Transparent Conducting Oxides on Polymeric Substrates · 2013-07-14 · given to illustrate the diversity of materials and deposition techniques. In chapter 3 the basis principles

Transparent Conducting Oxides on Polymeric Substrates

by Pulsed Laser Deposition

Cover: The diffraction pattern of a polycrystalline In2O3 thin film imaged by transmission electron microscopy. The colors are inverted. Each dot that appears on the detector originates from electrons being diffracted on one single grain. The rings correspond to different lattice spacing of the atomic planes within the grains.

Ph.D. committee: Chairman and secretary:

prof. dr. ir. J.A.M. Kuipers (University of Twente)

Supervisor: prof. dr. ing. D.H.A. Blank (University of Twente)

Assistant supervisor: dr. ing. A.J.H.M. Rijnders (University of Twente)

Members: prof. dr. S.J. Picken (Delft University)

prof. dr. M. Pollnau (University of Twente)

prof. dr. K.-J. Boller (University of Twente)

Referents: dr. C.P.G. Schrauwen (TNO, Arnhem)

dr. J. Wienke (ECN, Petten)

The work described in this thesis was carried out at the Inorganic Materials Science department at the Faculty of Science and Technology and the MESA+ Institute for Nanotechnology, University of Twente, P.O. Box 217, 7500 AE, Enschede, The Netherlands. This research was financially supported by the Dutch SenterNovem-IOP program: Oppervlaktetechnologie (No. IOT00002). J.M. Dekkers Transparent Conducting Oxides on Polymeric Substrates by Pulsed Laser Deposition, Ph.D. thesis University of Twente, Enschede, The Netherlands. ISBN: 978-90-365-2462-9 Printed by Wöhrmann Print Service, Zutphen. © J.M. Dekkers, 2007

TRANSPARENT CONDUCTING OXIDES

ON POLYMERIC SUBSTRATES

BY PULSED LASER DEPOSITION

PROEFSCHRIFT

ter verkrijging van de graad van doctor aan de Universiteit Twente,

op gezag van de rector magnificus, prof. dr. W.H.M. Zijm,

volgens besluit van het College voor Promoties in het openbaar te verdedigen

op donderdag 8 maart 2007 om 15:00 uur

door

Jan Matthijn Dekkers

geboren op 8 augustus 1977 te Almelo

Dit proefschrift is goedgekeurd door:

prof. dr. ing. D.H.A. Blank (promotor) en

dr. ing. A.J.H.M. Rijnders (assistent-promotor)

Contents

1 Introduction 9 1.1 Motivation ………………………………………………………………… 10 1.2 Outline of this thesis …………………………………………………….. 11 1.3 References ……………………………………………………………….. 13

2 Fundamentals of transparent oxide semiconductors 15 2.1 Introduction ……………………………………………………………….. 16 2.2 Electrical properties of wide bandgap semiconductors ……………… 17 2.2.1 Electronic conductivity and band structure …………………. 17 2.2.2 Scattering mechanisms ………………………………………. 19 2.2.3 Intrinsic conductivity limit ……………………………………... 22 2.3 Optical properties ………………………………………………………... 24 2.3.1 Transparency ………………………………………………….. 24 2.3.2 Correlation of optical and electrical properties …………….. 26 2.4 General properties ………………………………………………….…… 27 2.5 Concluding remarks ……………………………………………………... 28 2.6 References ……………………………………………………………….. 31 3 Fabrication and characterization of transparent semiconductors 35 3.1 Introduction ……………………………………………………………….. 36 3.2 Thin film growth by pulsed laser deposition …………………………... 36 3.2.1 Basic principles ……………….……………………………….. 36 3.2.2 Experimental setup ………………….………………………… 37 3.3 Thin film characterization …………….…………………………………. 39 3.3.1 Electronic transport analysis ……….………………………… 39 3.3.2 Optical analysis ………….…………………………………….. 40

Contents

3.3.3 Structural analysis …………….………………………………. 41 3.3.4 Morphology ……………….……………………………………. 42 3.3.5 Compositional analysis ……….………………………………. 42 3.4 Substrates ……………….……………………………………………….. 43 3.4.1 Polymers ………….……………………………………………. 43 3.4.2 Mechanical properties of metal-oxides on polymers …….… 45 3.4.3 Non-polymer substrates ………….…………………………… 46 3.5 Concluding remarks ………….………………………………………….. 47 3.6 References ………….……………………………………………………. 48

4 Indium (Tin) Oxide 51 4.1 Introduction ……………………………………………………………..… 52 4.2 PLD growth of indium tin oxide on polymers …………..……………... 53 4.3 Performance of ITO thin films on PET ….……………………………... 56 4.3.1 Results ………….………………………………………………. 56 4.3.2 Electrical properties …………………………..……………….. 57 4.3.3 Optical properties ……………………………………………… 59 4.4 Microstructure ……………………………………………………………. 61 4.4.1 TEM analysis ………….……………………………………….. 61 4.4.2 Surface morphology ……….………………………………….. 62 4.5 Discussion ………………………………………………………………... 65 4.6 Concluding remarks ……………………………………………………... 67 4.7 References ….……………………………………………………………. 69

5 Doped Zinc Oxides 73 5.1 Introduction ……………………………………………………………….. 74 5.2 ZnO and AZO deposition on PET ……………………………………… 75 5.3 Results ……………………………………………………………………. 78 5.3.1 Undoped ZnO thin films ………………………………………. 78 5.3.2 Al2O3 doped ZnO ………………………………………………. 79 5.3.3 Review and discussion ……………………………………….. 81 5.4 Indium zinc oxide ………………………………….…………………….. 82 5.5 Concluding remarks …..…………………………….…………………… 86 5.6 References …………….…………………………………………………. 88

6 P-type conducting ZnM2O4 91 6.1 Introduction ……………….………………………………………………. 92 6.2 Bandgap phenomena in spinel structures ………………….…………. 93 6.2.1 Ligand field theory ….…………………………………………. 93 6.2.2 d- orbital splitting in octahedral structures …….…………..... 94 6.2.3 Bandgap engineering in ZnM2O4 …………….………………. 96

Contents

6.3 Bulk ZnM2O4 synthesis ……………………………….…………………. 97 6.4 ZnM2O4 growth and analysis ……….…………………………………... 99 6.4.1 Thin film deposition ………………….………………………… 99 6.4.2 Structural analysis ……………….……………………………. 100 6.4.3 Morphology and stoichiometry ……….………………………. 102 6.4.4 Electronic band structure …………….……………………….. 103 6.5 Electrical and optical properties ……………….……………………….. 104 6.5.1 Optical analysis ……………………….……………………….. 104 6.5.2 Electrical conduction behavior …………………….…………. 105 6.5.3 Discussion ……………………..………………………………. 107 6.6 Concluding remarks ……………………………………………………... 108 6.7 References ………………………………………………………….……. 110 7 Amorphous p-type TOS on polymer substrates 113 7.1 Introduction …………………………………………………….…………. 114 7.2 Thin film deposition by “eclipse” PLD …………….……………………. 114 7.3 Structural and chemical composition …………….……………………. 118 7.3.1 Structural analysis …………………………………………….. 118 7.3.2 Stoichiometry of “eclipse deposited” films …….……………. 119 7.3.3 Chemical composition ……………………….………………... 120 7.3.4 Iridium valency ….…….……………………………………….. 123 7.4 Performance of amorphous films ………………………………………. 125 7.4.1 Optical behavior ………………………………….……………. 125 7.4.2 Electrical conduction behavior …………………….…………. 126 7.4.3 Conduction paths ………………………….…………………... 127 7.5 Transparent devices …………………………….………………………. 129 7.5.1 P-N junctions ……………………………….………………….. 129 7.5.2 Transparent electronics on plastic ……………….………….. 131 7.6 Concluding remarks ……………………….…………………………….. 133 7.7 References ………………………………….……………………………. 134 Summary 137 Samenvatting (Summary in Dutch) 141 Dankwoord 145

Contents

Chapter 1

Introduction

Abstract

This research on transparent conducting oxides on polymers is motivated by the increasing need for applications on flexible, cheap and light-weight substrates. This thesis describes various wide bandgap semiconducting materials that are grown on polymers by pulsed laser deposition. Their electrical and optical performance is in a large extent determined by fundamental properties as structure and composition, which can be tuned and manipulated in the deposition process.

Chapter 1 10

1.1 Motivation In our daily lives transparent conducting oxide (TCO) materials are used in

numerous devices. These applications are mostly found in display technology, (organic) light-emitting-diodes (OLEDs), thin-film solar photovoltaics and ‘smart windows’. Both in existing as well as new applications the implementation of a polymer support instead of the commonly used rigid glass has become significant in the field of research, for example in the realization of flexible displays. As polymer materials with functional or protective coatings are combined, the flexibility, light-weight and cost-effectiveness of the substrate offer many new advantages for device engineering. With regard to processing however, the manufacturing temperature of TCO materials on polymers is limited, which negatively influences their performance [1].

At present, Indium Tin Oxide (ITO) is commonly used as transparent electrodes, especially in large-area applications such as displays. As can be seen from figure 1.1 displays have quickly grown to become the primary end user of the world’s indium productionI. The supply of indium has become unreliable during the last 20 years [2], whereas the display market is still expanding. Indium prices are therefore subjected to major instabilities. Alternative materials are welcome to manufacturers of flat panel displays, and much research is conducted on suitable candidates.

With the discovery of p-type TCOs a new field in optoelectronics device technology has opened up [3]. In contrast to passive electrodes, the use of transparent conductors as active components became possible. This has led to new applications as blue or ultra-violet (UV) light-emitting-diodes [4,5,6]. Furthermore, it

I Market research of AIM Specialty Materials, Cranston, Rhode Island, USA.

Semiconductors 3%

Alloys 10%

Compounds 8% R&D

4%

Thin Films 75%

Figure 1.1: World-wide indium consumption by end use in 2004. The majority of indium used in thin filmapplications is the display industry.

Introduction 11

became possible to design all-transparent devices [7,8]. Although the fabrication of transparent circuits is nowadays limited to laboratory research [9], they are believed to find revolutionary applications in the next decades. However, appropriate p-type TCOs for this purpose are not available yet [10]. Moreover, most are far from suitable for processing on polymeric materials.

It is evident that devices employing TCOs become ever more sophisticated [11], and that the demand for implementation of polymer materials is increasing. The process conditions of TCOs need to be adjusted to allow coating of such heat-sensitive substrates. Under these circumstances it is important to optimize the TCO performance in order to meet the technological requirements. Understanding the fundamental properties is an important for the improvement of existing materials. Moreover, these insights are of great scientific importance for the design and synthesis of new type of TCOs. The use of a deposition technique that allows indirect, but precise control of the fundamental properties can assist the research of high-performance TCOs on polymer substrates. 1.2 Outline of this thesis

The aim of this thesis is the fabrication of TCOs on polymeric substrates without deteriorating the high electrical and optical performance commonly achieved, and to create an understanding of the fundamental aspects in order to optimize and develop these materials.

In the research described in the following chapters pulsed laser deposition is employed to coat polymers with transparent conducting oxide thin films. The investigated compounds are indium (tin) oxide, (doped) zinc oxide and zinc containing compositions, both n- and p-type. Extensive analysis is conducted by several measuring tools on the substrates, bulk materials and thin films in order to obtain insight on their relevant properties.

In chapter 2 a brief overview is given on the fundamental aspects of transparent conducting oxides. In particular the electrical and optical properties are discussed, since these have the most importance influence on the performance of these materials. An overview of transparent n- and p-type conducting oxides is given to illustrate the diversity of materials and deposition techniques.

In chapter 3 the basis principles and the growth procedure by pulsed laser deposition technique is described. The analytical tools to obtain information on the electrical and optical properties, structure, composition and morphology of the deposited thin films are summarized. Furthermore, the choice of the (polymer) substrates used in the experiments for coating with TCOs is motivated.

In chapter 4 the experiments on indium (tin) oxide are described. The dependence of the electrical properties on the doping content of tin oxide is

Chapter 1 12

investigated. Methodical structural analysis is performed in order to relate the observed electrical properties to the microstructure of the thin films. The sources of scattering are modeled and the influence on the conduction behavior is discussed.

In chapter 5 investigations on zinc related materials are presented. The influence on impurity doping of zinc oxide on the electrical performance is examined. Furthermore, zinc oxide is used in mixed compounds with indium oxide to combine their good properties. The influence of this composition on the optical and electrical performance when deposited on polymers is described.

From chapter 6, the research will focus on p-type transparent conductors. In this chapter, the synthesis of a new crystalline semiconductor, spinel ZnIr2O4, is presented. A hypothesis from theory with regard to the bandgap of this material is postulated. The fundamental optical and electrical aspects are investigated and related to ZnCo2O4 and ZnRh2O4.

In chapter 7, thin films of ZnIr2O4 are deposited at room temperature on polymers. The influence of the oxidation state in the amorphous thin films on the transparency and conductivity is described. Using this material, p-n junctions are fabricated and analyzed. Subsequently, the applicability of this material in devices on polymer substrates is demonstrated.

Introduction 13

1.3 References 1. F. Hanus, A. Jadin, and L.D. Laude, Appl. Surf. Sci. 96-98, 807 (1996). 2. B. O’Neill, Indium: Supply, Demand & Flat Panel Displays, Presented at

Minor Metals 2004, London, June 2004. 3. A.N. Banerjee & K.K. Chattopadhyay, Prog. Cryst. Growth Charact. Mater.

50, 52 (2005). 4. H. Ohta, K.-I. Kawamura, M. Orita, M. Hirano, N. Sarukura, & H. Hosono,

Appl. Phys. Lett. 77, 475 (2000). 5. K. Tonooka, H. Bando & Y. Aiura, Thin Solid Films 445, 327 (2003). 6. H. Hosono, H. Ohta, K. Hayashi, M. Orita & M. Hirano, J. Cryst. Growth 237-

239, 496 (2002). 7. G. Thomas, Nature 389, 907 (1997). 8. J.F. Wager, Science 300, 1245 (2003). 9. R.E. Presley, D. Hong, H.Q. Chiang, C.M. Hung, R.L. Hoffman & J.F. Wager,

Solid-State Electron. 50, 500 (2006). 10. H. Kawazoe, H. Yanagi, K. Ueda & H. Hosono, MRS Bull. 25, 28 (2000). 11. D.S. Ginley & C. Bright, MRS Bull. 25, 15 (2000).

Chapter 1 14

Chapter 2

Fundamentals of transparent oxide semiconductors

Abstract

Optical transparency and electrical conductivity can coexist in a specific class of oxide materials. These so-called transparent conducting oxides are semiconductors that transmit visible light because of a wide bandgap. By introducing doping, it is possible to control the conductivity. At very high doping concentration the semiconductor becomes degenerate, and the conductivity is limited by scattering of the charge carriers. Furthermore, the optical properties are affected by the doping concentration. Both the bandgap energy and the plasma wavelength change with the carrier density.

Chapter 2 16

2.1 Introduction A combination of transparency and conduction can be achieved in two

classes of material. The first group is formed by extremely thin (~10 nm) metal films, especially Ag, Au or Cu. The luminous transmittance can go up to 50%, and even higher when the metal is sandwiched between anti-reflecting layers [1]. The conductivity is strongly thickness dependant, and therefore quite low for these very thin films. The second class of materials consists of wide bandgap semiconductors. As early as 1907, coexistence of electrical conductivity and optical transmittance was first observed in cadmium oxide [2]. However, the technological advances in Transparent conducting oxides (TCOs) emerged only after 1940, as the potential applications in industry and research became evident. In 1956 Thelen et al. [3] found transparency and conduction in indium oxide, which was later used in applications for heated windows. Years of extensive research finally led to SnO2 doped In2O3 (known as indium tin oxide or ITO) with excellent electrical and optical properties [4].

Limitations of TCOs become more critical as passive and active devices based on these materials get more sophisticated. For example, as displays become larger and writing speeds get faster, it becomes increasingly important to decrease resistivity while the transparency is maintained [5]. Simply increasing the thickness is not acceptable since the optical absorption will increase. A profound understanding of the fundamental aspects of transparent semiconductors is therefore required in order to improve either the properties of existing materials, or design new type of TCOs. These insights are of great scientific importance whether the realization of high-performance TCOs on polymer substrates is possible.

