, the ratio of the existing oxide phases that will determine the final thin film properties.

This chapter presents a novel technique to obtain vanadium oxide thin films (VO<sup>x</sup> or V2 O5 ) using different annealing conditions with starting materials previously deposited and patterned at room temperature. The key objective of the method is to offer compatibility with the lift-off microstructuring technique in order to allow the definition of the bolometer active material geometry without need of a dry etching process. The bolometer functional materials are usually deposited on top of thin dielectric membranes (Si3 N4 /SiO<sup>2</sup> ) required to accomplish an adequate thermal isolation with respect to the silicon substrate. The thickness and stress level of materials used in this membrane can be crucial for bolometer mechanical stability and a partial etch of these layers during a reactive-ion etching (RIE) step used to conform the active layer (due to an over-etch linked to a spatial variability or a lack of selectivity of the RIE process) can compromise future structure reliability. For that reason, the solution proposed splits the thin film formation into two stages: a first deposition of a precursor thin film on samples with a photoresist layer ready for the lift-off process, obtained from a VO<sup>2</sup> or a metallic V target using RF magnetron sputtering at room temperature; and a second annealing step at a high temperature after the photoresist removal, to promote the oxidation and obtain the desired oxide phase mix required for the application.

The morphological, structural, and optical characterization of the thin films obtained with this method under different annealing conditions were performed using field emission scanning electron microscopy (FE-SEM), X-ray diffraction (XRD), and Raman spectroscopy. The electrical conductivity of the samples was measured as a function of the temperature using a probe station attached to a heated chuck. The DC-conduction data measured for pure phase V2 O5 thin films was fitted using the Mott's small polaron hopping transport model, taking into account the Schnakenberg phonon distribution model equation. The consistency between the different parameters measured for the V<sup>2</sup> O5 samples processed with the optimal conditions and the assumptions of strong electron-phonon interaction, existence of small polarons, and the nonadiabatic regime for the hopping of charge carriers has been checked, suggesting that small polarons hopping is the main conduction mechanism in pure phase V<sup>2</sup> O5 thin films obtained with this method.

### **2. Vanadium oxides**

Oxide (VO<sup>x</sup>

crystalline VO<sup>x</sup>

given by TCR = −Δ*W*/*k*T2

with a multilayer of V<sup>2</sup>

sitions of two known phases, VO<sup>2</sup>

O5 /V/V2 O5

are usually deposited on top of thin dielectric membranes (Si3

desired oxide phase mix required for the application.

) thin films were among the first functional materials chosen for this application

, V<sup>2</sup> O3

, or Ref. [7] in which authors controlled this ratio

. However, the most extended method in literature controls

N4 /SiO<sup>2</sup> , and V<sup>2</sup>

O5

thin films as bolometer

oxide phases,

thin films

thin films.

 or V2 O5 )

or a metal-

) required to accomplish

due to its simple integration with MEMS technology, which led to the development of the first IR image sensors based on focal plane arrays (FPAs) with thermoresistive transducers in the

functional material were its high TCR value (for most commercial devices TCRs are in the range of −2 to –3%/K [2]) and its low 1/*f* noise. First microbolometers were achieved using poly-

with a thickness of 500–1000 Å, a resistivity of 20kΩ/sq. and a TCR (measured at room temperature) of −2%/K [3]. Since then, numerous works have been published presenting methods to improve the performance of these films and the associated devices. The best balance between

with a mixing V−O proportion equal to (*x* = 1.8) [4]. The TCR value of a specific material can be obtained experimentally by measuring the slope of the variation of the film resistivity with the temperature, which is described by the expression TCR = 1/R(dR/dT) [1]. But, as will be shown later, it can be also linked to a material parameter known as the material activation energy Δ*W*,

Two main methodologies able to control the mixing ratio x of the oxide phases present in the

The first one achieves the control of the mixing proportion by using multilayers of pure phase materials. Among the published works based on this method, we can highlight Ref. [6] in which authors were able to control the mixing proportion with a sequence of successive depo-

the film-mixing ratio by managing the oxidation rate of the material taken from a pure metallic precursor target (pure vanadium material, 99.9%) used for thin film growth. The deposition process parameters are tuned during the growth cycle to adjust the oxidation rate and, consequently, the ratio of the existing oxide phases that will determine the final thin film properties.

using different annealing conditions with starting materials previously deposited and patterned at room temperature. The key objective of the method is to offer compatibility with the lift-off microstructuring technique in order to allow the definition of the bolometer active material geometry without need of a dry etching process. The bolometer functional materials

an adequate thermal isolation with respect to the silicon substrate. The thickness and stress level of materials used in this membrane can be crucial for bolometer mechanical stability and a partial etch of these layers during a reactive-ion etching (RIE) step used to conform the active layer (due to an over-etch linked to a spatial variability or a lack of selectivity of the RIE process) can compromise future structure reliability. For that reason, the solution proposed splits the thin film formation into two stages: a first deposition of a precursor thin film on

lic V target using RF magnetron sputtering at room temperature; and a second annealing step at a high temperature after the photoresist removal, to promote the oxidation and obtain the

[5], where *k* is the Boltzmann constant and *T* the temperature.

1970s–1980s [1]. The main reasons leading to the election of the VO<sup>x</sup>

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

mixed oxide thin films, formed by a blend of VO<sup>2</sup>

the resistivity and the TCR reported for a commercial device was achieved using VO<sup>x</sup>

sample have been described in the literature for the experimental synthesis of VO<sup>x</sup>

O5

This chapter presents a novel technique to obtain vanadium oxide thin films (VO<sup>x</sup>

samples with a photoresist layer ready for the lift-off process, obtained from a VO<sup>2</sup>

and V2

Vanadium is a transition metal with a [Ar] 3d3 4s2 electron configuration for the ground state and a centered cubic crystal structure [8]. As a consequence of its multivalent character, it has a number of possible oxidation states (V+2, V+3, V+4, V+5), which form an extensive list of binary V−O systems. Some of them are grouped in the so-called "Magneli phases," with stoichiometric formula VnO2n−1, and others in the Wadsley phases, with stoichiometric formula VnO2n+1.

The most commonly used phases, found in various applications due to their particular properties, are the VO, VO<sup>2</sup> , V<sup>2</sup> O3 , and V<sup>2</sup> O5 oxide phases. Their main characteristics are the following:

**VO** is one of the many vanadium oxide phases with crystalline cubic structure and good electrical conductivity due to the partially filled conduction band and the delocalization of electrons in the 2 g orbital [9].

**VO2** is an amphoteric compound with the unique property of changing from a semiconductor monoclinic phase to a (semi)metal tetragonal rutile phase at a temperature around 340 K and, therefore, its electrical resistivity together with the optical properties also change up to several orders of magnitude between these two states [10].

**V2 O3** phase, like VO<sup>2</sup> compound, presents an abrupt conductivity change at a temperature around 160 K, evidenced of a metal-insulator transition. In addition, it presents a thermochromic behavior in the infrared band [9].

**V2 O5** is the most stable of all vanadium oxide phases, and the preferred one to be used as thermoresistive material in microbolometer arrays for thermal imaging due to its high TCR value. Vanadium pentoxide is a semiconductor with a bandgap of 2.1–2.4 eV, which presents the following polymorphs: α-V<sup>2</sup> O5 (orthorhombic), β-V<sup>2</sup> O5 (monoclinic or tetragonal), and γ-V<sup>2</sup> O5 (orthorhombic), being the α-polymorph the most stable one [11].

There are many other V−O binary compounds with unique properties beyond the most used ones. A complete guide to the various V=O phases can be found in Ref. [9], including a diagram that represents the different V−O oxide phases as a function of their oxygen atomic fraction, obtained by thermodynamic calculations.

#### **3. Synthesis of VOx films**

A number of deposition techniques can be highlighted among the different synthesis methods reported in the literature for the synthesis of a multiplicity of Vanadium oxide compounds: sputtering [12], sol-gel process [13], chemical vapor deposition (CVD) [14], pulse laser deposition (PLD) [15], atomic layer deposition (ALD) [9], molecular beam epitaxy (MBE) [16], aqueous solution process [17], and the reactive vacuum evaporation [6].

Nevertheless, to obtain a specific oxide phase, it is necessary to complement one of these deposition methods with a thermal annealing process in order to enhance the crystallinity of the film, as well as to modify its stoichiometry [18]. Two illustrative cases can be found in Refs. [19, 20], where VO<sup>x</sup> samples were subjected to a thermal annealing process in air at different temperatures (up to 300°C) and a film recrystallization was detected. Furthermore, it was observed that an increase of the substrate temperature after film deposition promotes the loss of oxygen atoms, which results in a modification of the VO<sup>x</sup> film stoichiometry [12].

The temperature conditions required to achieve a particular oxide phase will depend on the enthalpy of formation of the various vanadium oxide compounds [19]. On the basis of the change in the Gibbs free energy associated with each stable Vanadium oxide phase, it can be established that V<sup>2</sup> O5 formation requires temperatures above 434°C, and the stable oxide phases formation sequence will be given in the order: VO<sup>2</sup> → V2 O<sup>3</sup> → VO → V2 O5 [21].

Consequently, it has been necessary to explore the influence of the temperature and the annealing time conditions used in the second step of the proposed thin film growth method in order to establish the optimal set of parameters that lead to attain the desired vanadium oxide phase combination.

#### **3.1. Preparation and synthesis of VOx films**

The sputtering technique, one of the most common physical vapor deposition methods [22], has been the process selected for the growth of vanadium oxide thin films. The key advantages of the sputtering technique over other alternative methods include the resultant film uniformity, the easy scalability to larger substrates, and the high deposition efficiency [23]. There are three main sputtering deposition modes available: DC, RF, and magnetron. The first reported growth of a VO<sup>2</sup> thin film deposited with a reactive sputtering method is typically attributed to Fuls and collaborators, from Bell Telephone Labs in 1967 [24].

A wide range of sputtering process conditions have been reported in the literature for the growth of VO<sup>x</sup> thin films with different compositions. The main process parameters that can be customized to control the final composition of the resultant material include the selection of a starting target with a virtually pure precursor material, such as metallic-V [25], VO2 [12], V<sup>2</sup> O5 [11]; the choice of a specific sputtering deposition method, RF [25] or DC [26]; the alteration of the substrate temperature during the deposition process [27]; the subsequent annealing in vacuum or in an oxidizing atmosphere with a controlled O<sup>2</sup> /N<sup>2</sup> ratio [28]; or the application of a varying voltage to the substrate during the deposition [29].

that represents the different V−O oxide phases as a function of their oxygen atomic fraction,

A number of deposition techniques can be highlighted among the different synthesis methods reported in the literature for the synthesis of a multiplicity of Vanadium oxide compounds: sputtering [12], sol-gel process [13], chemical vapor deposition (CVD) [14], pulse laser deposition (PLD) [15], atomic layer deposition (ALD) [9], molecular beam epitaxy (MBE) [16], aque-

Nevertheless, to obtain a specific oxide phase, it is necessary to complement one of these deposition methods with a thermal annealing process in order to enhance the crystallinity of the film, as well as to modify its stoichiometry [18]. Two illustrative cases can be found

different temperatures (up to 300°C) and a film recrystallization was detected. Furthermore, it was observed that an increase of the substrate temperature after film deposition promotes

The temperature conditions required to achieve a particular oxide phase will depend on the enthalpy of formation of the various vanadium oxide compounds [19]. On the basis of the change in the Gibbs free energy associated with each stable Vanadium oxide phase, it can

Consequently, it has been necessary to explore the influence of the temperature and the annealing time conditions used in the second step of the proposed thin film growth method in order to establish the optimal set of parameters that lead to attain the desired vanadium

The sputtering technique, one of the most common physical vapor deposition methods [22], has been the process selected for the growth of vanadium oxide thin films. The key advantages of the sputtering technique over other alternative methods include the resultant film uniformity, the easy scalability to larger substrates, and the high deposition efficiency [23]. There are three main sputtering deposition modes available: DC, RF, and magnetron. The first

A wide range of sputtering process conditions have been reported in the literature for the

be customized to control the final composition of the resultant material include the selection of a starting target with a virtually pure precursor material, such as metallic-V [25],

 **films**

attributed to Fuls and collaborators, from Bell Telephone Labs in 1967 [24].

samples were subjected to a thermal annealing process in air at

formation requires temperatures above 434°C, and the stable oxide

thin film deposited with a reactive sputtering method is typically

thin films with different compositions. The main process parameters that can

[11]; the choice of a specific sputtering deposition method, RF [25] or DC [26];

film stoichiometry [12].

O5 [21].

O<sup>3</sup> → VO → V2

obtained by thermodynamic calculations.

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

 **films**

ous solution process [17], and the reactive vacuum evaporation [6].

the loss of oxygen atoms, which results in a modification of the VO<sup>x</sup>

phases formation sequence will be given in the order: VO<sup>2</sup> → V2

**3. Synthesis of VOx**

in Refs. [19, 20], where VO<sup>x</sup>

be established that V<sup>2</sup>

oxide phase combination.

reported growth of a VO<sup>2</sup>

O5

growth of VO<sup>x</sup>

[12], V<sup>2</sup>

VO2

**3.1. Preparation and synthesis of VOx**

O5

In this work, the conditions explored to achieve the desired vanadium oxide thin films include the use of two different sputtering target materials: VO<sup>2</sup> and metallic-V. The thin films obtained with the sputtering process have been patterned using the lift-off technique, and finally, the samples have gone through a thermal annealing step to obtain the desired film composition. The experimental conditions used for this final annealing step in the case of the metallic-V target have been investigated in order to establish the set of parameters that are required to obtain VO<sup>x</sup> and V2 O5 thin films.

The first step in sample preparation has been the arrangement of a base substrate on which the different layers of vanadium oxide compounds have been deposited. The substrate used has been a Si wafer (100) single-side polished, with a precoating of SiO<sup>2</sup> and Si3 N4 , which has acted as an electrical insulator and as a mechanical buffer layer for subsequent thermal annealing. The surface was subjected to a standard cleanroom cleaning process previously to the deposition.

In order to confirm the compatibility of the proposed growth method with the lift-off microstructuring technique, the vanadium oxide thin films have been deposited on top of substrates partially covered with a photoresist layer. The preparation of the samples starts with the deposition of a layer of TI35E image reversal photoresist (MicroChemicals GmbH) on top of the Si substrates, which is selectively exposed using a photolithographic mask taking advantage of the image reversal feature of the photoresist to obtain negative sidewall profiles adequate for the liftoff process. The photoresist has been finally removed in exposed areas using an AZ developer.

The deposition of the Vanadium oxide thin films has been accomplished by using a BOC Edwards auto 500 Sputtering system, which was equipped with both RF and DC sources, using a maximum substrate temperature of 80°C to avoid damaging the photoresist. A set of vanadium oxide thin film depositions has been conducted under two different conditions.