A vast amount of literature is available on different TCOs deposited by various growth techniques and their fundamental aspects. Reviews on TCOs concerning these issues have been reported earlier [6,7]. In this chapter an overview of the most important fundamental aspects on transparent oxide semiconductors (TOSs) will be discussed. The sections 2.2 and 2.3 discuss the electronic and optical properties of TOSs, respectively. Most of the considerations refer to indium oxide and indium tin oxide as an example. Still, they hold for all other types of n-type transparent and conducting oxides if the correct material related numbers are used in the calculations. In most cases this also holds for p-type semiconductors. The carrier density then refers to the hole concentration in the material. However, mobility and carrier density in p-type wide bandgap semiconductors usually differ by a few orders of magnitude. The conduction mechanism can therefore be quite different. Some specific electrical transport phenomena considering p-type TOSs will be discussed in more detail in chapter 6

Fundamentals of transparent oxide semiconductors 17

and 7. Finally, section 2.4 deals with the general aspects of TCOs, and an overview of different materials is tabulated.

2.2 Electrical properties of wide bandgap semiconductors

2.2.1 Electronic conductivity and band structure The conductivity σ is a product of the number of charge carriers n in a

material, and the mobility µ of these charge carriers, times the elementary electron charge e. The resistivity ρ is defined as the inverse of the conductivity.

1e nσ µρ

= ⋅ ⋅ = (2.1)

For thin films of uniform thickness d, the electrical resistance is sometimes expressed as the sheet resistance (Rs=ρ/d). Other than the thickness, the sheet resistance is independent of the film dimensions.

In order to promote conductivity, the number of charge carriers can be increased by doping. Dependant on the material this can be done by substitutional doping, creation of vacancies or implantation of interstitials. Dependant on the valence of dopants or vacant sites, acceptor or donor states will induce p- or n-type conductivity.

Another possibility to enhance the conductivity is to increase the mobility. However, the mobility is dependant on intrinsic scattering mechanisms, and can therefore not be controlled directly. In general these mechanisms limit the mobility as the carrier density increases. As a result, the mobility is the most important parameter influencing the total conductivity.

The presence of a bandgap, providing low absorption in the visible range, is an essential feature of TOSs. Metal-oxide semiconductors having a bandgap of at least 3 eV meet this condition. The top of the valence band is mostly formed by the oxygen 2p bands, whereas the bottom of the conduction band is composed of a single and highly dispersed metal s band. The O 2p orbitals are low in energy and it is therefore that a large band gap can be obtained in oxides. In intrinsic stoichiometric oxides, coexistence of electrical conductivity besides visible transparency is not possible. However, substitutional doping by cationic donors or anion vacancies can create charge carriers, i.e. electrons. The donor (or acceptor) states alter the electronic band structure of the material. For increased donor density, the donor states merge with the conduction band at a certain critical density nc, whose magnitude can be estimated by Mott’s criterion [8,9].

1/3 *0 0.25cn a⋅ ≈ (2.2)

The effective Bohr radius a0* is given by:

Chapter 2 18

20

0 2**

m

c

hae mε ε

π= (2.3)

Where εm is the static dielectric constant of the host lattice (equals 8.9 for In2O3 [10]), and mc* is the effective mass of the electrons in the conduction band. After comparing experimental data of several authors, Köstlin et al. [11] calculated the effective mass of In2O3 to be about a third of the electron restmass; i.e. mc*=0.35·me. Using these numbers, one can obtain a0*≈1.3 nm, and the critical density nc is calculated as 6.4 x1018 cm-3. In general, the carrier density in TCOs is by far larger than this (up to 1021 cm-3 in ITO). Above this Mott critical density, free electron behavior can be expected. The donor states have merged with the conduction band and the material is said to be degenerate. The Fermi energy EF is determined by the highest occupied state of the conduction band, and one can write:

2 2

2 *Fc

kEm

= (2.4)

The band structure of indium (tin) oxide is approximated by parabolic functions of k close to the band edges. The schematic representation in figure 2.1 is valid for most binary metal oxide wide bandgap semiconductors. The valence band maximum and conduction band minimum are both located at k=0, so the material is

2 20

0( )2 *C g

C

kE k Em

= + 0( ) ( ) ( )C C CE k E k k= + ∑

2 20( )

2 *VV

kE km

= − 0( ) ( ) ( )V V VE k E k k= + ∑

0gE gE

ω ω

k kFk

E∆

Figure 2.1: Schematic representation of the bandstructure of undoped (a) and doped (b) metal oxide wide bandgap semiconductor in the vicinity of the top of the valence band and bottom of theconduction band. The grey areas denote the occupied states.

(a)

(b)


said to be a direct band semiconductor. If optical transitions or carrier transport occur far from the conduction band bottom or valence band top, the shape of the bands are slightly distorted. This band distortion is caused by interactions with other bands and many-body effects. It is therefore hard to determine the effective mass theoretically in these regions, since this value is determined by the band curvature. For theoretical calculations on the band structure, non-parabolic effects should be taken into consideration [12]. 2.2.2 Scattering mechanisms

As mentioned earlier, the conductivity of TCOs is reflected from the mobility that can be expressed as:

* *mfp

F

eem m V

λτµ⋅⋅

= =⋅

(2.5)

The relaxation time τ depends on the drift velocity VF and mean free path λmfp of the charge carriers. The parameters are affected by the different scattering mechanisms; lattice scattering, ionized impurity scattering, neutral impurity scattering, electron-electron scattering, electron-impurity scattering and grain boundary scattering [13]. The total mobility can therefore be written as:

1 1itot iµ µ

= ∑ (2.6)

In crystalline TCOs all scattering mechanisms have little effect, except the ionized impurity scattering [10]. However, TCOs deposited at lower temperatures posses a lower crystalline nature, and high doping concentrations result in the formation of neutral complexes [14]. In these cases grain boundary and neutral impurity scattering should also be taken into consideration. Ionized Impurity Scattering The Coulomb interaction between the ionized (donor) impurities and the free electrons provide a source of scattering that is intrinsic to the doped material, and can therefore set a lower limit to the resistivity, regardless of other scattering mechanisms such as neutral impurities, grain boundaries or structural disorder. The attainable resistivity due to ionized impurity scattering was first calculated by Brooks (1955) [15] and Dingle (1955) [16] using the Born approximation band Thomas-Fermi screening. The theory assumes that the ionized impurities form a uniform background of immobile charges. The mobile charges (electrons or holes) provide screening. The relaxation time of an electron of wave vector k is then given by:

Chapter 2 20

2 4

2 3 30

*1 ( )( ) 8 ( )

i

r

N Z e m f kk kτ π ε ε

= (2.7)

Here is Ni the number of scattering centers per unit volume and Z the charge on the impurity, ε0 represents the permittivity of free space whereas εr is the low frequency relative permittivity. f(k) is the screening function and is dependant on the Thomas-Fermi screening wave vector. For a degenerate system it is only necessary to consider the scattering of electrons at the Fermi surface (k=kF). The upper limit of the mobility can now be expressed by:

3 2 30

2 3 2

24 ( ) 1* ( )

riis

i F

nN Z e m f kπ ε εµ = (2.8)

Later, Moore [17] made corrections to the original work of Dingle based on second order terms in the Born approximation. This resulted in a slightly lower mobility limit. Pisarkiewicz et al. [18] applied corrections to this model by taking into account the non-parabolic band shape. This caused a larger effect on the intrinsic limit. According to his approximation, the effective mass at the Fermi-energy is dependant on the carrier concentration as;

22 2/3

00

* * 1 2 (3 )*

m m C nm

π= + (2.9)

where C is a constant used to fit experimental data. Equations 2.7 and 2.8 can be used to calculate the intrinsic mobility limit as a function of carrier concentration n. Grain boundary scattering

The grain boundary is a complex structure, usually consisting of a few atomic layers of disordered atoms. These incomplete atomic bonds induce a large number of defects. Here charge carriers are trapped; i.e. immobilized. The traps become electrically charged, creating a potential energy barrier, reducing the mobility of free carriers from moving from one crystallite to the other. Seto [19] assumed a barrier with thickness d, which is considerably smaller than grainsize L. Further the grain boundary contains Qt traps per cubic centimeter located at energy Et with respect to the intrinsic Fermi level. All charges in the region with thickness d around the grain boundary are assumed to be trapped, and form a depletion layer. The grain boundary can now be thought of as a potential barrier for electrons characterized by its height Eb. The contribution µgb is thermally activated and can be described by the Petritz [20] relation:


0 exp( )bgb

b

Ek T

µ µ= − , with 0 2 * b

Lqm k T

µπ

= (2.10)

Where q is the charge of the trap, kb is Boltzmann constant and T is the absolute temperature. The energy barrier height is:

2 2

08t

br d

q QENε ε

= , with d tN n Q d= + (2.11)

Nd is the total number of ionized donors. Free fitting parameters are; the number of traps Qt and barrier thickness d. This model was later applied to TCO materials [18], and data of polycrystalline ITO [21], F:SnO2 [22] and ZnO [23] films was fitted to this relation. It has been shown that the energy barrier height is approximately 0.01 eV [13,21]. The grain boundary scattering has an effect on the total mobility only if the grainsize is approximately of the same order as the mean free path of the charge carriers (L~λmfp). The mean free path can be calculated for known carrier density and mobility. Using a highly degenerate electron gas model [24] the electron velocity VF is:

2 1/3 1/3(3 ) ( / *)FV m nπ= (2.12)

Substitution in equation 2.5 results in the following expression for the mean free path λmfp:

2 1/3 1/3* (3 ) ( / )F FmV V e n

eµλ τ π µ= = = (2.13)

Neutral impurity scattering The solubility of dopants in oxide semiconductors can be quite high. The

solid solubility limit of Sn in ITO for instance is found to be well over 60% [11]. At high doping concentrations, an increasing amount of dopant material remains inactive, and can form a variety of neutral complexes [25]. Therefore neutral impurity scattering might be taken into consideration. The contribution to the mobility µN is given by [26,27,28]:

3

30

*20N

r N

m en

µε ε

= (2.14)

Here nN is the concentration of neutral impurities. This value is often estimated from the difference between measured carrier density and carrier concentration. The

Chapter 2 22

ionized impurity and grain boundary scattering on the other hand, show a direct relation with the carrier concentration n. 2.2.3 Intrinsic conductivity limit According to the previous sections, the maximum conductivity of wide bandgap semiconductors is restricted by the theoretical upper limit of mobility. This intrinsic mobility limit is dependant on carrier concentration for each scattering mechanism. Modeling the upper limit of separate or combined mechanisms for different TCO materials has been reported frequently [13-29]. In figure 2.2 the upper of mobility was simulated for polycrystalline indium oxide with intermediate grain size (50 nm). The total mobility limit is determined by the contributions of the scattering mechanisms from relations 2.8, 2.10 and 2.11. The non-parabolicity of the conduction band has been taken into account in the calculations. It can be seen

Figure 2.2: The upper limit of mobility (solid line) influenced by the effect of scattering on ionizedimpurities (dashed curve), grain boundaries (dotted curve) and neutral impurities (dashdotted curve). µiis was calculated for double ionized donors (oxygen vacancies; Ni•Z2=2n) resulting in larger contribution. The average grainsize for calculating µgb was taken as 50 nm, and the barrier height was fixed at 0.01 eV. The number of neutral impurities was taken as 50% of the amount of donors. Striped areas indicate the spread of data points in literature.The effect of ionized impurities is dominant (horizontal lines) or grain boundaries can play a role (vertical lines).

1019 1020 1021 1022100

101

102

103

104

105

µtot

µiis µgb

mob

ility

µ (c

m2 V

-1s-1

)

carrier concentration n (cm-3)

µN


that the importance of electron scattering on neutral impurities has negligible effect on the total mobility. The amount of neutral impurities in relation 2.11 was even overestimated since it was set to half the number of free carriers. The values for the mobility reported in literature are all located in the marked areas of figure 2.2. Best performing films are found in the region where the ionized impurity scattering is dominant (horizontal lines). If the grain size is small enough, TCO thin films up to intermediate carrier density are affected by grain boundary scattering as well (vertical lines).

Although most authors agree on the mobility limit due to ionized impurity scattering, the mobility due to grain boundary scattering is modeled differently in many cases. Apparently the free parameters are used to fit the obtained data exactly, whereas in general the average grainsize is just a rough estimate. Therefore the influence of grain boundary scattering in TCOs is still under debate [30]. However, it is commonly accepted that the mobility is only affected if the grain size is considerably small. And moreover this effect is only noticeable up to intermediate carrier concentrations, as for high electron density the ionized impurity scattering dominates. From figure 2.2 it becomes evident that the structure of TCO thin films affects the electrical properties on going from single- to polycrystalline phase. As the grainsize decreases further, the thin film is said to be nanocrystalline or even amorphous. Although the grain boundary density increases to infinity in the amorphous state, the mobility does not drop to zero. Namely in some amorphous oxide materials a reasonable mobility can be obtained, resulting in amorphous

(a)

(b)

(c)

Figure 2.3: Schematic illustration of the electron density in the atomic orbitals responsible for the carriertransport paths in transparent conducting oxides consisting of light metal cations (a), heavymetal cations in crystalline phase (b) and amorphous phase (c). In ‘ns orbitals’, the n denotes the principal quantum number.

Chapter 2 24

conductivity. In most n-type transparent conductor, the bottom of the conduction band consists of a highly dispersed single s-band. This high dispersion and s-type character results in a rather uniform distribution of the electron charge density. The scattering of these states is relatively low, and provides a large mobility [31]. As mentioned earlier, a high mobility is critical in attaining a high conductivity. Since the mobility is proportional to the width of the conduction bands, a large overlap between the relevant orbitals is required [32]. The magnitude of this overlap is also of great importance in case of polycrystalline or amorphous materials. A large overlap is namely quite insensitive to the structural randomness of the amorphous state. The spatial spreading of the spherical s-orbitals should be large enough to create continuous conduction paths. In particular for post transition metal cations, a direct overlap between neighboring metal s-orbitals is possible [31,33]. The conduction paths become insensitive to the distorted metal-oxygen-metal bonds in amorphous materials as schematically illustrated in figure 2.3 [34]. For this reason, the mobility’s in amorphous oxide semiconductors can be similar to the corresponding crystalline phase (>10 cm2s-1V-1). 2.3 Optical properties 2.3.1 Transparency

An important feature of TCOs is the existence of a transmission window covering most part of the visible spectrum. In literature, the optical transmission is defined as the ratio between incoming light intensity and transmitted intensity averaged over all values in between 400 nm and 700 nm. The typical spectral dependence of TOSs (from [35]) is schematically shown in figure 2.4. The

Figure 2.4: Spectral dependence of semiconducting transparent materials: λgap and λp are the wavelengths at which the bandgap absorption and free electron plasma absorption takesplace.

λpλgap

ReflectingTransmittingAbsorbing

Tran

smitt

ance

Wavelength (nm)


transmission window is defined by two regions where no light is transmitted due to different phenomena. At low wavelengths (λ<λgap) the absorption due to the fundamental bandgap dominates. The photon energy in this near-and deep-UV part of the spectrum is high enough to equal the bandgap energy (3-4 eV). This energy is absorbed and transformed to band to band transitions, and no light is transmitted because of this quantum phenomenon. For longer wavelengths, in the (near) infrared (IR) part of the spectrum no light is transmitted due to the plasma edge (λ>λp). Here the light is electronically reflected which can be best described by the classical Drude free electron theory [36].

In the free electron model, the electrons may be thought of as a plasma whose density is set into motion by the electric field component of the electromagnetic field. The plasma oscillates at a natural frequency ωp, the resonance or plasma frequency. This frequency corresponds to the plasma wavelength λp and is of the order of 1-4 µm for TCOs [37]. The interaction of free electrons with the electromagnetic field influences the relative permittivity ε of the material, which is expressed as a complex number:

2( ')N ikε = − (2.15)

The real and imaginary part is the refractive index (Ν) and the extinction coefficient (k’) respectively. These parameters determine the reflectance and absorptance of the material. Close to the plasma frequency the properties of the material changes drastically. In the infrared (IR) part of the spectrum, below this critical value (ω<ωp, or λ>λp) the imaginary part of formula (2.12) is large, and the penetrating wave drops off exponentially [38]. The real part is negative, and the material has near-unity reflectance. For ω>ωp (or λ< λp) the imaginary part tends to zero, and absorption is small. The refractive index is positive and almost constant with frequency according to:

1/ 22

1 pNω

εωε∞⎛ ⎞

= − ∞⎜ ⎟⎝ ⎠

(2.16)

Here ε∞ is the high frequency permittivityI. The TCO behaves like a dielectric and is transparent in the region for ω>ωp [39]. In this transparent regime the film is weakly absorbing (k’2<<Ν2) and the transmission can be expressed as [40]:

(1 )exp( )T R dα= − − (2.17)

I The relative permeability µ0 is set to 1, which holds for weakly magnetic materials as

most transparent semiconductors.

Chapter 2 26

R is the zero degree incidence reflectance, d is the film thickness and α is the absorption coefficient and is dependant on the wavelength according to:

4 'kπαλ

= (2.15)

Close to λgap the reflectance is zero and the absorption coefficient as a function of wavelength can be obtained easily from the transmission curve. The following relation applies for direct allowed transitions:

ghv Eα ∝ − (2.18)

where hν is the photon energy. The bandgap energy is calculated from the Tauc plot [41], α2 versus the photon energy hν. If α2 is extrapolated to the x-axis intersection (α2=0), formula 2.16 implies that the photon energy equals the band gap energy (hv≈Eg). This method is commonplace to extract bandgap energies from transmission data. 2.3.2 Correlation of optical and electrical properties

The optical parameters of TCOs are affected by the electrical properties of the material. The earlier mentioned plasma resonance frequency is not a fixed value, but varies with the electron concentration. The plasma frequency in relation to the carrier concentration is expressed by:

2

0 *pc

n em

ωε ε∞

⋅= (2.19)

At this frequency the dielectric-like visible transmittance equals the metallic-like IR reflectance (T=R). Thus the IR reflectivity of the material can be tuned, which is Important for heat reflecting or low emissive window applications. For example, the plasma wavelength of ITO films can be tuned from 1.5 µm to 4 µm due to the carrier concentration by changing the composition and deposition parameters [42,43]. The refractive index of commonly TCO materials varies in between 1.7 and 2.1 [7]. Moreover, a large spread for similar material is frequently reported. Main reason is that also the refractive index is dependant on the carrier concentration according to:

22

2

4*optNne

mπε

ω= − (2.20)


Here ω is the frequency of incoming light. Changing the refractive index is useful for waveguide applications. For instance in ITO this value can be tuned between 1.70 and 2.05 [44,45] For heavily doped oxide semiconductors a gradual shift of the bandgap towards higher energy as the electron density increases is generally observed [10, and the references therein]. This well known effect is attributed to the Burstein-Moss shift (BM shift) [46]. In heavily doped semiconductors, the lowest states in the conduction band are blocked. Hence transition can only take place to energies above EF (see figure 2.1), enlarging the effective optical gap. The energy gap between the top of valence band and lowest empty state in the conduction band (both assumed parabolic) can be given by:

0BM

g g gE E E= + ∆ − ∑ (2.21)

In this formula Eg0 is the intrinsic bandgap and the BM shift given by:

22 2 /3

* (3 )2

B Mg

VC

E nm

π−∆ = (2.22)

Here m*VC is the reduced effective mass of the electron carriers given by:

* * *

1 1 1

VC C Vm m m= + (2.23)

Where m*C and m*V are effective mass of the carriers in the conduction and valence band respectively. The term ћΣ in the equation (2.19) represents self energies due to electron-electron and electron-impurity scattering, causing a band gap narrowing that counteracts the BM shift. This effect is of importance at very high carrier concentrations (order 1021 cm-3). 2.4 General properties

The method of preparation is of great importance for growing high quality TCO films. The physical properties of TCO thin films are strongly dependant on the structure, morphology and composition of the thin films, and the nature of the impurities. These factors are influenced by the deposition parameters of the different growth techniques. For TCO thin film preparation, a wide variety of growth techniques has been reported and examined extensively.

Effective TCOs should possess both high electrical conductivity and low absorption of visible light. Thus the ratio between the transmittance (T) and sheet resistance (Rs) of transparent conducting films is a figure of merit (FOM) for rating these materials. Later Haacke defined a more suitable measure as for some

Chapter 2 28

applications optical absorption is too low at maximum FOM [47]. This new figure of merit is defined as:

10TC ST RΦ = (2.24)

Besides the electrical and optical properties, other criteria influence the choice of material and deposition method. For material processing the etching properties, production cost or toxicity are important. In view of applications the plasma frequency, mechanical hardness; thermal or chemical stability and minimum deposition temperature are essential. An overview of these criteria tailored to different application is reported by Gordon [48]. At the end of this chapter, an overview on n-type and p-type TCOs is given in table 2.1 and 2.2, respectively. The references in these tables are not selected on best performance of the TCO, but are merely a selection of the vast amount of literature on these materials. It illustrates the diversity of materials and corresponding properties deposited by different deposition methods. Among these are chemical vapor deposition (CVD), sputtering techniques (DC, RF, magnetron), evaporation (reactive, thermal or e-beam), spray pyrolysis and pulsed laser deposition (PLD). Particular deposition techniques for TCO manufacturing are chosen for several reasons as thickness uniformity, low production costs or high throughput. However, the electrical and optical performance is not directly related to the deposition method. They are more dependent on the intrinsic properties as structure, morphology and composition of the thin film. Each deposition method and conditions can influence the intrinsic behavior differently. Although each technique has its own advantages or limitations, they are all capable of tuning the intrinsic properties within a specific range in order to optimize the TCO performance. 2.5 Concluding remarks Coexistence of transparency and conductivity in semiconductors is possible if the bandgap is large enough to avoid visible light absorption (~3 eV). This gap is situated in between parabolic O 2p and metal s bands, forming the valence and conduction band respectively. Intentional doping creates free carriers (electrons or holes) which are responsible for the conductivity. Electron densities are in general high for n-type TCOs (order 1020 cm-3) resulting in degenerate electron systems. The electrical conductivity is dominated by scattering mechanisms, which are strongly related to the electron concentration. All electronic scattering contributions together influence the mobility, and determine the upper limit of conductivity. For single-crystalline materials the ionized impurity scattering is considered to be the most important. Though, as the structural nature of the thin film decrease, the


contribution of grain boundaries becomes ever more important in polycrystalline films. The transparency window for electromagnetic waves between UV and near-IR is typical for TCOs. This window is on the short wavelength side determined by the bandgap and on the long wavelength side by the plasma frequency. Both values are dependant on the electron concentration, and can therefore be tuned to serve specific applications.

The electrical conductivity of common n-type materials is around 10-4 Ωcm, whereas the transmission can be as high as 90% in the visible regime. The conductivity of p-type TCOs is in general at least a factor of 1000 lower. Although the figure of merit is a measure for the TCO electrical and optical performance, in practice the properties are tailored for the different applications. Many deposition methods can be used to grow TCOs, and varying performances are reported. Each deposition method has its own influence on the intrinsic properties as thin film structure and composition. It is these properties and not the deposition method by itself determining the electrical and optical performance of the TCO.

Chapter 2 30

Table 2.1: Reported properties of n-type transparent conducting oxides TCO

Deposition

method ρ

(Ωcm) T II %

FOM (10-3Ω-1)

Eg

(eV) N

n (cm-3)

µ (cm2/s/V)

ref.

SnO2 Spray 4.3 x10-3 97 1.45 4.11 2.2 [49] SnO2 Sputtering 6.1 x10-3 95 56.4 4.13 1.3 x1020 7.7 [50]

SnO2:F Spray 5 x10-4 >80 4.41 4.6 x1020 28 [51] SnO2:Mo Reactive ev. 3 x10-3 >85 >0.07 4.10 8.0 x1020 10 [52] SnO2:Sb Spray 10-3 85 19.6 3.75 7.0 x1020 10 [53] Cd2SnO4 Sputtering 5 x10-4 >80 2.05 5.0 x1020 40 [54] Cd2SnO4 Sputtering 5 x10-4 93 34.6 2.7 5.0 x 1020 22 [55] CdIn2O4 Sputtering 2.7 x10-4 90 69.7 3.24 4.0 x 020 57 [56]

In4Sn3O12 Sputtering 3.5 x10-4 >80 3.5 2.1 7.0 x1020 11.5 [57] In2O3 Thermal ev. 2 x10-4 >90 3.56 4.0 x1020 70 [58] In2O3 PLD 2 x10-4 86 9.0 x 020 37 [59]

In2O3:F CVD 2.9 x10-4 >85 3.9 3.5 [60] GaInO3 Sputtering 2.5 x10-3 90 14 4 x1020 10 [61]

ITO e-beam ev. 2.4 x10-4 90 3.85 2.0 8.0 x1020 30 [62] ITO CVD 1.7 x10-4 90 183 3.9 8.8 x1020 43 [63] ITO Sputtering 2,4 x10-4 95 70.4 4.0 2.0 1 x1020 12 [64] ITO PLD 8.5 x10-5 85 72.9 1.4 x1021 53.5 [65] ITO Sol-gel 5.0 x10-3 1.9 x1020 12 [66]

ITO:F Sputtering 6.7 x10-4 >80 3.5 6.0 x1020 16 [67] In2O3:Mo Sputtering 5.9 x10-4 90 7.7 4.3 5.2 x1020 20.2 [68]

ZnO Reactive ev. 10-3 88 3.3 1 x1020 10 [69] ZnO Sputtering 2 x10-3 >80 4 1.2 x1020 16 [70]

ZnO:Al Sputtering 10-2 90 3.52 4.7 x1020 1.47 [71] ZnO:Al CVD 3.3 x10-4 85 49.2 8.0 x1020 35 [72] ZnO:Al PLD 3.7 x10-4 90 28.3 3.8 1.98 8.0 x1020 18 [73] ZnO:Ga Sputtering 10-3 >85 3.59 10 x1020 10 [74] ZnO:In Sputtering 2 x10-2 >80 3.29 1.85 7 x1019 1.9 [75]

Zn3In2O6 PLD 1.0 x10-3 85 3.4 4.0 x1020 20 [76] ZnSnO3 Sputtering 4 x10-3 >80 1 x1020 10 [77]

Table 2.2: Reported properties of p-type transparent conducting oxides

TCO

Deposition method

σIII (Scm-1)

T %

Eg

(eV) S

µV/K Ea

(eV) n

(cm-3) µ

(cm2/s/V) ref.

CuAlO2 PLD 0.34 70 3.5 +214 0.22 2.7 x1019 0.13 [78] CuAlO2 CVD 2 <70 3.75 0.12 1.8 x1019 0.16 [79] SrCu2O2 PLD 4.8 x10-2 75 3.3 +260 0.10 6.1 x1017 0.46 [80]

CuYO2:Ca Thermal ev. 1.0 50 3.5 +275 0.13 >1.0 [81] AgCoO2 Sputtering 0.2 50 4.15 +220 0.07 [82] CuGaO2 PLD 6.3 x10-2 80 3.2 +560 1.7 x1018 0.23 [83] ZnO:P Sputtering 1.7 3.35 1.0 x1017 0.53 [84] ZnO:N PLD 0.5 85 + 6 x1018 0.1 [85]

II In some cases the authors use the transmission at the fixed wavelength of 500 nm

instead of an average over the range 400-700 nm. III In literature on p-type TCOs the electrical performance is often represented by the

conductivity (σ) instead of the resistivity (ρ) generally used for n-type materials.


2.6 References 1. C.G. Granqvist, A Hultåker, Thin Solid films, 411, 1, (2002) 2. K. Badeker, Ann. Phys. 22, 749 (1907) 3. A. Thelen, H König, Naturwissenschaften, 43, 297, (1956) 4. R. Groth, E Kauer, Philips Technical Review, 26, 105, (1965) 5. D.S. Ginley & C. Bright, MRS Bull. 25, 15 (2000). 6. K.L. Chopra, S. Major & D.K. Pandya, Thin Solid Flms 102, 1 (1983). 7. H.L. Hartnagel, A.L. Dawar, A.K. Jain & C. Jagadish, Semiconducting

Transparent Thin Films (Institute of Physics Publishing, Bristol and Philadelphia, 1995).

8. N.F. Mott, Can. J. Phys, 34, 1356 (1956). 9. N.F. Mott, Philos. Mag. 6, 284 (1961). 10. I. Hamberg & C.G. Granqvist, J. Appl. Phys. 60, R123 (1986). 11. H. Köstlin, R. Jost & W. Lems, Phys. Status Solidi A 29, 87 (1975). 12. J.S. Blakemore, J. Appl. Phys. 53, R123 (1982). 13. M. Chen, Z.L. Pei, X. Wang, Y.H. Yu, X.H. Liu, C. Sun & L.S. Wen, J. Phys.

D: Appl. Phys. 33, 2538 (2000). 14. G. Frank & H. Köstlin, Appl. Phys. A 27, 197 (1982). 15. H. Brooks, in Advances in electronics and electron physics, Chapter 7, p85,

(1955). 16. R.B. Dingle, Philos. Mag. 46, 831 (1955). 17. E.J. Moore, Phys. Rev. 160, 618 (1967). 18. T. Pisarkiewicz, K. Zakrzewska & E. Leja. Thin Solid Films 174, 217 (1989) 19. J. Y. W. Seto, J. Appl. Phys. 46, 5247 (1975). 20. R.L. Petritz, Phys. Rev. 104, 1508 (1956). 21. A.K. Kulkami & S.A. Knickerbocker, J. Vac. Sci. Technol. A 14, 1709 (1996). 22. J. Bruneaux, H. Cachet, M. Froment & A. Messad, Thin Solid Films 197, 129

(1991). 23. K. Ellmer, J. Phys. D: Appl. Phys. 34, 3097 (2001). 24. C. Kittel, in Introduction to solid state physics, 5th edition, Maruzen, Tokyo,

(1985), Chap. 6. 25. G. Frank & H. Köstlin, Appl. Phys. A 27, 197 (1982). 26. C. Erginsoy, Phys. Rev. 79, 1013 (1950). 27. S. Noguchi & H. Sakata, J. Phys. D: Appl. Phys. 14, 1523 (1981). 28. Y. Shigesato & D.C. Paine, Appl. Phys. Lett. 62, 1288 (1993). 29. T. Minami, MRS Bull. 38, 15 (2000). 30. T.J. Coutts, J.D. Perkins, D.S. Ginley & T.O. Mason, Conference Paper,

NREL/CP-520-26640, (1999).

Chapter 2 32

31. A.J. Freeman, K.R. Poeppelmeier, T.O. Mason, R.P.H. Chang & T.J. Marks,

MRS Bull. 25, 45 (2000). 32. H. Hosono, N. Kikuchi, N. Ueda & H. Kawazoe, J. Non-Cryst. Solids 198-200,

165 (1996). 33. M. Orita, H. Ohta & M. Hirano, Phil. Mag. B 81, 501 (2001). 34. K. Nomura, H. Ohta, A. Takagi, T. Kamiya, M. Hirano & H. Hosono, Nature

432, 488 (2004). 35. F. Simonis, M. van der Leij & C.J. Hoogendoorn, Solar Energy Mater. 1, 221

(1979). 36. P. Drude, Ann. Phys. 3, 369 (1900). 37. R.G. Gordon, MRS Bull. 25, 52 (2000). 38. E. Hecht, in Optics, 2nd edition, Addison-Wesley Publishing Company, Inc.

(1990). 39. T.J. Coutts, D.L. Young & X. Li, MRS Bull. 25, 58 (2000). 40. J.C. Manifacier, M. De Murcia, J.P. Fillard, E. Vicario, Thin Solid Films 41,

127 (1977). 41. J. Tauc, R. Grigorovici & A. Vancu, Phys. Status Solidi 15, 627 (1966). 42. P.K. Biswas, A. De, N.C. Pramanik, P.K. Chakraborty, K. Ortner, V. Hock &

S. Korder, Mater. Lett. 57, 2326 (2003). 43. H. Kim, J.S. Horwitz, A. Piqué, C.M. Gilmore & D.B. Chrisey, Appl. Phys. A

69, S447 (1999). 44. R.T. Chen & D. Robinson, Appl. Phys. Lett. 60, 1541 (1992). 45. Y. Ohhata, F. Shinoki & S. Yoshida, Thin Solid Films 59, 255 (1979). 46. E. Burstein, Phys. Rev. Lett. 93, 632 (1954). 47. G. Haacke, J. Appl. Phys. 47, 4086 (1976). 48. R.G. Gordon, MRS Bull. 25, 52 (2000). 49. H.H. Afify, R.S. Momtaz, W.A. Badawy & S.A. Nasser, J. Mater. Sci.: Mater.

Electron. 2, 40 (1991). 50. A. De & S. Ray, J. Phys. D: Appl. Phys. 24, 719 (1991). 51. E. Santhani, A. Banerjee, V. Dutta & K.L. Chopra, J. Appl. Phys. 53, 1615

(1982). 52. V. Casey & M.I. Stephenson, J. Phys. D: Appl. Phys. 23, 1212 (1990). 53. H. Kaneko & K. Miyake, J. Appl. Phys. 53, 3629 (1982). 54. S. Schiller, G. Beister, E. Buedke, H.J. Becker & H. Schicht, Thin Solid Films

96, 113 (1982). 55. N. Miyata, K. Miyake, K. Koga & T. Fukushima, J. Electrochem. Soc. 127,

918 (1980). 56. K. Budzylqska, E. Leja & S. Skrkzypek, Sol. Energy Mater. 12, 57 (1985).


57. T. Minami, T. Kakumu, K. Shimokawa & S. Takata, Thin Solid Films 317, 318

(1998). 58. C.A. Pan, T.P. Ma, J. Electron. Mater. 10, 43 (1981). 59. F.O. Adurodija, H. Izumi, T. Ishihara, H. Yoshioka & M. Motoyama, J. Vac.