A first set of samples (labelled as type A) has been obtained from a VO<sup>2</sup> target (brand MCSE, 99.9% purity, 3″ diameter, and 6 mm thick) with the RF source, using a 20 sccm argon flow, a 200 W RF power, and 30 min of deposition time. During deposition, the partial pressure on the chamber was 3.31 × 10−3 mBar, and sample holder was maintained at room temperature.

A second set of samples (labelled as type B) has been obtained from a metallic-V target (brand Kurt J. Lesker, 99.5% purity, 3″ diameter, and 6 mm thick) with the DC source, using a 30 sccm argon and 10 sccm oxygen flow, a 400 W DC power, and 30 min deposition time. During deposition, the partial pressure on the chamber was 6.61 × 10−3 mBar, and sample holder was maintained at 80°C.

Sample preparation was concluded with the lift-off process, which removes the photoresist and the material deposited on top of it with an acetone solvent. At this point, films are ready to endure the thermal treatment required to modify the crystallography and stoichiometry.

Type A samples were subjected to annealing processes using different temperatures: 80, 280, 400, and 475°C and times 8, 6, 4, and 3 h, respectively, in an air atmosphere making use of an electrical oven with a temperature control. In contrast, with type B samples, it was essential to promote the additional oxidation of the metallic-V and, therefore, all samples were subjected to an annealing process at a higher temperature of 500°C for 30 min, making use of a vacuum oven with a controlled argon/oxygen environment (1 sccm oxygen and 3 sccm argon flow).

After the thermal annealing of the samples, the electrical contacts with the vanadium oxide films were obtained making use of an aluminium metallization in a subsequent process step. The samples resulting from this process sequence are illustrated in **Figure 1**.

**Figure 1.** VOx film conformation by lift-off and metal interconnections: (a) VO<sup>x</sup> after lift-off, (b) VO<sup>x</sup> with Al metal connections.

#### **4. Characterization of VOx films**

The morphological characterization of the samples has been performed using a field emission scanning electron microscope (FE-SEM, FEI Helios Nanolab 650) with a low acceleration potential, 3 kV. The morphology of the VO<sup>x</sup> thin films that result from the proposed growth method is shown in **Figure 2**, with FE-SEM images corresponding to the samples A and B annealed at the highest temperatures, 475 and 500°C, respectively. A good film uniformity can be observed in both the cases as well as in an enhancement of the crystal grain size and shape due to the effect of the thermal annealing [30]. The measured grain sizes are between 100 and 250 nm in both the cases, values that are similar to those reported in the literature for similar materials [31].

The structural characterization was performed by X-ray diffraction (XRD), using a CuKa cathode = 1.5406 Å (XRD-Philips PW 1730/10) configured in a standard Bragg-Brentano powder diffraction geometry. The analysis of the peak coincidences for each sample was performed using the PDF2 database with license from the International Center for Diffraction Data (ICDD). The crystalline phases present in the different samples were determined after the identification of the peak coincidences in the XRD patterns. The most significant results have been obtained for samples annealed at highest temperatures, A@475 and B@500, due to the presence of the dominant V<sup>2</sup> O5 phase, as shown in **Figure 3**.

Vanadium Oxide Thin Films Obtained by Thermal Annealing of Layers Deposited by RF Magnetron Sputtering... http://dx.doi.org/10.5772/67054 157

**Figure 2.** Microscopy Image performed by FE-SEM: (a) sample A@475°C, (b) sample B@500°C.

an electrical oven with a temperature control. In contrast, with type B samples, it was essential to promote the additional oxidation of the metallic-V and, therefore, all samples were subjected to an annealing process at a higher temperature of 500°C for 30 min, making use of a vacuum oven with a controlled argon/oxygen environment (1 sccm oxygen and 3 sccm

After the thermal annealing of the samples, the electrical contacts with the vanadium oxide films were obtained making use of an aluminium metallization in a subsequent process step.

The samples resulting from this process sequence are illustrated in **Figure 1**.

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

 **films**

film conformation by lift-off and metal interconnections: (a) VO<sup>x</sup>

The morphological characterization of the samples has been performed using a field emission scanning electron microscope (FE-SEM, FEI Helios Nanolab 650) with a low acceleration

method is shown in **Figure 2**, with FE-SEM images corresponding to the samples A and B annealed at the highest temperatures, 475 and 500°C, respectively. A good film uniformity can be observed in both the cases as well as in an enhancement of the crystal grain size and shape due to the effect of the thermal annealing [30]. The measured grain sizes are between 100 and 250 nm in both the cases, values that are similar to those reported in the literature for

The structural characterization was performed by X-ray diffraction (XRD), using a CuKa cathode = 1.5406 Å (XRD-Philips PW 1730/10) configured in a standard Bragg-Brentano powder diffraction geometry. The analysis of the peak coincidences for each sample was performed using the PDF2 database with license from the International Center for Diffraction Data (ICDD). The crystalline phases present in the different samples were determined after the identification of the peak coincidences in the XRD patterns. The most significant results have been obtained for samples annealed at highest temperatures, A@475 and B@500, due to the

phase, as shown in **Figure 3**.

thin films that result from the proposed growth

after lift-off, (b) VO<sup>x</sup>

with Al metal

argon flow).

**Figure 1.** VOx

connections.

**4. Characterization of VOx**

similar materials [31].

presence of the dominant V<sup>2</sup>

O5

potential, 3 kV. The morphology of the VO<sup>x</sup>

**Figure 3.** XDR peak matching analysis: (a) sample A@475°C, (b) sample B@500°C.

The analysis of the diffractogram corresponding to A@475 sample shows the presence of a practically pure V<sup>2</sup> O5 oxide phase. Peaks identified correspond to V<sup>2</sup> O5 reflection planes V<sup>2</sup> O5 (010) = 20.25°, V<sup>2</sup> O5 (110) = 21.72°, and V<sup>2</sup> O5 (020) = 41.22°.

In contrast, the presence of a mixed oxide phase can be observed on the diffractogram corresponding to B@500 sample, where main peaks can be attributed to a V<sup>2</sup> O5 phase and secondary ones to a V<sup>6</sup> O13 phase. The reflection planes identified are V<sup>2</sup> O5 (010) = 20.25°, V2 O5 (110) = 21.72°, V<sup>2</sup> O5 (111) = 33.27°, V<sup>2</sup> O5 (020) = 41.42°, and V<sup>6</sup> O13 (002) = 17.77°, V<sup>6</sup> O<sup>13</sup> (003) = 27.08°, V<sup>6</sup> O13 (005) = 45.78,° and V<sup>6</sup> O13 (006) = 55.53° for the secondary V<sup>6</sup> O13 phase.

Optical vibrational modes have been analyzed by means of Raman spectroscopy using a LabRAM HR Raman system (Horiba Jobin Yvon) equipped with a confocal microscope and a charge coupled device (CCD) detector. The 514.5 nm emission line of an Ar + laser has been used as an excitation source. The material databases for peak identification were included in the instrument. The prevailing V<sup>2</sup> O5 phase is once more clearly identified in Raman spectrums obtained for A@475 and B@500 samples due to the perfect matching of the spectrum peaks with data extracted from the database. **Figure 4** displays the clear correspondence between the measured spectrum for both the samples and the one relative to the V<sup>2</sup> O5 reference material in the database.

The presence of two main emission peaks can be clearly spotted in the Raman spectrum of both samples; a first one in 141.6 cm−1, which corresponds to the V–O–V Raman vibration mode and a second one in 993 cm−1, associated with the V=O double-bond vibration mode [32]. In contrast, the strong emission peak observed in the Raman spectrum of A@475 sample (**Figure 4a**), in 520 cm−1, cannot be assigned to any V–O vibration mode and has been related to the silicon crystalline substrate used as a support for the samples.

**Figure 4.** Raman characterization: (a) sample A@475°C, (b) sample B@500°C.

Finally, the electrical conductivity of the Vanadium oxide films has been measured as a function of the sample temperature, using a semiconductor parameter analyser Keithley 4200 CSC connected to a probe station and a thermal chuck equipped with an electronic temperature controller. Resistance measurements were performed using four collinear probes to implement a four-wire measurement scheme. A temperature swept, from room temperature up to 100°C, was applied to the chuck, and the actual temperature achieved by samples was recorded with the help of an additional thermocouple in contact with sample surface.

The measurement of the electrical conductivity as a function of the sample temperature has been used to establish the activation energy corresponding to the different materials by fitting the experimental data with Eq. (1):

$$
\sigma\_{(l)} = \sigma\_0 \exp\left(\frac{-\Delta W}{kT}\right),
\tag{1}
$$

where *σ*<sup>0</sup> is a constant, Δ*W* is the activation energy, *k* is the Boltzmann constant expressed in eV/*k*, and *T* is the temperature in Kelvin. The activation energy was obtained by fitting the experimental data obtained for the electrical conductivity vs. temperature (log(*σ*) vs. 1000/*T*) to a linear expression derived from Eq. (1) using the least squares method (**Figure 5**). Such linear fits over the experimental measurements provide activation energy values of Δ*W* = 0.267 eV for the A@475 sample, and Δ*W* = 0.056 eV for the B@500 sample.

**Figure 5.** Electrical characterization of samples: (a) sample A@475°C, (b) sample B@500°C.

These activation energy values corroborate that A@475 sample presents a single V<sup>2</sup> O5 phase, as it presents an activation energy similar to that reported by Ioffe [33] for the V<sup>2</sup> O5 singlecrystal material with a purity of 99%, Δ*W* = 0.27 eV. In contrast, the activation energy that has been measured for the B@500 sample is much lower due to the presence of other mixed vanadium oxide phases. Consequently, the TCR value will be higher for the A@475 sample than for the B@500 sample, but the higher value of electrical conductivity in B type samples can represent a benefit in terms of the 1/*f* noise. The TCR estimated for each sample from the measured activation energy using the formerly stated equality TCR = −Δ*W*/*k*T2 [5] is TCRA@475 = 3.44%/K and TCRB@500 = 0.72%/K, respectively.

### **5. Electrical charge transport**

with data extracted from the database. **Figure 4** displays the clear correspondence between

The presence of two main emission peaks can be clearly spotted in the Raman spectrum of both samples; a first one in 141.6 cm−1, which corresponds to the V–O–V Raman vibration mode and a second one in 993 cm−1, associated with the V=O double-bond vibration mode [32]. In contrast, the strong emission peak observed in the Raman spectrum of A@475 sample (**Figure 4a**), in 520 cm−1, cannot be assigned to any V–O vibration mode and has been

Finally, the electrical conductivity of the Vanadium oxide films has been measured as a function of the sample temperature, using a semiconductor parameter analyser Keithley 4200 CSC connected to a probe station and a thermal chuck equipped with an electronic temperature controller. Resistance measurements were performed using four collinear probes to implement a four-wire measurement scheme. A temperature swept, from room temperature up to 100°C, was applied to the chuck, and the actual temperature achieved by samples was

recorded with the help of an additional thermocouple in contact with sample surface.

the experimental data with Eq. (1):

where *σ*<sup>0</sup>

*σ*(*T*) = *σ*<sup>0</sup> exp(

**Figure 4.** Raman characterization: (a) sample A@475°C, (b) sample B@500°C.

The measurement of the electrical conductivity as a function of the sample temperature has been used to establish the activation energy corresponding to the different materials by fitting

\_\_\_\_ -Δ*W*

is a constant, Δ*W* is the activation energy, *k* is the Boltzmann constant expressed in

eV/*k*, and *T* is the temperature in Kelvin. The activation energy was obtained by fitting the experimental data obtained for the electrical conductivity vs. temperature (log(*σ*) vs. 1000/*T*) to

*<sup>κ</sup><sup>T</sup>* ), (1)

O5

reference mate-

the measured spectrum for both the samples and the one relative to the V<sup>2</sup>

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

related to the silicon crystalline substrate used as a support for the samples.

rial in the database.

The different morphological and structural analysis performed with samples A@475 and B@500 has shown that V<sup>2</sup> O5 is the prevailing vanadium oxide phase formed in all samples due to the use of an annealing temperature well above 434°C, but a slightly different behavior has been perceived in electrical conductivity measurements. The morphological tests determine that A@475 sample presents an almost pure V<sup>2</sup> O5 phase, while a set of mixed vanadium oxide phases VOx has been obtained for all B type samples, even those that have been processed at the highest temperature, 500°C. Nevertheless, even for these B type samples, the clear preponderance of the V<sup>2</sup> O5 phase when compared with other identified oxide phases, as can be seen on the Raman spectrums in **Figure 4**, should be noted.

The electrical behavior measured for the different vanadium oxide samples, illustrated in **Figure 6a**, shows that electrical conductivity (*σ*) has an Arrhenius behavior as a function of temperature (*T*), indicating that electrical conduction in this material is the result of a thermally activated process. A Meyer-Neldel rule (MNR) relationship is found to hold in a wide variety of such activated processes as the electron conduction in extended states, the ionic conduction, or the thermally activated hopping [34–36]. According to the MNR, the prefactor in Eq. (1) (*σ*<sup>0</sup> ) and the activation energy (Δ*W*) are related by Eq. (2), where *σ*00 is a constant and Δ*WMNR* is the Meyer-Neldel activation energy

$$
\log \langle \sigma\_o \rangle = \log \langle \sigma\_{o0} \rangle + \frac{\Delta W}{\Delta W\_{\text{MOR}}}.\tag{2}
$$

The activation energies for the processed samples range from 0.093 to 0.217 eV, according to the experimental data obtained for the electrical conductivity, measured in a temperature range between 25 and 100°C. As can be seen in **Figure 6b**, all of them obey the MNR, with a value of *σ*<sup>00</sup> = 4.2 × 10−3 Ω−1 cm−1 and Δ*WMNR* = 0.9374 eV, with a goodness of fit *R*<sup>2</sup> = 0.8944.

**Figure 6.** Compliance of activation energy for different processing temperatures with Meyer-Neldel rule: (a) log(σ) vs. 1000/T for samples A and B, (b) fit of ΔW to Meyer-Neldel rule.

The good correspondence observed between the activation energies derived from the electrical conductivity measurements carried out with the different samples and the conventional Meyer-Neldel rule can be attributed to a modification of the location of the Fermi energy level *εF* with respect to the conduction band *εC* with the different annealing conditions applied, similar to the behavior noted in amorphous semiconductors [37].