Sci. Technol. A 18, 814 (2000). 60. T. Maruyama & K. Fukui, Jpn. J. Appl. Phys. 29, L1705 (1990). 61. J.M. Phillips, J. Kwo, G.A. Thomas, S.A. Carter, R.J. Cava, S.Y. Hou, J.J.

Krajewski, J.H. Marshall, W.F. Peck, D.H. Rapkine & R.D. van Dover, Appl. Phys. Lett. 65, 115 (1994).

62. S.A. Agnihotry, K.K. Saini, T.K. Saxena, K.C. Napgal & S. Chandra, J. Phys. D: Appl. Phys. 18, 2087 (1985).

63. T. Maruyama & K. Fukui, J. Appl. Phys. 70, 3848 (1991). 64. K. Screennivas, T. Sudersena Rao & A. Mansingh, J. Appl. Phys. 57, 384

(1985). 65. A. Suzuki, T. Matsushita, T. Aoki, A. Mori & M. Okuda, Thin Solid Films 411,

23 (2002). 66. P.K. Biswas, A. De, N.C. Pramanik, P.K. Chakraborty, K. Ortner, V. Hock,

S. Korder, Mater. Lett. 57, 2326 (2003). 67. C. Geoffroy, G. Campet, J. Portier, J. Salardenne, G. Couturier, M. Bourrel,

J.M. Chabagno, D. Ferry & C. Quet, Thin Solid Films 202, 77 (1991). 68. W-n. Miao, X-f. Li, Q. Zhang, L. Huang, Z.-j. Zhang, L. Zhang & X.-j. Yan,

Thin Solid Films 500, 70 (2006). 69. G.H.G. Swamy & P.J. Reddy, Semicond. Sci. Technol. 5, 980 (1990). 70. J.B. Webb, D.F. Williams & M. Buchanan, Appl. Phys. Lett. 39, 640 (1981). 71. S. Ghosh, A. Sarkar, S. Bhattacharya, S. Chaudhuri & A.K. Pal, J. Cryst.

Growth 108, 534 (1991). 72. J. Hu & R.G. Gordon, J. Appl. Phys. 71, 880 (1991). 73. H. Kim, U,A. Pique, J.S. Horwitz, H. Murata, Z.H. Kafafi, C.M. Gilmore & D.B.

Chrisey, Thin Solid Films 377-378, 798 (2000). 74. B.H. Choi, H.B. Im, J.S. Song & K.H. Yoon, Thin Solid Films 193-194, 712

(1990). 75. A. Sarkar, S. Ghosh, S. Chaudhuri & A.K. Pal, Thin Solid Films 204, 255

(1991). 76. N. Naghavi, A. Rougier, C. Marcel, C. Guéry, J.B. Leriche & J.M. Tarascon,

Thin Solid Films 360, 233 (2000). 77. T. Minami, H. Sonohara, S. Takata & H. Sato, Jpn. J. Appl. Phys. 33, L 1693

(1994). 78. H. Yanagi, S-i Inoue, K. Ueda, H. Kawazoe, H. Hosono & N. Hamada, J.

Appl. Phys. 88, 4159 (2000).

Chapter 2 34

79. H. Gong, Y. Wang & Y. Luo, Appl. Phys. Lett. 76, 3959 (2000). 80. A. Kudo, H. Yanagi, H. Hosono & H. Kawazoe, Appl. Phys. Lett. 73, 220

(1998). 81 M.K. Jayaraj, A.D. Draeseke, J. Tate & A.W. Sleight, Thin Solid Films 397,

244 (2001). 82. J. Tate, M.K. Jayaraj, A.D. Draeseke, T. Ulbrich, A.W. Sleight, K.A. Vanaja,

R. Nagarajan, J.F. Wager & R.L. Hoffman, Thin Solid Films 411, 119 (2002). 83. K. Ueda T. Hase, H. Yanagi, H. Kawazoe, H. Hosono, H. Ohta, M. Orita, & M.

Hirano, J. Appl. Phys. 89, 1790 (2001) 84. K.-K. Kim, H.-S. Kim, D.-K. Hwang, J.-H. Lim & S.-J. Park, Appl. Phys. Lett,

83, 63 (2003). 85. X.-L. Guo, H. Tabata & T. Kawai, Opt. Mater. 19, 229 (2002).

Chapter 3 Fabrication and characterization of transparent semiconductors

Abstract

Pulsed laser deposition is a powerful tool for thin film research. A large freedom of choice in independently controllable deposition parameters allows the researcher to quickly obtain results on the exploration and optimization of new materials. The ablated species can be tuned over a large energetic range, enabling optimum conditions at lower substrate temperatures. This makes the deposition of TCOs on heat resistive substrates possible. The polymer substrates in this research are commercial available and commonly used materials. Analysis of the PLD grown films is done by a variety of measuring tools to obtain information on the electrical, optical and structural properties as well as thin film composition.

Chapter 3 36

3.1 Introduction Chemical vapor deposition (CVD), spray and sputter techniques are used

nowadays to coat large areas of TCOs on glass for commercial applications. In the emerging field of transparent conductors the need for new materials or deposition on unconventional substrates as plastics is needed. Pulsed laser deposition (PLD) is (yet) not a very suitable technique for large area applications. However, it is a powerful tool in the field of thin film research and offers many advantages compared to conventional deposition techniques. In the last decades PLD has been successfully used to grow high-Tc superconductors and semiconductors thin films of excellent quality [1,2]. PLD is a physical vapor deposition technique based on the evaporation of material by an intense laser beam. The vaporized material forms a plume, and condenses on a substrate which is placed in the vacuum chamber opposite to the target. A thin film will form on this substrate, owing its properties to a wide variety of process parameters.

The combination of polymer materials with functional or protective coatings offers a number of advantages above conventional glass, for example light-weight, flexibility and cost-effectiveness. Moreover, polymers can match the optical qualities of glass in the visible part of the spectrum. Once the TCO thin film is fabricated, the electrical and optical performance must be tested. In order to improve or influence the TCO performance, knowledge on structural, compositional and morphological properties are essential. The use of dedicated analysis tools is therefore indispensable. In this chapter the pulsed laser deposition method used for growing TCOs will be addressed. An overview of the analytical tools used to characterize bulk or thin film materials are discussed briefly in section 3.3. In section 3.4 some aspects of the (polymer) substrates will be described, 3.2 Thin film growth by pulsed laser deposition 3.2.1 Basic principles

In the deposition process a pulsed laser beam is focused on a piece of source material, in general referred to as the target. The focusing of the laser beam results in a locally very high energy density (fluency) on the target surface. The electromagnetic energy is converted into thermal energy via electronic processes, and results in evaporation of the source material [3]. This process takes place on a very short timescale, shorter than the laser pulse duration (nanoseconds). Therefore, also laser energy will be absorbed from the incident beam by the evaporated material close to the surface. Extremely high temperatures are reached resulting in high density plasma that starts to expand due to the pressure gradient close to the target surface. The ablated species move in a forward direction

Fabrication and characterization of transparent semiconductors 37

perpendicular to the target surface into the vacuum chamber. These high energetic particles interact with the background gas and emit photons in this process, resulting in the characteristic luminous plasma plume. The surrounding gas pressure influences the mean free path of the species in the plume, and determines their deceleration. Control of this background gas pressure allows the kinetic energy of the ablated species to be tuned from high values (~100 eV) in vacuum to low energies (~0.1 eV) at high background pressures [4]. The poor thickness uniformity of PLD grown thin films, especially important in transparent electrodes for displays, is a disadvantage. Therefore only small substrates can be used (typical dimensions 10x10 mm2). From a research point of view however, the exploration of (transparent conducting) oxide materials can benefit from the different aspects typical for PLD thin film growth: 1. The kinetic energy of the particles arriving at the substrate can be tuned by

changing the ambient gas parameters; mass and pressure. In this way the thin film growth can be modified since this parameter affects the diffusivity and absorption and desorption probability of the energetic particles at the film surface [5]. Furthermore, the presence of high-energy species offers the possibility to deposit crystalline films at lower deposition temperatures [6], beneficial for deposition on polymer substrates.

2. Stoichiometric transfer from target to substrate. Ablation of a wide range of materials is possible from multi-component targets.

3. The decoupling of the vacuum hardware and the energy source inducing evaporation makes pulsed laser a flexible deposition system. The process parameters can be controlled almost independently from each other and is possible over a wide range. Switching to different operational modes and optimization of film growth is easy and can be done quickly.

4. Compared to other deposition techniques, the deposition rate is high. The rate can be as high as 1 Ǻ/pulse. The growth takes place in between subsequent pulses at a much shorter time scale than the laser frequency. Actual growth rate is therefore determined by the laser frequency.

3.2.2 Experimental setup The experimental set-up is schematically represented in figure 3.1. The

energy source used in the experiments is a KrF (λ=248 nm) excimer laser (Lambda Physic, Compex 205) capable of operating at frequencies between 1-50 Hz. The pulse width is about 25 ns (FWHM) and the maximum pulse energy is 650 mJ. A mask (maximum dimension 102 mm2) is placed in the beampath to select the center part of the beamprofile where the energy spatial variation is below 10%. The laser beam is focused by a lens (focal length 453 mm at 248 nm), passes through the laser window and is projected on the target at 45° incidence. The spotsize is in

Chapter 3 38

the order of a few square millimeters and can be changed by adjusting the lens and mask position. The energy is measured behind the lens and, by taking into account the energy loss due to the laser window, the fluency on the rotating target can be calculated accordingly. Measured pulse-to-pulse energy variation was below 5% at any time. Prior to deposition, the vacuum chamber is pumped down to a base pressure of 10-6 mbar. Oxygen and argon can be used as deposition gas, and enter the vacuum chamber by mass-flow controllers (0-40 ml/min.). The pressure during deposition is controlled by the combination of flow rate and the pump speed. A butterfly control-valve adjusts the effective pump speed through a small restriction, bypassing the closed main valve. This allows the deposition pressure to be adjusted in between 10-3-100 mbar with an accuracy of 0.001 mbar, which is the range normally used to grow oxide materials. The PID controlled temperature of the heater can adjust the substrate temperature from room temperature (RT) up to 950 °C. The heater stage is able to move and allows a target to substrate distance between 35-70 mm. A maximum of

Figure 3.1: Schematic view of the pulsed laser deposition setup, and a photograph of the actual systemand laser (background) used in the experiments.

to pump

vacuumdepositionchamber

target

gas inlet

plasma

lens mask

substrateon heater

KrF excimer laser( = 248 nm)λ

main valvecontrol valve


four targets can be mounted in the system. Multilayer stacks can be deposited without breaking the vacuum. 3.3 Thin film characterization 3.3.1 Electronic transport analysis

Room temperature measurements of the electrical resistivity and Hall mobility were performed using a standard van der Pauw four-point probe technique [7]. Thin films were either patterned by shadow mask deposition, or wiring was attached to the outermost corner of the square samples to obtain best results [8,9 10]. In general the electrical connections are either wire bondedI, or attached to the film surface by conducting silver glueII. In some cases probes were pressed directly to the film surface.

The electrical measurement setup consist of a Keithley Instruments model 6221 current source, model 7001 switchboard equipped with the 7065 Hall-effect card with unity-gain buffers and a 2182A nanovoltmeter. This setup allows measurements on low as well as high resistive samples. In order to perform Hall measurements, the samples can be placed in a permanent magnetic field. The field in the middle between the two 5.08 cm diameter permanent magnets was calibrated at 720 mT. The field was found to be homogeneous within 10 mT over a 1x1 cm2 area. The obtained data provides information concerning carrier type and the carrier concentration according to:

HI BV

q n d×

=⋅ ⋅

(3.1)

If a current I is applied to a sample with thickness d in a perpendicular magnetic field B, the moving charge carriers are deflected by the Lorenz force. A potential difference VH will build up perpendicular to both the current flow and magnetic field. The accumulation of charge continues until the counteracting electrostatic force cancels the magnetic force on the moving charges. From the polarity of the potential difference VH it is possible to distinguish between electrons or holes as the majority of charge carriers. If VH is plotted against B, the carrier type and carrier density can be deduced from the sign and slope of the graph, respectively. Combined with the resistivity data, the carrier mobility can be calculated from equation 2.1.

I Aluminum wire contacts are ultrasonically attached to the film surface by a wire-

bonder, Mech-el Industries Inc., Wolburn, Massachusetts, USA. II Silver filled epoxy EE129-4, Epoxy Technology, Inc. Billerica, USA.

Chapter 3 40

Temperature dependant conductivity measurements were performed in a physical property measuring system (Quantum Design 6000). The temperature range can be varied from 300 K down to 4 K. The superconducting magnets in this system can produce an adjustable magnetic field up to 9 Tesla, allowing Hall measurements at very low temperatures.

Thermo-power or Seebeck effect measurements are useful in determining the type of charge carriers in a semiconductor (p or n). Electrons (or holes) will thermally diffuse to one end of the material to the other if a temperature gradient is applied to it. The electric field due to the difference in charge carriers on both sides of the sample induces an electric current counteracting the thermal diffusion current. In equilibrium a potential difference between both sides of the sample can be measured and is known as the Seebeck voltage. The Seebeck coefficient is defined as the ratio between this voltage difference and temperature difference on both sides of the sample:

( , *)nVS f e mT

∆= =∆

(3.1)

The sign of the Seebeck coefficient determines whether electrons (negative) or holes (positive) are the major charge carriers. Furthermore, the magnitude of this coefficient is exponentially dependant on the amount of charge carriers [11]. This value is in the order of µV·K-1 for metals and mV·K-1 for semiconductors.

For the thermo-power experiments the sample is brought into thermal contact with a glass slide that is placed on a hot plate by one end and on a metal heat sink by the other end. If the hot plate is turned on, the glass slide will warm up, whereas the other side remains at a constant temperature. A temperature gradient is generated over the glass slide, and accordingly over the length of the sample. The temperature on both ends of the sample is monitored by thermocouples. If the system is in equilibrium, the potential difference over both sides of the sample is measured using a nanovoltmeter. 3.3.2 Optical analysis

Optical transmission and reflection data was collected with a UV-Visible spectrophotometer (Cary 50, Varian Instruments), capable of measuring in the wavelength range 190-1100 nm. This apparatus uses a Czerny-Turner 0.28m monochromator and a double silicon diode detector. A xenon flashlight is used as illumination source, and the dual beam operation mode is able to correct for intensity fluctuations of the source our surroundings. The resolution and scan speed can be adjusted. In most experiments the wavelength interval was 1 nm measured at 10 nm/s. A baseline can be recorded in order to use the bare substrate as a


reference. By using the 45º incidence specular reflectance attachment, reflection mode data can be recorded.

From the obtained optical data, values for the absorption can be deduced, and subsequently the bandgap value can be obtained as described in section 2.3.1). Moreover, the film thickness can be calculated for known refractive index from a fringe pattern, or vice versa. 3.3.3 Structural analysis

Structural characterization of thin films and deposition targets was done by x-ray diffraction (XRD). Measurements were performed on a multi-purpose diffractometer (X’Pert MPD, Philips, The Netherlands) as well as a four-circle single crystal diffractometer (CAD4, Enraf Nonius, Delft, The Netherlands), both using Cu-Kα radiation. The X’Pert system uses an x-ray source operating at 50 kV and 35 mA, and the minimum resolution is as small as 0.001º. The system is equipped with automatic slits to control the illuminated and observed length on the sample (0.5-20 mm). A graphite monochromator is used in the diffracted beampath, and eliminates high-order reflections of crystalline substrates. The CAD4 system operates at 50 kV and 26 mA maximum, the minimum resolution is 0.01º. The circular spot has a diameter of about 1 mm on the sample (angle dependant). This four circle diffractometer is capable of measuring asymmetric scans on epitaxial film samples. Furthermore, reciprocal space maps can be recorded in order to measure in-plane lattice parameters. Most performed XRD scans are symmetric ω-2θ scans; i.e. in (00l)- direction where the diffraction vector is perpendicular to the film surface. A small offset in ω is used if epitaxial films are measured to compensate for the vicinal angle of the crystalline substrate. In some cases the polycrystalline samples are measured by a constant grazing incident angle (2θ scan). This geometry increases the effective x-ray pathlength through the film, whereas the diffracted intensity of the substrate is reduced. Overall, the thin film signal-to-noise ratio increases significantly.

The X’Pert Highscore plus software package (PANalytical B.V., The Netherlands) coupled to the ICDD databaseIII is used to analyze the data and fit the obtained peak profile. Furthermore the grain size can be obtained from the spectral line broadening. A Gaussian function (FWHM) fitted to the structural broadened peaks of polycrystalline films [12]. These calculations are based on the Scherrer’s formula [13], with the instrumental broadening taken as 0.03º (X’Pert system).