#### **5.1. Electrical transport in V<sup>2</sup> O5**

It is well known that the presence of the V<sup>2</sup> O5 oxide phase is responsible for the high TCR value observed at room temperature in mixed VO<sup>x</sup> samples. Small polaron hopping between localized states is the prevailing mechanism that handles the electrical charge transport in the V2 O5 phase. Mott's works [34, 38] established the basis for the study of this transport mechanisms in transition metals [39], and the proposed models were successively used to experimentally identify this behavior in both single crystal [33] and amorphous V2 O5 [40] materials. In this section, Mott's models are used to confirm that experimental measurements obtained with sample A@475, with an almost pure V<sup>2</sup> O5 phase, are consistent with the small polaron hopping charge transport mechanism.

temperature (*T*), indicating that electrical conduction in this material is the result of a thermally activated process. A Meyer-Neldel rule (MNR) relationship is found to hold in a wide variety of such activated processes as the electron conduction in extended states, the ionic conduction, or the thermally activated hopping [34–36]. According to the MNR, the prefactor

The activation energies for the processed samples range from 0.093 to 0.217 eV, according to the experimental data obtained for the electrical conductivity, measured in a temperature range between 25 and 100°C. As can be seen in **Figure 6b**, all of them obey the MNR, with a value of *σ*<sup>00</sup> = 4.2 × 10−3 Ω−1 cm−1 and Δ*WMNR* = 0.9374 eV, with a goodness of fit *R*<sup>2</sup> = 0.8944.

The good correspondence observed between the activation energies derived from the electrical conductivity measurements carried out with the different samples and the conventional Meyer-Neldel rule can be attributed to a modification of the location of the Fermi energy level

**Figure 6.** Compliance of activation energy for different processing temperatures with Meyer-Neldel rule: (a) log(σ) vs.

with respect to the conduction band *εC* with the different annealing conditions applied,

O5

localized states is the prevailing mechanism that handles the electrical charge transport in the

In this section, Mott's models are used to confirm that experimental measurements obtained

 phase. Mott's works [34, 38] established the basis for the study of this transport mechanisms in transition metals [39], and the proposed models were successively used to experi-

oxide phase is responsible for the high TCR

samples. Small polaron hopping between

O5

[40] materials.

similar to the behavior noted in amorphous semiconductors [37].

mentally identify this behavior in both single crystal [33] and amorphous V2

**O5**

) and the activation energy (Δ*W*) are related by Eq. (2), where *σ*00 is a constant and

Δ*WMNR*

. (2)

in Eq. (1) (*σ*<sup>0</sup>

*εF*

V2 O5

**5.1. Electrical transport in V<sup>2</sup>**

It is well known that the presence of the V<sup>2</sup>

value observed at room temperature in mixed VO<sup>x</sup>

1000/T for samples A and B, (b) fit of ΔW to Meyer-Neldel rule.

Δ*WMNR* is the Meyer-Neldel activation energy

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

log(*σ*0) <sup>=</sup> log(*σ*00) <sup>+</sup> \_\_\_\_\_\_ <sup>Δ</sup>*<sup>W</sup>*

Charge transport through polarons is a well-known effect observed under conditions of strong electron-phonon interaction. One evident effect of strong electron-phonon interaction in a material is the dependence of the electrical conductivity with temperature as seen in analyzed samples. However, other less obvious effect is revealed under this condition, an increase in the effective mass of the electrons due to the interaction with heavy ion nuclei [41]. The assembly formed by the electron and its associated field deformation is known as polaron.

Polarons can be classified taking into account the size of the field deformation radius with respect to the lattice constant, which gives rise to large or small polarons. Charge transport mechanism depends on polaron size; while for large polarons, charges are moved in a unique band, with small polarons, the charge remains trapped on a single ion most of the time. The interaction between the lattice vibration and the localized electron induces charges to jump from one atom to a neighboring one. This process is called conduction by hopping charge carriers and takes place through thermal activation at high temperature [42].

Based on Mott's model, the DC conductivity for the hopping of polarons in a nonadiabatic approximation, above the Debye temperature *θD* is given by Eq. (3) [34, 38]:

$$
\sigma = \nu\_0 N e^2 R^2 \frac{C(1 \cdot C)}{\kappa T} \exp(-2\alpha \mathbb{R}) \exp\left(\frac{-\Delta W}{kT}\right),
\tag{3}
$$

where *ν*<sup>0</sup> is the optical phonon frequency, *N* is the number of transition metal ion sites per unit volume, *e* is the electron charge, *α* is the wave function decay constant, *C* is the ratio of ion concentration (V+4 vs. V+5), *R* is the hopping distance, and *θD* is the Debye temperature given by *θ<sup>D</sup>* = *hν*<sup>0</sup> /k, where *h* and *k* are the Plank and Boltzmann constants. Schnakenberg proposed a simplified model formulation taking into account the phonon distribution [43] in which the dependence of the electrical conductivity with the temperature can be expressed with Eq. (4) [44],

$$\ln\left(\sigma T\right) = \ln\left(\sigma T\right)\_0 \cdot \frac{W\_D}{2kT} \cdot \frac{W\_H}{kT} \frac{\tanh\left(h\,\nu\_0/4kT\right)}{h\,\nu\_0/4kT} \,, \tag{4}$$

where *σ*<sup>0</sup> is a constant, *WD* is the activation energy for hopping due to disorder, *WH* is the polaron hopping energy, and *T* the temperature in *K*. Eq. (4) is valid for the hopping of polarons in the nonadiabatic approximation, above *θD*/2 temperature, where *θD* is the Debye temperature. Total activation energy ∆*W* can be derived from previous parameters using an expression proposed by Austin-Mott in Eq. (5) [34, 38],

$$\begin{cases} \Delta W = W\_{\rm H} + \frac{W\_{\rm D}}{2} & \text{for } T > \Theta\_{\rm D}/2, \\ \Delta W = W\_{\rm D} & \text{for } T < \Theta\_{\rm D}/4. \end{cases} \tag{5}$$

In order to verify the nature of the electronic transport mechanism in sample A@475, with an almost pure V<sup>2</sup> O5 phase, the experimental data obtained for the electrical conductivity was fitted by using the Mott-Schnakenberg model, as shown in our previous work [45]. The fitting parameters derived by least squares for *σ*<sup>0</sup> , *WH*, *WD*, and *hν*<sup>0</sup> have been used to check the condition for polaron existence and to identify the particular type of polaron that is responsible of charge transport. The parameters obtained from the best linear fit, with a 95% confidence interval limits, were *WH* = 0.1682 ± 0.0121 eV, *WD* = 0.2241 ± 0.0139 eV, and *hν*<sup>0</sup> = 0.02755 ± 0.00994 eV; with a goodness of fit R<sup>2</sup> = 0.9827 [45]. The Debye temperature *θ<sup>D</sup>* = *hν*<sup>0</sup> /k that corresponds to these values is *θ<sup>D</sup>* = 319 K.

#### **5.2. Verification of the polaron charge transport in V<sup>2</sup> O5**

The values extracted for Mott's model parameters with the fit of the electrical conductivity data have been used to check the consistency of the results with the several hypotheses used by Schnakenberg in Mott's model simplified formulation, which involve significant assumptions regarding the transport mechanism. Specifically, we have certified that the strong electronphonon interaction condition proposed by Austin-Mott, the small polaron formation condition established by Emin-Holstein, the minimum mobility condition for conduction by hopping formulated by Cohen, and the condition proposed by Holstein to set the limit between the adiabatic and nonadiabatic regimes are verified for the temperature range used in measurements.

The strong electron-phonon interaction condition was verified by calculating the ratio of the polaron effective mass *mP* to the rigid-lattice effective mass *m\**. The higher this ratio, the greater the electron-phonon coupling. The ratio between these two parameters was obtained using Eq. (6) [34, 38],

$$m\_p = \left(\hbar^2 \middle| \!\!\!/\!\!\!/\!\!/\!\!/^2\right) \exp\left(\begin{smallmatrix} \gamma \end{smallmatrix}\right) = m^\* \exp\left(\begin{smallmatrix} \gamma \end{smallmatrix}\right) \tag{6}$$

where *J* is the polaron bandwidth, *R* is the mean separation between the transition metal ions, and *γ* is the electron-phonon interaction parameter, which is given by Eq. (7) [34, 38],

$$\gamma \,\, \gamma \,\, = \, 2 \left( \frac{\mathcal{W}\_{\mathcal{H}}}{h \,\, \nu\_{\boldsymbol{\alpha}}} \right) . \tag{7}$$

The condition of small polaron formation was verified using the expression proposed by Emin-Holstein in Ref. [46], an inequality that relates the polaron hopping energy *WH* with the bandwidth of the polaron *J*, as expressed in Eq. (8),

$$J \ll \frac{\mathcal{W}\_{\rm H}}{3}.\tag{8}$$

The bandwidth of the polaron *J* was estimated from the fitting parameters using the approximate expression proposed by Holstein [47], as shown in the relationship in Eq. (9),

$$J \approx 0.67 \,\text{h} \,\text{v}\_{\text{o}} \,\text{(T/} \theta\_{\text{o}}\text{)}^{\text{u}}.\tag{9}$$

The type of conduction mechanisms followed by the charge carriers was established using the condition proposed by Holstein [47] to set the limit between the adiabatic and nonadiabatic regimes [34], given by Eq. (10),

Vanadium Oxide Thin Films Obtained by Thermal Annealing of Layers Deposited by RF Magnetron Sputtering... http://dx.doi.org/10.5772/67054 163

$$H = \left< 2kT \, \mathcal{W}\_{\mathbb{H}} | \pi \right>^{14} \left< \hbar \nu\_a | \pi \right>^{12}. \tag{10}$$

The polaron conduction regime is determined by the inequality,

The fitting parameters derived by least squares for *σ*<sup>0</sup>

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

**5.2. Verification of the polaron charge transport in V<sup>2</sup>**

/k that corresponds to these values is *θ<sup>D</sup>* = 319 K.

*θ<sup>D</sup>* = *hν*<sup>0</sup>

using Eq. (6) [34, 38],

*mP* = (ℏ<sup>2</sup>

bandwidth of the polaron *J*, as expressed in Eq. (8),

*J* < \_\_\_

regimes [34], given by Eq. (10),

, *WH*, *WD*,

**O5**

check the condition for polaron existence and to identify the particular type of polaron that is responsible of charge transport. The parameters obtained from the best linear fit, with a 95% confidence interval limits, were *WH* = 0.1682 ± 0.0121 eV, *WD* = 0.2241 ± 0.0139 eV, and *hν*<sup>0</sup> = 0.02755 ± 0.00994 eV; with a goodness of fit R<sup>2</sup> = 0.9827 [45]. The Debye temperature

The values extracted for Mott's model parameters with the fit of the electrical conductivity data have been used to check the consistency of the results with the several hypotheses used by Schnakenberg in Mott's model simplified formulation, which involve significant assumptions regarding the transport mechanism. Specifically, we have certified that the strong electronphonon interaction condition proposed by Austin-Mott, the small polaron formation condition established by Emin-Holstein, the minimum mobility condition for conduction by hopping formulated by Cohen, and the condition proposed by Holstein to set the limit between the adiabatic and nonadiabatic regimes are verified for the temperature range used in measurements. The strong electron-phonon interaction condition was verified by calculating the ratio of the polaron effective mass *mP* to the rigid-lattice effective mass *m\**. The higher this ratio, the greater the electron-phonon coupling. The ratio between these two parameters was obtained

/2*J R*<sup>2</sup>

and *γ* is the electron-phonon interaction parameter, which is given by Eq. (7) [34, 38],

*γ* = 2(

where *J* is the polaron bandwidth, *R* is the mean separation between the transition metal ions,

The condition of small polaron formation was verified using the expression proposed by Emin-Holstein in Ref. [46], an inequality that relates the polaron hopping energy *WH* with the

*WH*

The bandwidth of the polaron *J* was estimated from the fitting parameters using the approxi-

*J* ≈ 0.67 *h υ*<sup>0</sup> (*T*/ *θD*)<sup>1</sup>/<sup>4</sup> . (9)

The type of conduction mechanisms followed by the charge carriers was established using the condition proposed by Holstein [47] to set the limit between the adiabatic and nonadiabatic

mate expression proposed by Holstein [47], as shown in the relationship in Eq. (9),

\_\_\_

and *hν*<sup>0</sup>

) exp(*γ*) = *m*\* exp(*γ*), (6)

*WHh <sup>υ</sup>*0). (7)

<sup>3</sup> . (8)

have been used to

The polaron conduction regime is determined by the inequality, 
$$ \begin{cases} I > H & \text{for adiabatic hopping.} \\ I < H & \text{for non-adiabatic hopping.} \end{cases} \tag{11} $$

Finally, for the case of the nonadiabatic regime, the mobility *μ* has been evaluated using Eq. (11) based on the model proposed by Murawsky [48],

$$
\mu = \left(\frac{\varepsilon R^2}{\hbar kT}\right) \left(\frac{\pi}{4\,\mathcal{W}\_{\text{fl}}\,kT}\right)^{12} \exp\left(\frac{\Lambda W}{kT}\right). \tag{12}
$$

**Table 1** condenses the values obtained for Mott's model parameters from the fit of the electrical conductivity experimental data achieved with A@475 sample [45], together with the values derived for other parameters using the proposed equations. It can be seen that for temperatures in the range used for measurements (from 25 to 100°C), these values are consistent with the simultaneous verification of the four implicit hypotheses used by Schnakenberg in Mott's model simplified formulation and, therefore, allows us to deduce that the small polarons hopping is the conduction mechanism in V<sup>2</sup> O5 thin films obtained for A@475 samples.


**Table 1.** Mott's model parameters obtained for V<sup>2</sup> O5 thin films [45].

As a final remark, it is noteworthy that the low value of mobility has been achieved with A type samples, *μ* = 1.5 × 10−5 cm2 /Vs. Cohen condition [49] establishes that mobility *μ* has to be much lower than 10−2 cm2 /Vs in the case of a hopping conduction process. This condition was completely met in our V<sup>2</sup> O5 thin film type A samples, with a mobility of *μ* = 1.5 × 10−5 cm2 /Vs, while for V<sup>2</sup> O5 , monocrystal mobility value of *μ* = 0.15–0.5 × 10−2 cm2 /Vs [33] is the limit of Cohen condition.

#### **6. Summary and conclusion**

A new method for the preparation of vanadium oxide thin films has been proposed with the aim of providing compatibility with the lift-off microstructuring technique. The thin film formation has been separated into two phases: a first deposition at low temperature using RF magnetron sputtering and an additional thermal annealing at high temperature to adjust film structure after patterning with the lift-off technique. Different starting materials (sputtering targets) and annealing conditions have been analyzed in order to obtain films with high TCR values for application as infrared microsensors (bolometers).

Samples were analyzed using a variety of characterization techniques comprising SEM, XRD, and Raman. I-V curves were measured in a probe station to establish the dependence of electrical conductivity with the temperature for the different samples.

Structural and optical characterization by XRD and Raman shows that V<sup>2</sup> O5 is the predominant oxide phase identified in all samples, even if mixed phases are observed in samples obtained from the metallic-V target (type B samples).