III Electronic database of the International Centre for Diffraction Data.

Chapter 3 42

Transmission electron microscopy (TEM) images were taken on a Philips CM30T-LaB6, and upgraded CM300ST-FEGIV. The maximum acceleration voltage is 300 kV, and the point-to-point resolution is 2.3 Å. Next to high resolution images the diffraction patterns can be imaged to determine the crystalline nature of the samples. Analysis on amorphous and polycrystalline samples was mainly performed on thin films of about 25 nm thickness deposited on copper grids covered with a layer of carbon. 3.3.4 Morphology

Atomic Force Microscopy (AFM) measurements are performed on a Nanoscope IV instrument (Veeco, Digital Instruments, Santa Barbara, USA) capable of measuring in contact or tapping mode. The resolution is dependant on tip shape, sample roughness and instrumental settings, but is in general sub-Angstrom in z-direction (height). From the topographical data the surface roughness can be determined.

In-situ AFM measurements were also performed during the growth of thin films by PLD. In contrast to in-situ RHEED this novel experimental technique, which is still under development, allows growth monitoring in real-space instead of its reciprocal counterpart [14]. From the start of a bare substrate, the morphological evolution of the thin film can be checked in between the separate laser pulses. The AFM-PLD principle works at deposition conditions up to a temperature of 530 °C.

Scanning Electron Microscopy (SEM) surface analysis was carried out on a JSM-5610 (Jeol, Tokyo, Japan) operating between 0.5 and 30 keV. The lateral resolution is 3.5 nm at 30 keV and a working distance of 6 mm. The system is not only capable of imaging the topography by secondary electrons, but also measures backscattered electrons adding information on the chemical composition.

Film thickness determination was performed using a stylus profilometer (Alpha Step 250, Tencor Instruments). The stylus measures the step height of the film edge relative to the substrate. The film edge is either created by shadow mask deposition, or photolithographic patterning (lift-off or back-etch technique). 3.3.5 Compositional analysis

X-ray Photoelectron Spectroscopy (XPS) data is collected by a PHI Quantera Scanning ESCA MicroprobeIV. This analysis tool is based on the photo electrical effect. For each and every element, there will be a characteristic binding energy associated with each core atomic orbital. Each element will give rise to a characteristic set of peaks in the photoelectron spectrum at kinetic energies

IV These measurements are performed at the Central Materials Analysis Laboratory

(Mesa+, Twente University, The Netherlands).


determined by the photon energy and the respective binding energies. From this peak intensity ratio and positions information on the composition and chemical bonding of a material can be retrieved.

Monochromatic scanning soft x-rays (Al-Kα, 0-1487 eV) are used as excitation source. The compositional detection limit is ~0.1 %at. An argon-ion gun can be used for charge neutralization on the film surface. The technique is surface sensitive, but by using a sputter source chemical depth-profiling of the sample is possible. Furthermore, low energetic x-rays (0-60 eV) are used to examine the valence level.

XRF Philips PW1480 with Rh tube is used to determine the chemical composition of deposition targetsV. In contrast to XPS, this technique is not surface sensitive and is therefore ideal to examine bulk materials. The detection limit is dependant on the element to be analyzed and the sample matrix, but is in the order of 0.01%. 3.4 Substrates 3.4.1 Polymers Polymers replacing glass should obviously have a significant transmission in the visible part of the spectrum. In table 3.1 the thermal and mechanical properties of some transparent polymers are listed. Major differences can be observed in the heat sustainability and mechanical strength. An important parameter is the glass transition temperature (Tg), above this temperature the mechanical properties (elastic modulus) is strongly dependant on the temperature. It is therefore not desirable to exceed this temperature significantly during deposition conditions, as this may lead to deformations of the polymer surface and film.

A disadvantage of polymers is that they can absorb a substantial amount of water. The water uptake of the listed materials in table 3.1 varies up to a maximum of 0.4 %wt. [15]. When the substrate is brought under vacuum conditions for deposition, it starts to dry. Furthermore it is evaporated under exposure of the first plasma pulses. The remaining water on the surface can result in formation of hydroxides at the interface, but is not expected to influence the film adhesion [16].

Although all materials in table 3.1 are potential candidates, not all are commonly used. PEEK for instance is a high performance plastic used in applications where mechanical strength, high temperature and chemical resistance are of great importance. However, the price can be up to 100 times that of PET and it is only frequently used in the automotive and aerospace industry. On the other V XRF measurements are performed at the Catalytic Processes and Materials group

(Faculty of Science & Technology, Twente University, The Netherlands).

Chapter 3 44

hand PET, PC and PMMA are used in a broad range of applications as plating or packaging. The substrates used in this research are samples of commonly commercial available polymer sheets directly obtained from manufacturers. Several of these polymers were obtained from different suppliers. Table 3.1: Overview of the properties of transparent polymersVI.

Polymer Abbr. ρd

g/cm3

Tg

ºC

Tm

ºC

λc

W/m·K

αl

10-5 K

cp

kJ/kg·K

Eγ

MPa

St

MPa

Poly Ethylene Terephthalate PET 1.34 75 255 0.24 7 1.0 2950 65

Poly Carbonate PC 1.22 150 - 0.21 6.5 1.31 2100 62

Poly Methyl Methacrylaat PMMA 1.18 110 - 0.19 7 1.45 3300 65

Poly Ethylene Naphthalate PEN 1.36 80 265 . 2 . . 74

Polyvinylchloride PVC 1.38 87 - 0.16 8 0.9 3150 65

Polyamide PA 1.13 50 260 0.20 8.5 1.7 1850 57

Polyether-etherketon PEEK 1.30 143 334 0.25 4.7 . 3700 90

From these materials, PET was selected for (most of) the experiments.

PMMA is quite brittle, which is a problem during material processing. The optical properties of PC on the other hand were found to be of less quality (see figure 3.2). The onset of adsorption at higher wavelength may obstruct the measurements of the thin film optical properties. PEN was not used as the thin substrates are non-rigid, preventing easy handling during the experiments. Considering the moderate thermal and mechanical properties of PET, successful film growth on this material is expected to be applicable for all polymers in table 3.1.

Unless stated otherwise, Vivak® PET substrates form Kubra BV are used in the experiments. This high transparent material is amorphous (A-PET), and from AFM measurements the surface roughness was found to be very smooth (RMS roughness <1 nm). The sheets are supplied with protective foils on both sides, preventing most substrate surfaces from damage or contamination. One large sheet with a thickness of 1 mm was cut to 10x10 mm2 sized samples. The substrates used in this research all originated from this single batch. Prior to deposition the polymer substrates are ultrasonically cleaned in ethanol, and subsequently blown

VI Values obtained from industrial technical datasheets and reference 15. Some

numbers cover a certain range from which the average is listed. Listed properties are; density (ρd), glass-transition temperature (Tg), melt temperature (Tm), thermal conductivity (λc), linear expansion coefficient (αl), heat capacity (cp), elastic modulus (Eγ) and tensile strength (St).


dry with nitrogen. PET however, can be quite soft, and vulnerable to scratching. After the cleaning process, the individual substrates were checked for damage under a microscope prior to use.

Table 3.2: Overview of the supplied polymers

Polymer Supplier Commercial

Name

Thickness

PET Kubra BV Vivak® 1 mm

DSM Arnite A02® 0.75 mm

Schreinemacher Ertalyte® 1 mm

Dupont Teijin Films Melinex® 0.125-0.250 mm

PC Schreinemacher Lexan® 1 mm

Teijn Chemicals LTD. Panlite® 1mm

Barlo Plastic Barlo PC® 1-2 mm

PMMA Barlo Plastic Barlo XT® 1 mm

Merrem Materials BV n.a. 1.5 mm

PEN Dupont Teijin Films Teonex® 0.3 mm

Kaladex® 0.125 mm

3.4.2 Mechanical properties of metal-oxides on polymers

Adhesion is fundamentally a surface property, often governed by a layer of molecular dimensions [17]. Sometimes it is therefore useful to modify the polymer surfaces for improved properties. Over the years, several methods have been developed for surface cleaning, strengthening, adhesion, wettability and other technologically important characteristics. These include mechanical treatments, wet-chemical treatments, exposure to flames, corona discharges, UV and plasma treatments [18,19.20]. Although it has been shown that pre-treatment of the PET and PC substrates can be beneficial for the mechanical properties [21,22,23], this research is focused on direct deposition without additional treatment steps.

Besides the adhesion, stress build-up during deposition is an important issue and can lead to crack onset or delamination of the thin film. The resultant stress in a thin film is the summation of the intrinsic, thermal and external stress [24]. The intrinsic stress is quite insensitive to the substrate used in case of polycrystalline films (no epitaxial relation). The thermal stress is dependant on the different expansion coefficients between film and substrate, which is quite large between oxides and polymer. To reduce this stress component, deposition at room temperature is favored. However, in many thin film fabrication methods the substrate temperature rises as a result of the deposition process if no external

Chapter 3 46

cooling is applied. This is a known issue for sputter deposited films at room temperature, whereas the temperature rise in during PLD growth is minimal [25].

However, the mechanical properties of the polymer-metal-oxide interface are not the scope of this research. Extensive research on these interfaces has been published elsewhere [22,26,27]. Moreover, the adhesion of metal-oxides as TCOs seems not to be a technological issue in this research. The mechanical properties of ITO films on polymer substrates for example are in general found to be very good [28,29]. It has even been shown that the adhesion of ITO films on polymeric substrates is better than on glass substrates [30]. Qualitative adhesion test measurements on some samples are performed and did not result in crack formation or delaminationVII. 3.4.3 Non-polymer substrates Other than transparent polymers, depositions are also performed on glass and quartz for comparison, and to allow thin film growth at higher temperatures. Glass substratesVIII are only used for experiments up to 300 °C to prevent

VII Crack-free films on polymers were tested with scotch tape peel test or multiple-

bending experiments (on thin substrates). VIII Glass substrates are cut from microscope slides (Menzel Glaser GmbH & Co,

Germany).

Figure 3.2: Transmission spectra of optical transparent polymer, amorphous and crystalline substrates.The thickness of all substrates is 1 mm.

200 300 400 500 600 700 800 900 10000

20

40

60

80

100

PET glass quartz YSZ PC

Tran

smis

sion

(%)

wavelength (nm)


outgassing of sodium [31]. Quartz substratesIX are used for experiments at higher temperatures. Moreover, because of its excellent ultra-violet (UV) transmission, quartz is a suitable substrate for optical measurements on materials in the high energetic part of the photometric spectrum. The transmission spectra of different optical substrates are plotted in figure 3.2.

Crystalline substrates are used in some of the experiments. In combination with a high substrate temperature an epitaxial relation can be obtained between the thin films and substrate. Furthermore the use of yttrium stabilized zirconia (YSZ) or sapphire (Al2O3) allows optical measurements since they are transparent in the visible part of the spectrum. Another reason for using crystalline substrates is the requirement of conducting or ultra-flat surfaces. Silicon and (niobium doped) strontium titanate (SrTiO3) can fulfill these conditions. 3.5 Concluding remarks PLD is a physical vapor deposition technique that uses an intense laser beam to evaporation the target material that condenses on a substrate. Most process parameters can be independently controlled over a large range. Optimization of thin film growth is done fast and easily. This makes PLD a powerful research tool for the exploration of (new) TCOs and their fundamental properties. Moreover, the tune-ability of ablated species PLD is a suitable method for depositing these materials on polymer substrates.

PET was used as polymeric substrates in most of the experiments. Besides the excellent optical properties, this low-cost material is commonly used in several large-scale applications. Furthermore, because of the rather standard properties of this material, the deposition of thin TCO films with similar properties is possible on other type of polymers. The performance of grown TCO materials is tested by several electrical measuring techniques. Besides the (temperature dependant) resistivity measurements, Hall and thermo power measurements are used to obtain information on carrier density, mobility and type (p or n). Photo spectroscopy was used for optical analysis. Data on structure, morphology and composition was obtained by XRD, TEM, SEM, AFM and XPS.

IX Quartz substrates are slides of amorphous fused silica, supplied by SurfaceNet

GmbH, Germany.

Chapter 3 48

3.6 References 1. D. Dijkkamp, T. Venkatesan, X.D. Wu, S.A. Shaheen, N. Jiswari, Y.H. Min-

Lee, W.L. Mclean & M. Croft, Appl. Phys. Lett. 51, 619 (1987). 2. H.S. Kwok, P. Mattocks, L. Shi, X.W. Wang, S. Witanachchi, Q.Y. Ying, J.P.

Zheng & D.T. Shaw, Appl. Phys. Lett. 52, 1825 (1988). 3. D.B. Chrisey & G.K. Hubler, In Pulsed Laser Deposition of Thin Films, John

Wiley & Sons Inc., New York (1994) 4. P. te Riele, F. Vroegindeweij, A. Janssens, F. Roesthuis, G. Rijnders &

D.H.A. Blank, in Nevac Blad, Dutch Vacuum Society, issue 3, p71 (2006). 5. D.H.A. Blank, G.J.H.M. Rijnders, G. Koster & H. Rogalla, Appl. Surf. Sci. 138-

139, 17 (1999). 6. M. Okuyama, H. Sugiyama & M. Noda, Appl. Surf. Sci. 154-155, 411 (2000) 7. L.J. Van der Pauw, Philips Res. Repts. 13, 1 (1958) 8. D.W. Koon & C.J. Knickerbocker, Rev. Sci. Instrum. 63, 207 (1992). 9. D.W. Koon & C.J. Knickerbocker, Rev. Sci. Instrum. 64, 510 (1993). 10. D.W. Koon & C.J. Knickerbocker, Rev. Sci. Instrum. 67, 4282 (1996). 11. R. Horiuchi, K. Okano, N. Rupesinghe, M. Chhowalla & G.A.K. Amaratunga,

Phys. Stat. Sol. A 193, 457 (2002). 12. H.P. Klug, L.E. Alexander, X-ray diffraction procedures for polycrystalline and

amorphous materials, chapter 9, Wiley, New York (1954). 13. P. Scherrer, Göttinger Nachrichten 2, 98 (1918). 14. A.H.J.M. Rijnders, The initial growth of complex oxides: study and

manipulation, PhD thesis, University of Twente, The Netherlands (2001). 15. A.K. van der Vegt, in Polymeren, van keten tot kunststof, 4th edition, Delft

University Press, The Netherlands (1999). 16. C.P.G. Schrauwen, R.A. Tacken, T. van Oudheusden & R. Renders, In

Adhesion of copper metallization films on ABS: effects of sputtering parameters and the influence of water, a model study, contribution to: 13.NDVaK, Beschichtigung und Modifizierung von Kunststofoberflächen, Dresden (2005).

17. M.R. Wertheimer, L. Martinu & E.M. Liston, in Plasma Sources for Polymer Surface Treatment, part E3.0. of Handbook of Thin Film Process Technology, edited by D.A. Glocker and S.I. Shah, Institute of Physics Publishing, Bristol, UK and Philadelphia, USA (1995).

18. L.J. Gerenser, in Surface Chemistry of Plasma-Treated Polymers, part E3.1. of Handbook of Thin Film Process Technology, edited by D.A. Glocker & S.I. Shah, Institute of Physics Publishing, Bristol, UK and Philadelphia, USA (1995).


19. E. Arenholz, J. Heitz, M. Wagner, D. Bäuerle, H. Hibst & A. Hagemeyer, Appl.

Surf. Sci. 69, 16 (1993). 20. W. Wong, K. Chan, K.W. Yeung & K.S. Lau, J. Mater. Process. Technol.

6020, 1 (2002). 21. Y. Han, D. Kim, J.-S. Cho, Y.-W. Beag & S.-K. Koh, Thin Solid Films 496, 58

(2006). 22. Y. Leterrier, Prog. Mater. Sci. 48, 1 (2003). 23. S. Petit, P. Laurens, J. Amouroux & F. Arefi-Khonsari, Appl. Surf. Sci. 168,

300 (2000). 24. S. Tamulevičius, Vacuum 51, 127 (1998). 25. X. Xu, J. Appl. Phys. 77, 6715 (1995). 26. F.J. Boerio, G.D. Davies, J.E. de Vries, C.E. Miller, K.L. Mittal, R.L. Opila &

H.K. Yasuda, Crit. Rev. Surf. Chem. 3, 81 (1993). 27. P. Spaepen, Acta Mater. 48, 31 (2000). 28. Y. Letterier, L. Médico, F. Demarco, J.-A.E. Månson, U. Betz, M.F. Escolà, M.

Kharrazi Olsson & F. Atamny, Thin Solid Films 460, 156 (2003). 29. J. Herrero & C. Guillén, Vacuum 67, 611 (2002). 30. J.H. Shin, S.H. Shin, J.I. Park & H.H. Kim, J. Appl. Phys. 89, 5199 (2001). 31. R.G. Gordon, MRS Bull. 25, 52 (2000).