The electrical characterization showed a negative exponential behavior with temperature for all samples, with an activation energy of 0.267 eV in the case of the pure V<sup>2</sup> O5 phase observed in sample A@475, which corresponds to a TCR value of 3.44%/K.

Regarding the electron transport mechanism in processed samples, it has been found that the electrical conductivity measurements performed with the B type VO<sup>x</sup> samples annealed at different temperatures found correlation with a conventional Meyer-Neldel rule, suggesting a thermally activated conduction mechanism.

Finally, it has been found that electrical measurements performed with type A samples, with a pure V2 O5 phase, are consistent with the electron transport model proposed by Mott for the small polarons hopping. The experimental data was fitted using the Schnakenberg simplified formulation obtaining a polaron hopping energy *WH* = 0.1682 eV and an activation energy for hopping due to disorder *WD* = 0.2241 eV. The type of charge transport in type A samples was verified by checking the consistency of the resulting fitting parameters with the implicit hypothesis in Schnakenberg formulation: the conditions for the strong electron-phonon interaction, for the existence of small polarons, for the nonadiabatic regime of the hopping of charge carriers, and for the maximum mobility limit. Overall, results suggest that small polarons hopping is the prevalent mechanism driving the electron transport in V<sup>2</sup> O5 thin films obtained with the proposed method.

### **Acknowledgments**

This research leading to these results and chapter edition has been partially funded by the Spanish Ministerio de Economía y Competitividad, through project TEC2013-48147-C6-6-R (TEMINAIR) and supported by grant PRH 2007 No. 203 (PAE 37079), funded by National Institute of Industrial Technology (INTI) and the National Agency for Promotion of Science, Ministry of Science and Technology of Argentina. The authors also wish to thank B. Halac and E. Di Liscia for the Raman spectroscopy, S. Amore for the XRD diffractograms measurements, and L. Patrone for the SEM images performed on samples.

## **Author details**

formation has been separated into two phases: a first deposition at low temperature using RF magnetron sputtering and an additional thermal annealing at high temperature to adjust film structure after patterning with the lift-off technique. Different starting materials (sputtering targets) and annealing conditions have been analyzed in order to obtain films with high TCR

Samples were analyzed using a variety of characterization techniques comprising SEM, XRD, and Raman. I-V curves were measured in a probe station to establish the dependence of elec-

nant oxide phase identified in all samples, even if mixed phases are observed in samples

The electrical characterization showed a negative exponential behavior with temperature for

Regarding the electron transport mechanism in processed samples, it has been found that the

different temperatures found correlation with a conventional Meyer-Neldel rule, suggesting

Finally, it has been found that electrical measurements performed with type A samples, with

small polarons hopping. The experimental data was fitted using the Schnakenberg simplified formulation obtaining a polaron hopping energy *WH* = 0.1682 eV and an activation energy for hopping due to disorder *WD* = 0.2241 eV. The type of charge transport in type A samples was verified by checking the consistency of the resulting fitting parameters with the implicit hypothesis in Schnakenberg formulation: the conditions for the strong electron-phonon interaction, for the existence of small polarons, for the nonadiabatic regime of the hopping of charge carriers, and for the maximum mobility limit. Overall, results suggest that small polar-

This research leading to these results and chapter edition has been partially funded by the Spanish Ministerio de Economía y Competitividad, through project TEC2013-48147-C6-6-R (TEMINAIR) and supported by grant PRH 2007 No. 203 (PAE 37079), funded by National Institute of Industrial Technology (INTI) and the National Agency for Promotion of Science, Ministry of Science and Technology of Argentina. The authors also wish to thank B. Halac and E. Di Liscia for the Raman spectroscopy, S. Amore for the XRD diffractograms measurements,

ons hopping is the prevalent mechanism driving the electron transport in V<sup>2</sup>

phase, are consistent with the electron transport model proposed by Mott for the

O5

O5

is the predomi-

phase observed

samples annealed at

O5

thin films

values for application as infrared microsensors (bolometers).

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

obtained from the metallic-V target (type B samples).

a thermally activated conduction mechanism.

obtained with the proposed method.

and L. Patrone for the SEM images performed on samples.

**Acknowledgments**

a pure V2

O5

trical conductivity with the temperature for the different samples.

in sample A@475, which corresponds to a TCR value of 3.44%/K.

Structural and optical characterization by XRD and Raman shows that V<sup>2</sup>

all samples, with an activation energy of 0.267 eV in the case of the pure V<sup>2</sup>

electrical conductivity measurements performed with the B type VO<sup>x</sup>

Hernan M. R. Giannetta1, 2, Carlos Calaza3 , Liliana Fraigi1, 2 and Luis Fonseca3

\*Address all correspondence to: hgiannetta@frba.utn.edu.ar

1 Centro de Micro y Nano Electrónica del Bicentenario (CMNB), Instituto Nacional de Tecnología Industrial (INTI), San Martín, Buenos Aires, Argentina

2 Universidad Tecnológica Nacional (UTN) - Facultad Regional Buenos Aires (FRBA), Argentina

3 Institute of Microelectronics of Barcelona (IMB-CNM, CSIC), Campus UAB, Bellaterra, Barcelona, Spain

### **References**


[22] Thornton J.A., Plasma-assisted deposition processes: theory, mechanisms and applications. Thin Solid Films. 1983;**107**(1):3–19. DOI: 10.1016/0040-6090(83)90003-2

[9] Bahlawane N., Lenoble D., Vanadium oxide compounds: structure, properties, and growth from the gas phase. Chemical Vapor Deposition, Wiley. 2014;**20**(7–9):299–311.

[10] Cardarelli F., editor. Materials Handbook: A Concise Desktop Reference. Springer,

[11] Zou C.W., Yan X.D., Han J., Chen R.Q. and Gao W., Microstructures and optical prop-

[12] Ruzmetov D., Zawilski K.T., Narayanamurti V. and Ramanathan S., Structure-functional property relationships in rf-sputtered vanadium dioxide thin films. Journal of Applied

[13] Chae B.G., Kim H.T., Yun S.J., Kim B.J., Lee Y.W. and Kang K.Y., Comparative analysis

[15] Pauli, S.A. and Herger, R. and Willmott, P.R. and Donev, E.U. and Suh, J.Y. and Haglund, R.F., X-ray diffraction studies of the growth of vanadium dioxide nanoparticles. Journal

[16] Rata A.D., Chezan A.R., Haverkort M.W., Hsieh H.H., Lin H.J., Chen C.T., Tjeng L.H.

by wet coating using polyvanadate solutions. Japanese Journal of Applied Physics.

[18] Suh J.Y., Lopez R., Feldman L.C. and Haglund R.F., Semiconductor to metal phase

[19] Melnik V., Khatsevych I., Kladko V., Kuchuk A., Nikirin V., Romanyuk B., Low-

ited by reactive sputtering. Thin Solid Films. 2006;**515**(4): 2519–2522. DOI: 10.1016/j.

thin film based on the sputtering oxidation coupling method. Applied Surface Science.

[21] Xu X., He X., Wang G., Yuan X., Liu X., Huang H., Yao S., Xing H., Chen X. and Chu J., The study of optimal oxidation time and different temperatures for high quality {VO<sup>2</sup>

ometry. Physical Review B. 2004;**69**(7):075404. DOI: 10.1103/PhysRevB.69.075404

Japanese Journal of Applied Physics. 2007;**46**(2R):738. DOI: 10.1143/JJAP.46.738

[14] Piccirillo C., Synthesis and functional properties of vanadium oxides: V<sup>2</sup>

Applied Physics. 2009;**42**(14):145402. DOI: 10.1088/0022-3727/42/14/145402

nanorods prepared by magnetron sputtering. Journal of Physics D:

deposited on glass by aerosol-assisted CVD. Chemical Vapor Deposition.