Chapter 3 50

Chapter 4

Indium (Tin) Oxide

Abstract

Deposition of ITO on polymer substrates is restricted to room temperature during the growth process. By careful control of the pulsed-laser deposition parameters, films with high optical transmittance (>85%) and low resistivity (4.1 x10-4 Ωcm) are fabricated at room temperature on polyethylene terephthalate substrates. The films ablated from Sn-doped targets are more resistive compared to samples of pure In2O3. Due to increased scattering, the charge carrier mobility in Sn-doped films is lower compared to the undoped samples. A relation between the structural properties and the amount of Sn-doping is observed. The electrical properties of films with different compositions are influenced by a different size and formation of grains during growth.

Chapter 4 52

4.1 Introduction Since the first synthesis of SnO2 doped In2O3 (Indium Tin Oxide or ITO)

decades ago [1], the electrical and optical properties of this material are still superior to other TCOs. At present, ITO is commonly used as transparent electrode in (organic) light emitting diodes (OLEDs), solar cells and display technology. Especially in large area applications, this material is very suitable. When deposited by physical vapor deposition methods, the material can be applied very uniformly to glass substrates, a critical parameter for large area LCDs. ITO can be rapidly and uniformly etched. Moreover, the films are very stable, thus having a long life time. These features make ITO the unique candidate for display applications. Besides the developments in flexible displays, in many other optoelectronic applications there is a strong need to replace the commonly used glass substrates for cheap, light-weighted and non-rigid polymers. To allow deposition of highly conducting ITO on polymers, precise understanding of the materials fundamental aspects and optimization of the fabrication process are required.

The electrical properties of ITO films have been shown to depend mainly on the oxygen pressure and substrate temperature during thin film growth [2,3,4]. These parameters strongly influence the electron density in the deposited thin films. Donor generation in ITO films is namely governed by two mechanisms; the creation of double charged oxygen vacancies, and the contribution of a single electron for each Sn4+-cation substituted on an In3+-site. This intentionally doping of In2O3 with SnO2 is governed by a self-compensating reaction where electro-neutrality is preserved by the formation of neutral complexes [5,20]. Using the Kröger-Vink notation this reaction can be described as:

2 InInx + 2 SnO2 → (2 SnIn

• O’’)x + In2O3 (4.1)

In reducing environment the oxygen interstitials are removed, leaving one electron for each substituted Sn-ion:

(2 SnIn• O’’)x → 2 SnIn

• + 2 e' + ½ O2(g) (4.2)

In conditions that are highly reducing, oxygen anions are removed leaving doubly charged vacancies:

OOx → VO

•• + 2 e' + ½ O2(g) (4.3)

Whereas the last is mainly dependant on the oxygen partial pressure during deposition, the reactions described in equation 4.1 and 4.2 are dependant on SnO2 doping concentration as well as deposition temperature, since it is a thermally activated mechanism. Polymer substrates however cannot sustain these high temperatures normally required (300-500 ºC) to grow ITO films with excellent properties, i.e. low resistivity and high optical transmittance. A precise control of the

Indium (Tin) Oxide 53

process parameters during growth is therefore required in order to obtain high quality films at low substrate temperatures, allowing deposition on polymer substrates. The for PLD unique possibility of individual fine-tuning of deposition parameters and growth kinetics are beneficial for this research.

A tremendous amount of research has been conducted on the intrinsic properties and the thin film preparation by various deposition techniques of this material. In general sputtering techniques [6-789101112] are used to grow ITO films with excellent properties. But also e-beam evaporation [13-14151617], sol-gel [18,19], spray pyrolysis [20,21], chemical vapor deposition [22,23] and reactive thermal evaporation [24,25] are used to deposit this TCO. Only during the last five years, researchers began to exploit the possibilities of the deposition of ITO on polymer substrates. All the techniques mentioned above have been used to deposit ITO on PC [26,27], PET [28,29], PMMA [30,31] or other polymer substrates [32,33] with properties varying from moderate to fairly good. PLD as a deposition technique has also been used frequently to grow ITO films [2,34,35,36], however the use of polymer substrates is limited [37,38].

In this chapter the deposition of tin doped indium oxide films on polymer substrates by PLD will be discussed. Results on the electrical and optical performance of these room temperature grown films are presented section 4.3, whereas section 4.4 deals with the structural aspects. In section 4.5 the role of Sn- doping in In2O3 thin films on the microstructure affecting the electrical properties will be considered [39]. 4.2 PLD growth of indium tin oxide on polymers

In the experiments, commercially available indium oxide and indium tin oxide deposition targets were usedI. Besides pure In2O3, the stoichiometry of these targets 2 %wt, 5 %wt and 10 %wt SnO2 doped In2O3, referred to as ITO 2%, ITO 5% and ITO 10% respectively. As mentioned earlier, the electrical and optical properties of ITO are to a large extent dependant on oxygen deposition pressure and temperature. The influence of other parameters as the fluency on the target, laser frequency and distance from target to substrate were investigated in order to optimize these deposition conditions. The fluency mainly determines the target morphology evolution under laser irradiation of multiple pulses. Multi-component targets can suffer from non-stoichiometric ablation largely caused by the thermal properties of the elements [40]. Above 1.2 J/cm2, smooth target morphology is observed after several laser pulses from SEM imaging (figure 4.1), indicating stoichiometric and droplet-free evaporation.

I Targets supplied by Umicore Specialty Materials (USA). Target purity is >99.99%, and

the target density >95%.

Chapter 4 54

Other than the deposition rate, the influence of laser frequency and target-substrate distance was found to be insignificant on the thin film growth. Therefore the frequency and distance was fixed at 10 Hz and 50 mm respectively. Using these parameters and a fluency of 3 J/cm2, the deposition rate was about 1 Ǻ per pulse. The film thickness in the different experiments ranges from 125-500 nm.

As described in the previous section, the electrical resistivity of ITO thin films is strongly improved as the deposition temperature increases. Sn-activation is enhanced, and electrons are less scattered by structural disorder. The resistivity of crystalline films on glass substrates is normally around 2 x10-4 Ωcm. Only for epitaxial grown films lower values are reported, for example on YSZ a resistivity as low as 7.7 x10-5 Ωcm could be obtained [41]. A somewhat increased temperature of the polymeric substrate is therefore expected to be beneficial for the electric properties.

The crystallization temperature of ITO is around 150 ºC [14,35]. This is in accordance with the data in figure 4.2, which shows the diffraction patterns of ITO 5% films grown at different temperatures on glass substrates. For films grown at temperatures above 150 °C diffraction peaks are visible. Although there is no epitaxial relation with the glass substrate, the (004) reflection is dominant indicating a preferred (001) growth orientation. Below this temperature almost no film related intensity is present. Films grown at room temperature are completely x-ray amorphous. Thus, the deposition temperature should be close to the crystallization temperature of ITO, but should not significantly exceed the glass transition temperature of PET (Tg=75 °C) to avoid deformation of the polymer.

×2.7000 5µm 11kV ×50 500µm 17kV

Figure 4.1: SEM image of the ITO 5% target surface morphology after 200 laser pulses at a fluency of3 J/cm2 magnified 50x (a) and 2700x (b). The oval shape corresponds to the projection of a 59 mm2 mask which was placed in the beam path.

(a) (b)


Depositions at temperature above and below Tg are performed up to 130 °C. All these films showed mechanical failure, cracking or rippling after cool down as shown in figure 4.3. The morphology of the cracked films was checked by AFM and profilometer measurements. The results showed accumulation of material (buckling) at the break lines, indicating the presence of compressive strain upon cooling. This suggests that the films are cracked because of a large difference in the thermal expansion coefficient between film and substrate (thermal stress). The resulting compressive strain is released upon cracking as schematically drawn in the figure. Sometimes the large difference in shrink rate between film and substrate leads to delaminating of the thin film (figure 4.3b). Mechanical failure of the thin film due to thermal stresses is a commonly known problem in the coating of plastics above room temperature [42].

In order to obtain crack-free ITO films on polymers, the thin film growth is restricted to room temperature. Unless the resulting lower crystalline nature of ITO, other deposition parameters have to be tuned such that growth of high performance coatings remains possible. The existence of an optimum oxygen pressure for ITO growth at room temperature, where a high transmittance and low resistivity can be

Figure 4.2: X-ray spectra of ITO 5% target and films grown at increasing deposition temperatures. The diffraction pattern of the bulk powder is added for comparison. The scan speed is 0.5º p/m, except for the film grown at T=RT (1º p/m.). The broad halo peak around 18º originates from the glass substrate.

10 20 30 40 50 60 70

T=100 ºC.

T=300 ºC.

T=200 ºC.

T=RT.

Target(004)(222)

c/s

2θ (°)

Chapter 4 56

obtained, has been shown earlier by other authors [34,43,44]. Therefore this pressure regime for films with different composition has been closely examined. For each composition a series of 125-500 nm thick films is deposited on the polymer substrates at room temperature. Next to the composition, the oxygen deposition pressure close to the optimum pressure in steps of 0.002 mbar. In this pressure regime (P=0.003-0.033 mbar) no noticeable thickness variations have been observed. 4.3 Performance of ITO thin films on PET 4.3.1 Results

The electrical and optical data are plotted in figure 4.3a and b, respectively. The resistivity is measured using the four point van der Pauw geometry. The data shows an optimum around 0.015 mbar for all compositions. In this regime the resistivity is in the order of 5 x10-4 Ωcm. However, in contrast to high temperature deposition [2], lowest resistivity is observed for In2O3 films without additional SnO2 doping. The In2O3 film grown at optimum conditions has a resistivity of 4.1 x10-4 Ωcm. This value is comparable to that of films grown at high temperature or annealed samples [45].

The optical transmittance is measured between 200 and 1100 nm using a bare PET substrate as a reference. From these measurements, the average transmission in the visible part of the spectrum (400-700 nm) was calculated. In the same narrow pressure window around 0.015 mbar, where the lowest resistivity is

50 µm

Figure 4.3: Cracked ITO thin film surface on polymer substrate after cooldown from the deposition temperature of 80 ºC. AFM micrograph of a 150x150 µm film surface area, the z-scale is 2 µm (a). Section analysis and schematically representation of mechanical failure upon thermal stress (b) and optical microscope image showing partial delaminating film (c).

(a) (b) (c)


found, an optimum in the optical data (figure 4.4b) is observed. At this optimum, the average transmission is close to 90% for all film compositions.

Experiments showed that the data in figure 4.4 can be reproduced if the total deposition pressure is replaced by a partial oxygen pressure. An oxygen argon gas mixture with a total pressure of 0.050 mbar was used, and the gas flow ratio determines the oxygen pressure. Although at 0.050 mbar the mean free path of ablated particles is still larger than the target-substrate distance, the reproducibility proves that the resistivity curve is merely related to the oxygen incorporation and not associated to the kinetics during the ablation process. 4.3.2 Electrical properties

In order to understand the resistivity data, Hall measurements are performed to obtain the carrier density, and calculate the electron mobility. The results are plotted in figure 4.5 for two compositions; pure In2O3 and ITO 10%II. As the oxygen pressure decreases, more oxygen vacancies are incorporated in the film, increasing the carrier density for both compositions. However, for ITO 10% these values are somewhat lower than for pure In2O3. In contrast to high temperature deposition, the addition of SnO2 does not contribute to an increased carrier density, but it seems to deteriorate the film properties instead. The majority

II Data of ITO 2% and 5% was found to be in between the values of In2O3 and ITO 10%

for both carrier concentration and mobility. However, these points are not presented in this graph for clarity.

0.005 0.01 0.02 0.03 0.043

5

10

20

ρ

(x10

-4 Ω

cm)

pO2 (mbar)0.005 0.01 0.02 0.03 0.04

50

60

70

80

90

Av.

Tra

nsm

. (%

)

pO2 (mbar)

(a) (b)

Figure 4.4: Resistivity versus oxygen deposition pressure of In2O3 (), ITO 5% () and ITO 10% () films deposited at PET at T=RT (a). Average optical transmission (400-700 nm) of these films in the same pressure regime (b).

Chapter 4 58

of Sn-atoms are not thermally activated, i.e. substituting In3+, and are believed to form neutral complexes with the oxygen anions [20]. The charge carrier formation in ITO films grown at room temperature can therefore be mainly ascribed to oxygen vacancies, created in the low oxygen deposition pressure environment. It is therefore difficult for non-vacuum techniques to obtain high conductivity in these films. Note that when grown at high temperatures, the carrier density of SnO2 doped In2O3 can be as high as 1021 cm-3.

The mobility of the charge carriers is dependant on the scattering centers; in crystalline ITO films these are mainly ionized impurities [46]. The number of ionized impurities (oxygen vacancies and Sn-ions) is directly related to the carrier density. As oxygen deposition pressure increases, the amount of charge carriers decreases due to a smaller number of oxygen vacancies. As can be seen in figure 4.5, the total mobility is therefore higher since the number of ionized impurities decreases. However, independent of the difference in charge carrier concentration, the mobility for ITO films is considerably lower than for In2O3 for each deposition pressure. Apparently, in our amorphous/polycrystalline films, other scattering mechanisms play a role and are dependant on SnO2 doping. As mentioned earlier, Sn-atoms can form neutral complexes and precipitates, acting as neutral scattering centers. The amount of neutral defects will be substantially higher in our room temperature grown films, and will no longer play a marginal role as stated by Hamberg et al. [47]. Moreover, the low crystallinity also induces scattering at grain boundaries, further lowering the mobility. A relation between the microstructure and the electrical properties can therefore be expected, and will be discussed in the next paragraph.

Figure 4.5: Carrier density (circles) and mobility (squares) versus oxygen deposition pressure of In2O3

(closed symbols) and ITO 10% (open symbols) films deposited on PET at T=RT.

0.005 0.01 0.02 0.03 0.040.2

1.0

5.0

00

10

20

30

40

50

60

70

n (x

1020

cm

-3)

pO2 (mbar)

µ (c

m2 V -1

s-1)


4.3.3 Optical properties Transmission curves of ITO 5% deposited at different oxygen pressures are

illustrated in figure 4.6a, showing typical features of the ITO optical data. From the fringes caused by interference in the measured transmission or reflectance spectrum, the thickness can be determined for known refractive index (or vice versa). Although the refractive index is slightly dependant on carrier concentration and wavelength (section 2.3.2), an average value is used for these calculations. The obtained thickness from optical data is in good agreement with profilometer thickness measurements on the same samples.

From the graph it is evident that the total transmission of films deposited at low pressure (0.009 mbar) is decreased. The increasing oxygen deficiency causes a darkening of these films. This explains the steep drop in average transmission below 0.013 mbar in figure 4.4b for the 125 nm films. For higher pressures, the difference in transmission is less pronounced but drops gradually. These films are yellowish, and average transmittance is decreased. A photograph of the samples deposited at increasing oxygen pressure is shown in figure 4.6b. The grayscales show the trend of the films turning from opaque to clear transparent and finally grayish (corresponding to yellow).

300 400 500 600 700 800 900 10000

20

40

60

80

100

350 4000

20

40

60

pO2= 0.009mbar pO2= 0.013mbar pO2= 0.019mbar pO2= 0.025mbar

% T

rans

mis

sion

Wavelength (nm)

<0.001

pO2 (mbar)

>0.035

Figure 4.6: Transmission curves of ITO 5% films deposited at different oxygen pressures. The inset shows a magnification of the difference in transmission around the absorption edge (a). Photograph showing the coloring of samples as a result of the change in oxygen depositionpressure (b).

(b)

Chapter 4 60

The small decrease in transmission at high pressures is a result of the narrowing of the bandgap, i.e. shifting to lower energies. The inset of figure 4.6a shows a magnification of the transmittance around 400 nm of films deposited at 0.013 mbar, 0.019 mbar and 0.025 mbar. A small shift in the absorption edge to longer wavelengths can be observed. More light at shorter wavelengths in the visible spectrum is absorbed, coloring the film yellow. This phenomenon is known to be caused by the Burstein-Moss (BM) shift [48] as discussed in chapter 2, and is dependant on the carrier concentration. For indium (tin) oxide films having high carrier concentrations, the narrowing bandgap behavior can always be observed [47, and the references therein]. The bandgap is calculated from the extrapolation of α2 versus the photon energy hν, as described in section 2.3.1. Figure 4.7a shows the extrapolation of the absorptivity for three 125 nm In2O3 films at different oxygen pressures. The x-axis intersection, corresponding to the optical gap, decreases at higher oxygen deposition pressure as a result of the difference in carrier concentration. The obtained bandgap values are expected to increase with the carrier concentration by an n2/3 dependence according to BM-theory (section 2.3.2). In figure 4.7 the measured optical bandgap values for In2O3 and ITO 10% corresponding to the 125 nm films in figure 4.3b are plotted against n2/3.

The linear nature of both curves shows that the bandgap widening is indeed proportional to n2/3. The data of the bandgap values are within 0.03 eV of the fitted straight lines. However, the data point of ITO 10% at the highest carrier

3.0 3.2 3.4 3.6 3.80

1

2

3

4

5

6

7

a2 (x10

13 c

m-1)

hv (eV)0 1 2 3 4 5 6

3.3

3.4

3.5

3.6

Opt

cal b

andg

ap (e

V)

n2/3 (x1013cm-2)

Figure 4.7: Tauc plot of In2O3 films deposited at 0.015 mbar (), 0.021 mbar () and 0.027 mbar (∆) (a). Variation of the optical energy gap against carrier concentration n2/3 of In2O3 () and ITO 10% (∆) films deposited on PET at T=RT (b).

(a) (b)


concentration is omitted in the fit. Thin films of ITO 10% with these carrier concentrations can be found in films deposited around 0.007 mbar. As can be seen from figure 4.4b the transmittance for these films has dropped below 50%. Such low transmission prevents an accurate calculation of the optical gap from the Tauc plot. Furthermore, a bandgap much higher than 3.5 eV can not be measured on PET since above this photon energy the optical data will be affected by the absorption of the substrate. Assuming that only the BM-shift changes the bandgap, one can calculate the fundamental bandgap Eg0 and the reduced effective mass m*cv. The former can be obtained from the y-axis intercept, whereas the latter is determined by the slope of the fitted curve. The fundamental gaps of In2O3 and ITO 10% are 3.30 eV and 3.37 eV respectively. These values are somewhat lower than the intrinsic bandgap values for crystalline films [49], but resemble those of amorphous In2O3 and ITO films earlier reported by Bellingham et al. [50].

The reduced masses calculated from the slope of the linear fit are both for In2O3 and ITO around 1.03·me, which is in agreement with the values ranging from 0.49·me–1.2·me reported earlier [49]. However, if a nominal value of m*cv=0.55·me and m*c=0.35·me [51] is substituted in equation 2.21, a negative value for m*v is obtained. This would imply that the curvature of the valence band is negative, which is impossible. Hamberg et al. [47] and Roth et al. [52] suggested that in addition to bandgap widening, effects of bandgap narrowing removes this anomaly. The contribution is indicated as the term ћΣ in formula 2.19 (section 2.3.2). The downward shift of the conduction band is a consequence of electron-electron and electron-impurity scattering. The bandgap narrowing because of heavy doping is quite common in silicon [53], and seems to play an important role in ITO as well. This effect is also proportional to n2/3, and when taken in consideration it reduces the slope of the graph [49]. The lower value of m*cv results in a positive m*v, hence resulting in conduction band that is curved upwards as illustrated in figure 2.1 (section 2.2.1). 4.4 Microstructure 4.4.1 TEM analysis

XRD data shows no intensity that could be related to the crystallinity of room temperature grown films. The film is either completely amorphous, or the grains are too small to provide clear x-ray reflections. TEM measurements can obtain information on the nanocrystallinity of the samples. Using this technique, the difference in grain sizes of samples with different compositions is investigated. For this experiment films of 25 nm were deposited on Cu-grids at similar conditions as the thicker films on PET. Figure 4.8 shows transmission electron micrographs and diffraction patterns of undoped and doped In2O3. A clear difference in structural

Chapter 4 62

properties can be seen. In2O3 samples clearly contain crystallites, whereas the ITO 10% film has a nanocrystalline, nearly amorphous phase. The visible grains are not necessarily completely crystalline, since the existence of amorphous grains in ITO has been observed earlier [46]. However, the diffraction patterns of these samples confirm a decreased crystallinity of ITO compared to In2O3. The ITO 10% sample has smaller and less intense diffraction spots, combined with the hazy diffraction rings. Some of the crystallites of pure In2O3 are as large as several tenths of nanometers. Diffraction rings show sharp diffraction spots. The TEM micrographs of ITO 2% and ITO 5% doping are intermediate states complementing the trend of decreasing crystallinity as the doping level is increased. 4.4.2 Surface morphology

It is possible that the nature of the substrate (the Cu-grids) has influenced the grain growth. Therefore other experiments should confirm the relation between doping and microstructure on other substrates. For this purpose, AFM measurements are performed to check the surface morphology. AFM surface scans

Figure 4.8: TEM images of In2O3 (a), ITO 2% (b), ITO 5% (c) and ITO 10% (d) deposited at roomtemperature. The insets show the corresponding electron diffraction patterns of the selectedarea.


of In2O3 and ITO 10% films on polymer substrates are presented in figure 4.9a and b. In figure 4.9c, the surface roughness in relation to the doping content is shown.

The results confirm the found relation between doping and crystallinity Not only the average particle size is larger at the surface of the undoped film, also the surface roughness is increased significantly, indicating a different growth mode for these compositions. Figure 4.9c illustrates the decreasing root mean square (RMS) surface roughness for higher doping concentrations. Note that the average surface roughness of the polymer substrate is well below 1 nm.

From the TEM and AFM data it is observed that the grain size in room temperature grown films is decreasing on the addition of SnO2. This would imply that during PLD growth, Sn-atoms prevent the grains from growing. In order to check this hypothesis, a study on the initial growth during deposition is required. Only little work on in-situ monitoring of ITO films has been reported, and are all concerned with resistivity measurements during growth. These measurements were mainly used to determine the crystallization temperature [14] and growth mode analysis [54,55]. From this last work a transition from a 3D island growth mode at low temperatures to 2D growth at higher temperature was observed. This transition from the Volmer-Weber mechanism to the Frank van-der-Merwe mechanism [56] occurs at the crystallization temperature (~150 °C).

(a) (b)

(c)

Figure 4.9: AFM micrographs showing the surface of In2O3 (a) and ITO 10% (b) films. The relation between the RMS surface roughnesses in nanometers versus SnO2-doping concentration in weight percentage (c). The error margins are within 10 %.

Chapter 4 64

Instead of the thermodynamic growth behavior of doped In2O3, the diffusivity of the material is more important in order to explain the initial grain growth behavior. In this study the film growth is therefore studied by another approach, using in-situ AFM combined with PLDIII. This recently developed technique allows real-time and real-space imaging of the substrate surface during initial film growth [57]. In these experiments an atomically flat STO surface is used as ultra smooth substrate [58]. Prior to the experiments, the properties of thin films deposited in the AFM-PLD were confirmed to reproduce the films grown in the conventional PLD setup. The resolution of in-situ AFM is not sufficient to detect variations in the length scales of diffusion at room temperature. Since the diffusion of atomic species on a substrate surface is a thermally activated mechanism [59], experiments at higher temperatures increase the diffusion length. Therefore the experiments are

III At that time, the in-situ AFM-PLD was under development and the measurement was

part of the experimental test phase.

(a) (b)

Figure 4.10: In-situ AFM micrographs showing the STO surface after deposition of 6 pulses In2O3 (a) and ITO 10% (b).The lower profiles show the section analysis of the image corresponding to thewhite lines.


performed at a substrate temperature of 300 ºC, keeping the other parameters equal to those described in section 4.2. After each deposition pulse, AFM images are recorded to follow the initial thin film growth.

Results of in-situ AFM experiments are shown in figure 4.10. These images are the observed STO surface after 6 laser pulses. A difference can be distinguished between the initial growth behaviors of the films having different compositions. In2O3 forms relatively large grains separated about 1 µm from each other on average. The surface in between the grains is clean, even showing the substrate steps on closer look. Although some island coalescence can be observed for ITO, on average the grains are smaller. The inter-grain distance is less compared with In2O3, and material is found in between the islands. This difference indicates that Sn-atoms are less mobile than In-atoms and therefore act as nucleation center. A higher density of nucleation sites results in the formation of more, hence smaller grains. If the growth is continued after 6 pulses, all grains coalesce and form a closed film.

4.5 Discussion

From the experiments we learned that highly conductive and optical transparent films can be grown in a narrow deposition regime at room temperature by PLD. The films are more resistive as the Sn-doping content in In2O3 increased. Furthermore, from the TEM and AFM measurements one can conclude that room temperature grown films need to be considered as nanocrystalline instead of completely amorphous. TEM as well as (in-situ) AFM measurements confirm a decrease of the grain size as the doping content increases.

The decreasing grain size coincides with the increasing resistivity at higher doping concentrations. There is a continuous debate on the influence of grain boundary scattering in polycrystalline ITO thin films. In general it is presumed that grain boundary scattering does not contribute to a decrease in resistivity of the polycrystalline since the mean free path is in the order of only a few nanometers [60]. If in equation 2.13 the obtained values for the mobility and carrier concentration are used, the maximum mean free path is indeed about 5 nm. For ITO 10%, the average grain size drops below 10 nm. These small grains bring forth a significant amount of grain boundaries. This causes the effective grain size, defined as the distance between the depletion zones forming the grain boundary potential barriers, to be even smaller. In this case the mean free path of the free carriers is in the same range as the effective grain size. This results in the contribution of the grain boundary scattering mechanism to the total mobility as discussed in section 2.2. Hence, the intrinsic conductivity limit is lower for the nanocrystalline ITO films are lowered.

Chapter 4 66

In figure 4.11 the intrinsic mobility limit is modeled similar to the results presented in figure 2.4. However, the grain size is now set to only 5 nm, which consequently leads to a reduced µgb compared to the 50 nm used earlier. The mobility limit due to ionized impurities is modeled for two cases. The first extremity is for the films containing only oxygen vacancies (Z=2, Ni=n/2), which are double ionized and give the largest scattering contribution. This case is similar to In2O3 and approaches RT-grown ITO where no Sn is activated. The other upper limit of µiis is modeled for films where single ionized donors provide scattering (Z=1, Ni=n). This is approached by high-temperature deposited ITO films, where the electrons from Sn-activation are the dominant free carriers.

The data points in the graph correspond to In2O3 and ITO 10% samples deposited at deposition pressures close to the optimum of 0.015 mbar, i.e. where the resistivity is lowest. The data of these samples shows that in accordance with figure 4.5 the mobility is low, whereas the carrier concentration is high for samples deposited at pO2<0.015 mbar. The arrows in the graph indicate the direction of increasing oxygen deposition pressure of the samples during the fabrication

1019 1020 1021 1022100

101

102

103

Z=2

µiis + gb

µiis

µgb

mob

ility

µ (c

m2 V

-1s-1

)

carrier concentration n (cm-3)

Z=1

Figure 4.11: The upper limit of mobility µiis + gb (solid line) influenced by the effect of scattering on ionizedimpurities (dashed curves) and grain boundaries (dotted curve).The total mobility wascalculated with µiis for double ionized donors (Z=2; Ni=n/2). The upper limit of µiis for single ionized donors is also shown (Z=1, Ni=n). The average grain size for calculating µgb was taken as 5 nm. Data points for RT-grown In2O3 () and ITO 10% () are shown, the gray arrow indicate the increase of deposition pressure of the samples. One data point of ITO 5% deposited on YSZ at T=600 °C (×) is also shown.


process. The large amount of free carriers due to oxygen vacancies cannot compensate for the reduced mobility. As a result of a too high oxygen deficiency the conduction network is disturbed. As deposition pressure increases the conduction network is restored, and the mobility is only limited by the ionized impurity scattering if the grain size is larger than the mean free path of the electrons. This is the case for In2O3 and the data points approach the intrinsic mobility limit for scattering on oxygen vacancies. The mobility of ITO 10% samples however does not approach this limit. In contrast to In2O3 the resistivity of these samples is also affected by grain boundary scattering. Therefore the data points stay close to the total mobility limit affected by both (µis + gb).

It is clear that different mechanisms influence the conductivity of crystalline and the room-temperature deposited ITO films. For comparison a film of ITO 5% was deposited on YSZ at T=600 °C. From XRD measurements it is observed that this film is epitaxially grown on the substrate. The large grain size makes the influence of grain boundaries insignificant. In contrast to ITO grown at room temperature, most of the added Sn is activated and produces the majority of carriers. The mobility value of this sample is therefore closer to the higher mobility limit (Z=1). 4.6 Concluding remarks

Up to now, ITO is the most important TCO for commercial applications. In most of these applications the need for cheap, light-weight and flexible substrates as polymers is increasing. The implementation of polymer substrates requires an essential optimization in the fabrication process in order to deposit high performance TCOs at reduced substrate temperatures.

The electrical and optical properties of ITO are mainly determined by the process parameters from which the oxygen pressure and substrate temperature are most important. The former mainly determines the carrier density of this TCO, whereas the latter is related to the structure, i.e. crystallinity of the thin film. Deposition of ITO on polymers at elevated temperatures by PLD resulted in cracked films due to the different thermal expansion coefficient between film and substrate. The fabrication process is therefore restricted to room temperature.

Although the crystallization temperature of ITO is only 150 °C, RT-grown films are nanocrystalline. But precise tuning of the oxygen deposition pressure resulted in films with electrical and optical properties comparable with high temperature deposited ITO. The electrical conductivity of these room temperature grown films is mainly determined by oxygen deficiency instead of activated donors. Addition of SnO2 in In2O3 does not result in increased carrier concentrations. The majority of Sn-atoms are not thermally activated, i.e. substituting In3+ for Sn4+.

Chapter 4 68

Furthermore, the Sn-doped films are more resistive compared to pure In2O3 because of a structural difference. The presence of Sn during growth leads to the formation of smaller grains. The increased scattering and trapping of charge carriers deteriorates the electrical properties. In ITO the intrinsic mobility limit is not only limited by ionized impurity scattering, but the grain boundary scattering has a significant influence as well. For pure In2O3 films lowest resistivity of 4.1 x10-4 Ωcm is observed, and the optical transmittance well exceeds 85%.


4.7 References 1. R. Groth & E. Kauer, Philips Technical Review, 26, 105, (1965) 2. F. Hanus, A. Jadin, & L.D. Laude, Appl. Surf. Sci. 96-98, 807 (1996). 3. H. Kim, J.S. Horwitz, A. Pique, C.M. Gilmore & D.B. Chrisey, Appl. Phys. A.

69, S447 (1999). 4. S.H. Kim, N.-M. Park, T.Y. Kim & G.Y. Sung, Thin Solid Films 475, 262

(2005). 5. J.-H. Hwang, D.D. Edwards, D.M. Kammler & T.O. Mason, Solid State Ionics

129, 135 (2000). 6. J.C.C. Fan & J.B. Goodenough, J. Appl. Phys. 48, 3524 (1977). 7. S. Ray, R. Banerjee, N. Basu, A.K. Batabyal & A.K. Barua, J. Appl. Phys. 54,

3497 (1983). 8. L.-j. Meng, A. Macarico, R. Martins & L.-J. Meng, Vacuum 46, 673 (1995). 9. Y.S. Jung, J.Y. Seo, D. W. Lee & D.Y. Jeon, Thin Solid Films 445, 63 (2003). 10. K. Zhang, F. Zhu, C.H.A. Huan & A.T.S. Wee, Thin Solid Films 376, 225

(2000). 11. L.R. Cruz, C. Legnani, I.G. Matoso, C.L. Ferreira & H.R. Moutinho, Mater.

Res. Bull. 39, 993 (2004). 12. Y.S. Jung & S.S. Lee, J. Cryst. Growth 259, 343 (2003). 13 . Y. Shigesato & D.C. Paine, Appl. Phys. Lett. 62, 1268 (1993). 14. D.C. Paine, T. Whitson, D. Janiac, R. Beresford & C.O. Yang, J. Appl. Phys.

85, 8445 (1999). 15. S.A. Agnihotry, K.K. Saini, T.K. Saxena, K.C. Nagpal & S. Chandra, J. Phys.

D: Appl. Phys. 18, 2087 (1985). 16. M. Yamaguchi, A.I. -Ektessabi, H. Nomura & N. Yasui, Thin Solid Films 447,

115 (2004). 17. I. Hamberg, C.G. Granqvist, K.-F. Berggren, B.E. Semelius & L. Engström,

Phys. Rev. B: Condens. Matter, 30, 3240, (1984). 18. M.J. Alam & D.C. Cameron, Thin Solid Films, 76, 420, (2002). 19. P.K. Biswas, A. De, N.C. Pramanik, P.K. Chakraborty, K. Ortner, V. Hock &

S. Korder, Mater. Lett. 57, 2326 (2003). 20. G. Frank & H. Köstlin, Appl. Phys. A. 27, 197 (1982). 21. T. Kawashima, H. Matsui & N. Tanabe, Thin Solid Films 445, 241 (2003). 22. O.O. Akinwunmi, M.A. Eleruja, J.O. Olowolafe, G.A. Adegboyega & E.O.B.