/Si substrates by the sol–gel process.

thin films with controlled stoichi-

thin films prepared

O2

nanoparticles and thin films. Journal of

phase formation. Materials

thin films depos-

}

O3 , VO<sup>2</sup> ,

DOI: 10.1002/cvde.201400057

O5

erties of β-V<sup>2</sup>

of VO<sup>2</sup>

and V2

O5

tsf.2006.04.025

London-UK. 2008. ISBN 978-1-4471-3648-4

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

Physics. 2007;**102**(11):113715. DOI: 10.1063/1.2817818

2007;**13**(4):145–151. DOI: 10.1002/cvde.200606540

thin films prepared on sapphire and SiO<sup>2</sup>

of Applied Physics. 2007;**102**(7):073527. DOI: 10.1063/1.2786917

and Hibma T., Growth and properties of strained VO<sup>x</sup>

1996;**35**(4A):L438. DOI: 10.1143/jjap.35.l438

transition in the nucleation and growth of VO<sup>2</sup>

[17] Takahashi I., Hibino M. and Kudo T., Thermochromic V1−xW<sup>x</sup>

Applied Physics. 2004;**96**(2):1209–1213. DOI: 10.1063/1.1762995

temperature method for thermochromic high ordered VO<sup>2</sup>

Letters. 2012;**68**(1):215–217. DOI: 10.1016/j.matlet.2011.10.075

2011;**257**(21):8824–8827. DOI: 10.1016/j.apsusc.2011.04.068

[20] Fu G., Polity A., Volbers N., and Meyer B., Annealing effects on VO<sup>2</sup>


**Advances of Thin-film Technologies in Medicine and Biology**

[36] Yelon A., Movaghar B. and Crandall R.S., Multi-excitation entropy: its role in thermodynamics and kinetics. Reports on Progress in Physics. 2006;**69**:1145. DOI: 10.1088/00

[37] Spear W.E., Allan D., Comber P.L., Ghaith A., A new approach to the interpretation of transport results in a-Si. Philosophical Magazine Part B. 1980;**41**(4):419–438. DOI:

[38] Mott N., Conduction in glasses containing transition metal ions. Journal of Non-

[39] Davis E.A. and Mott N.F., Conduction in non-crystalline systems V. Conductivity, optical absorption and photoconductivity in amorphous semiconductors. Philosophical

[41] Kittel C., editor. Introduction to Solid State Physics. 8th ed. USA: Wiley; 2004. 704pp.

[42] Devreese J. T., "Polarons" in Encyclopedia of Applied Physics, Vol. 14, pp. 383 – 409. Ed.

[43] Schnakenberg J., Polaronic impurity hopping conduction. Physica Status Solidi (b).

[44] Culea E., Gheorghiu, C. and Nicula A., Electrical conductivity of vitreous

[45] Giannetta H.M.R., Calaza C., Lamas D., Fonseca L., Fraigi L., Electrical transport prop-

[46] Emin D., David E., Holstein T., Studies of small-polaron motion IV. Adiabatic theory of the Hall effect. Annals of Physics. 1969;**53**(3):439–520. DOI: 10.1016/0003-4916(69)90034-7

[47] Holstein T., Studies of polaron motion: Part I. The molecular-crystal model. Annals of

[48] Murawski L., Chung C.H. and Mackenzie J.D., Electrical properties of semiconducting oxide glasses. Journal of Non-Crystalline Solids. 1979;**32**(1):91–104. DOI: 10.1016/0022-3

[49] Cohen M., Review of the theory of amorphous semiconductors. Journal of Non-

Crystalline Solids. 1970;**4**:391–409. DOI: 10.1016/0022-3093(70)90068-2

Physics. 1959;**8**(3):325–342. DOI: 10.1016/0003-4916(59)90002-8

tron sputtering at room temperature. Thin Solid Films. 2015;**589**:730–734. DOI: 10.1016/j.

), Physica Status Solidi (a). 1986;**96**(1):K85–K88. DOI: 10.1002/

thin films obtained by thermal annealing of layers grown by RF magne-

O5

. Physica

Crystalline Solids. 1968;**1**(1):1–17. DOI: 10.1016/0022-3093(68)90002-1

Magazine. 1970;**22**(179):903–922. DOI: 10.1080/14786437008221061

Status Solidi (a). 1983;**76**(2):661–666. DOI: 10.1002/pssa.2210760232

[40] Sanchez C., Morineau R., Livage J. Electrical conductivity of amorphous V<sup>2</sup>

34-4885/69/4/R04

10.1080/13642818008245397

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

ISBN: 978-0-471-41526-8

O5 − 25%(As<sup>2</sup>

O5

75%V<sup>2</sup>

pssa.2210960163

erties of V<sup>2</sup>

tsf.2015.06.048

093(79)90066-8

by G.L. Trigg (VCH, Weinheim, 1996).

O3 ·B2 O3

1968;**28**(2):623–633. DOI: 10.1002/pssb.19680280220

### **Smart Thermoresponsive Surfaces Based on pNIPAm Coatings and Laser Method for Biological Applications Smart Thermoresponsive Surfaces Based on pNIPAm Coatings and Laser Method for Biological Applications**

Laurentiu Rusen, Valentina Dinca, Cosmin Mustaciosu, Madalina Icriverzi, Livia Elena Sima, Anca Bonciu, Simona Brajnicov, Natalia Mihailescu, Nicoleta Dumitrescu, Alexandru I. Popovici, Anca Roseanu and Maria Dinescu Laurentiu Rusen, Valentina Dinca, Cosmin Mustaciosu, Madalina Icriverzi, Livia Elena Sima, Anca Bonciu, Simona Brajnicov, Natalia Mihailescu, Nicoleta Dumitrescu, Alexandru I. Popovici, Anca Roseanu and Maria Dinescu

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

http://dx.doi.org/10.5772/66280

#### **Abstract**

Various applications within last decades such as bacterially resistant surfaces, soft robotics, drug delivery systems, sensors and tissue engineering are poised to feature the importance of the ability to control bio-interfacial interactions. An enhanced attention is dedicated to designing smart stimuli-responsive interfaces for DNA, drug delivery, protein and cell based applications. Within this context, the thermoresponsive materials, especially poly(N-isopropylacrylamide) (pNIPAm) have been intensively used in tissue engineering applications for a controlled detachment of proteins and cells with a minimum of invasive effect on protein and cell structural conformation. The properties of smart bio-interfaces can be controlled by its composition and polymer architecture. Therefore, appropriate methods for obtaining controlled coatings are necessary. Laser methods were successfully used in the last decades for obtaining controlled organic and inorganic coatings for various types of applications, from electronics to tissue engineering. Among these, Matrix-Assisted Pulsed Laser Evaporation (MAPLE) technique bring us a step forward to other laser methods by avoiding damage and photochemical decomposition of materials. In this chapter we describe materials and approaches used for design of smart bio-interfaces aimed at controlling protein and cells behavior *in vitro*, focusing MAPLE method for tuning coatings characteristics in relation with biological response.

**Keywords:** bio-smart interfaces, temperature responsive, laser processing, mammalian cells

© 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.