Ajayi, Opt. Mater. 13, 225 (1999). 23. K. Maki, N. Komiya & A. Suzuki, Thin Solid Films 445, 224 (2003) 24. P. Thilakan & J. Kumar, Vacuum 48, 463 (1997). 25. S. Noguchi & H. Sakata, J. Phys. D: AppI. Phys. 14, 1523 (1981). 26. D. Kim, Appl. Surf. Sci. 218, 71 (2003).

Chapter 4 70

27. Y. Hoshi, H.-o. Kato & K. Funatsu, Thin Solid Films 445, 245 (2003) 28. T. Ohishi, J. Non-Cryst. Solids 332, 87 (2003). 29. Y.-S. Kim, Y.-C. Park, S,G, Ansari, J.-Y. Lee, B.-S. Lee & H.-S. Shin, Surf.

Coat. Technol. 173, 299 (2003). 30. S.-S. Lin & J.-L. Huang, Ceram. Int. 29, 771 (2003). 31. N. Al-Dahoudi, H. Bisht, C. Gobbert, T. Krajewski & M.A. Aegerter, Thin Solid

Films 392, 299 (2001). 32. C. Nunes de Carvalho, G. Lavareda, E. Fortunato & A. Amaral, Thin Solid

Films 427, 215 (2003). 33. C. Guillén & J. Herrero, Thin Solid Films 431-432, 403 (2003). 34. J.P. Zheng & H.S. Kwok, Appl. Phys. Lett. 63, 1 (1993). 35 . X.W. Sun, H.C. Huang & H.S. Kwok, Appl. Phys. Lett. 68, 2663 (1996). 36. J.H. Kim, K.A. Jeon, G.H. Kim & S.Y. Lee, Appl. Surf. Sci. 252, 4834 (2006). 37. H. Kim, J.S. Horwitz, G.P. Kushto, Z.H. Kafafi & D.B. Chrisey, Appl. Phys.

Lett. 79, 284 (2001). 38. T.K. Yong, T.Y. Tou & B.S. Teo, Appl. Surf. Sci. 248, 388 (2005). 39. J.M. Dekkers, G. Rijnders & D.H.A. Blank, Appl. Phys. Lett. 88 151908

(2006). 40. L.M. Doeswijk, in Pulsed laser deposition of oxides on silicon: Exploring their

passivating qualities, Chap.2, PhD thesis, University of Twente, The Netherlands (2002)

41. H. Ohta, M. Orita, M. Hirano, H. Tanji, H. Kawazoe & H. Hosono, Appl. Phys. Lett. 76, 2740 (2000).

42. Y. Leterrier, Prog. Mater. Sci. 48, 1 (2003). 43. M.A. Morales-Paliza, M.B. Huang & L.C. Feldman, Thin Solid Films 429, 220

(2003). 44. M.A. Morales-Paliza, R.F. Haglund & L.C. Feldman, Appl. Phys. Lett. 80,

3757 (2002). 45. H.L. Hartnagel, A.L. Dawar, A.K. Jain & C. Jagadish, in Semiconducting

Transparent Thin Films, section 2.5, Institute of Physics Publishing, Bristol and Philadelphia (1995).

46. H.-C. Lee & O. Ok Park, Vacuum 75, 275 (2004). 47. I. Hamberg & C.G. Cranqvist, J. Appl. Phys. 60, R123 (1986). 48. E. Burstein, Phys. Rev. Lett. 93, 632 (1954). 49. L. Gupta, A. Mansingh & P.K. Srivastava, Thin Solid Films 76, 33 (1989). 50. J.R. Bellingham, W.A. Philips & C.J. Adkins, J. Phys.: Condens. Matter 2,

6207 (1990). 51. H. Köstlin, R. Jost & W. Lems, Phys. Status Solidi A 29, 87 (1975). 52. A.P. Roth, J.B. Webb & D.F. Williams, Phys. Rev. B 25, 7836 (1982).


53. L. Vina, C. Umbach, M. Candona, A. Compann & A. Axman, Solid State

Commun. 48, 457 (1983). 54. J.P. Zheng & H.S. Kwok, Thin Solid Film, 232, 99 (1993). 55. J.P. Zheng & H.S. Kwok, Appl. Surf. Sci. 106, 51 (1996) 56. D.B. Chrisey & G.K. Hubler, In Pulsed Laser Deposition of Thin Films, John

Wiley & Sons Inc., New York (1994) 57. J.J. Broekmaat, G. Rijnders & D.H.A. Blank, unpublished. 58. G. Koster, B.L. Kropman, G. Rijnders, D.H.A. Blank, H. Rogalla, Appl. Phys.

Lett. 73 (1998), 2920. 59. A.H.J.M. Rijnders, in The initial growth of complex oxides: study and

manipulation, PhD thesis, University of Twente, The Netherlands (2001). 60. T.J. Coutts, J.D. Perkins, D.S. Ginley & T.O. Mason, Conference Paper,

NREL/CP-520-26640, (1999).

Chapter 4 72

Chapter 5

Doped Zinc Oxides

Abstract

Transparent conducting oxides composed of (doped) zinc oxide and zinc containing compounds are of major interest since they are regarded as replacement material for ITO. The conductivity in these materials is completely governed by oxygen vacancies and is therefore strongly dependant on the oxygen partial pressure during thin film growth. Also in this material small amount of doping (Al2O3 ~1 %wt.) enhances the electrical and optical properties of the room temperature grown films. In order to grow smooth films by pulsed laser deposition on polymer substrates, the process parameters should be adjusted to prevent excessive substrate heating. In order to benefit from the advantages of both In2O3 and ZnO, the In2O3-ZnO compound system on PET substrates is investigated. The electrical resistivity can be tuned by the ZnO content as it increases one order of magnitude on going from pure In2O3 to pure ZnO. The optical transmission is high and constant over the whole composition range.

Chapter 5 74

5.1 Introduction

Although ITO films still are the best performing transparent conductors, the use of this material in future applications is believed to decrease. As mentioned in chapter 1, the price of indium is subject to major price fluctuations and has become an uncertain factor for manufacturers of flat panel displays. Since the display market is still expanding, researchers are intensively searching for alternatives. Zinc oxide (ZnO) has been nominated as the most important candidate. ZnO has some advantages above the nowadays predominant ITO. The raw mineral resources are cheap and abundant, ZnO is nontoxic and has good stability in hydrogen plasma, which is of significance for applications related to amorphous silicon solar cells [1,2,3]. However, the use of undoped ZnO has two major disadvantages. ZnO is more chemical unstable than ITO [4], and the electrical resistivity is higher by at least one order of magnitude [5]. On the other hand, ZnO is easily doped n-type [6], and numerous dopants are used in order to improve the electrical properties and chemical stability.

Highly transparent and conductive ZnO have been prepared by doping with a group III (B, Al, Ga or In), group IV (Si, Ge, Ti, Zr or Hf) [7] and even rare-earth elements using various deposition techniques [8]. Fluorine is used for anionic doping [9]. Most of the different dopants contribute to the enhancement of specific properties of undoped ZnO. Among these, Al is mostly used as a dopant because of the strong positive effects on the electrical conductivity and chemical stability.

Instead of impurity doped binary compounds (ZnO:Al or In2O3:Sn), multicomponent oxides have been attracting much attention as new TCO materials. Five closely grouped d10 cations are known to be constituents of transparent conductors, namely Zn2+, Cd2+, In3+, Ga3+ and Sn4+. Except for Ga, they form TCOs as single cation oxides. From combinations of these binary oxides different TCO compound systems can be formed [10]. Phases consisting of ternary, quarternary and quintenary systems are reported or under investigation by many groups [11]. The basic idea of this research is to combine the different material advantages as etchability, resistivity, transmittance and stability, to one new TCO material tailored to specific applications.

The benefits of ZnO can be combined with In2O3, which have been proven to be low resistive and highly conducting when deposited at room temperature by PLD. Therefore the ZnO-In2O3 system seems a good candidate as TCO on polymer substrates. In this chapter, deposition of undoped and impurity doped zinc oxide as well as the binary-binary compound ZnO-In2O3 system will be considered. The influence of the fabrication method to grow crack-free films on polymer substrates and the composition dependence will be discussed.

Doped Zinc Oxides 75

5.2 ZnO and AZO deposition on PET Al-doped ZnO (AZO) thin films deposited on rigid glass substrates have

been extensively studied in recent years [12-13141516], including pulsed laser deposited samples [17-18192021]. Recently the PLD growth of AZO films with a resistivity as low as 1.3 x10-4 Ωcm has been reported [22]. On the other hand, only few papers have been published for AZO films deposited on polymer substrates [23,24,25], These films were all deposited by sputter techniques. It is well known that during sputter deposition the substrate temperature can increase significantly due to the continuous bombardment of energetic particles. Although this temperature rise has never been an issue in the numerous reports on sputter deposition on polymer substrates, it seemed problematic in the few cases of AZO coating on plastic. Pei et al. reported a temperature increase of about 80 ºC during AZO deposition on PET [25]. In this case a 40 nm Al2O3 buffer layer was required to prevent damage to the substrate, whereas Fortunato et al. [24] used special polymer substrates resistant to moderate temperatures. Only Yang et al. [23] deposited AZO films directly onto heat sensitive polymeric substrates. The films on PET showed moderate resistivity of 2.7 x10-3 Ωcm, and a transmission of 84%. Since the substrate temperature rise during PLD growth is expected to be insignificant, this method can be advantageous for growing AZO thin films on polymers. Furthermore the stoichiometric material transfer from target to substrate is beneficial since AZO films are optimized at low Al2O3 concentrations (up to 3 %wt.).

The free charge carriers in undoped ZnO are generated by oxygen vacancies. Similar to Sn4+ substituting In3+ in In2O3, the replacement of Al3+ for Zn2+ in ZnO adds one extra electron to the valence band. This substitution reaction is also dependant on the process temperature. Hence, the process parameters of pulsed laser deposited zinc oxides at room temperature on polymeric substrates will strongly influence the TCO performance. Deposition targets were prepared from ZnO and Al2O3 powdersI. Besides pure ZnO, powders were mixed to form stoichiometric compounds containing 1, 2 and 3 weight percentage of Al2O3, referred to as AZO 1%, AZO 2% and AZO 3% respectively. The mixtures were ball-milled for 24 hours in ethanol, and after drying they were pressed to 16 mm pellets; uniaxially pre-pressed at 200 bar, and subsequently isostatically pressed at 4000 bar. The pellets were sintered at a temperature of 1100 °C in atmospheric conditions. All targets possessed a density of about 90% of the theoretical density. The Al2O3 content was checked by x-ray fluorescence (XRF) measurements, and was found to be consistent within 0.1 %wt.

I ZnO (99.999% pure) powder is purchased from Alfa Aesar GmbH, whereas Al2O3

(AKP50 99.999%) is purchased from Sumitomo Chemical. Co. Ltd.

Chapter 5 76

From the targets, thin films with typical thickness of 150 nm were deposited on PET substrates. Initially, the deposition parameters were chosen similar to those used for ITO growth. However, using these settings cracks appeared in the film surface after ZnO deposition on PET (figure 1a). Profilometer scans perpendicular to crack lines revealed a height increase. This indicates that the film cracks upon compressive strain, similar to ITO deposited on heated polymer substrates (chapter 4). Also if quartz is used, the films crack and delaminate from the substrate surface (figure 1b). An increased temperature of film and substrate surface is suspected to be responsible for building in this strain, which is later released by cracking.

As discussed in section 3.4.2, the deposition process can increase the substrate temperature during growth. In contrast to the continuous sputtering process, the perturbation of the substrate temperature in PLD is only significant within the top several microns of the substrate for a period of a few microseconds after each pulse [26]. The total temperature rise at the substrate surface is approximately inversely proportional to the square root of the thermal diffusivity κ of

Figure 5.1: Optical microscope photographs of ZnO thin films deposited using F=3 J/cm2 and d=50 mm on PET (a) and quartz (b). ZnO thin films deposited on PET using F=1.5 J/cm2 and d=50 mm (c), and d=70 nm (d).Photographs show the edge of the cracked film created by deposition mask and the undamaged substrate surface (upper part).

(a)

(c) (d)

(b)


the substrate materialII. By placing a thermocouple in thermal contact with the substrate surface during deposition of ZnO this temperature rise was monitored. From the results shown in figure 5.2 it can be observed that both PET and quartz show a rise in temperature of at least 10 ºC. This is at least a factor of 10 higher than for silicon, and is expected from values for the thermal diffusivity.

Despite the increased substrate surface temperature ITO films grow flat, whereas ZnO films crack using similar deposition parameters. One explanation can be the difference in thermal expansion coefficient of ZnO (4.75 x10-6 K-1) which is about 50% smaller than In2O3 [27]. Another important parameter is the molecular weight of the incident particles, since lighter species condense on the substrate in a shorter period of time [26]. Furthermore, plume characteristics such as energy distribution play an important role as well.

The energy of the species impinging the film and substrate is determined by the deposition parameters used in the experiment. Other than the oxygen pressure, which should be around 0.01 mbar for optimum results, the parameters are changed in order to prevent film cracking. Decreasing the fluency on the target resulted in less crack formation in the ZnO film. A too low fluency however can result in inhomogeneous ablation from multi-component targets as AZO. Although in literature values of 1.2 J/cm2 are reported, SEM analysis of the AZO targets

II Thermal diffusivity is defined as the ratio between the thermal conductivity (λc) and

volumetric heat capacity (ρd•cp). Values for polymers are listed in table 3.1

1T κ

−∆ ∝

Figure 5.2: Substrate surface temperature increase during ZnO deposition on different substrates. The values for the thermal diffusivity κ of the substrates are also listed. Used depositionparameters; F=3 J/cm2, f=10 Hz, P=0.035 mbar and d=50 mm. The laser is stopped after 360seconds, indicated by the vertical dotted line.

0 60 120 180 240 300 360 420 480 540 600

0

2

4

6

8

10

12

14

16

Tem

pera

ture

(°C

)

t (sec.)

PET, κ= 8 x10-3 cm2/s Quartz, κ= 1.4 x10-2 cm2/s Silicon, κ= 0.8 cm2/s

Chapter 5 78

showed irregularities on the ablated surface. This can be an indication of non-stoichiometric ablation of zinc and aluminum. Therefore the fluency could be decreased to a minimum value of 1.5 J/cm2 in the settings used. A small improvement is also observed if the laser frequency is reduced, related to the heating of the substrate. In addition fewer cracks are formed if the target – substrate distance is increased. This can be understood by the pressure – distance relation [28,29]. A larger distance at fixed deposition pressure corresponds to a larger deposition pressure at fixed distance.

Despite the alterations of the deposition parameters, no crack-free ZnO/AZO films could be obtained (figure 5.1d) at an oxygen deposition pressure of 0.01 mbar. After optimization of the parameters, it is experimentally determined that crack-free films can be grown from approximately 0.03 mbar and higher. For this reason the total pressure during deposition is chosen as 0.05 mbar with an oxygen argon mixture background gas. The oxygen partial pressure is thus determined by the individual oxygen and argon flow into the PLD chamber. 5.3 Results 5.3.1 Undoped ZnO thin films

Similar to the results for ITO films, an optimum oxygen pressure can be observed for ZnO grown films figure 5.3a). At 0.010 mbar the resistivity is at it lowest value of 2.8 x10-3 Ωcm. The resistivity increases at high oxygen pressure due to reduction of oxygen vacancies. At low pressures the higher carrier concentration reduces the mobility, hence increasing resistivity. Furthermore the optical transmittance is decreased as the material is more oxygen deficient (figure 5.4b). Films deposited in pure argon environment are black. The main difference to ITO films is the lower carrier concentration, resulting in a resistivity for ZnO that is a factor of four higher. At the optimum pressure the mobility is 35.8 cm2V-1s-1 and the carrier concentration is 6.3 x1019 cm-3, compared to 37.3 cm2V-1s-1 and 2.4 x1020 cm-3 for ITO 10%. respectively. For the same reason the optical gap of ZnO is not changing with oxygen partial pressure, i.e. carrier concentration, in contrast to ITO films. The highest carrier concentration is still too low to give a significant contribution from the BM-shift to the fundamental gap. By taking an effective mass of 0.24me [30], one can calculate the contribution from 6.3 x1019 cm-3 charge carriers, which is only 0.15 eV. The measured optical gap for this film is 3.35 eV. Taken the fundamental gap as 3.2 eV [5] this implies that the measured optical gap of 3.35 eV is in agreement with this prediction.

The optical gaps of all films in this pressure regime are not significantly affected by the BM-shift, and can also be found around 3.3 eV. In contrast to ITO, the average transmission of the films in relation to the deposition pressure is


therefore not changed. The bandgap is not varying, causing different absorption close to or in the visible part of the spectrum. For this reason the optical transmission for films deposited at pressures higher than 0.010 mbar is not decreasing. In figure 5.3b it can be seen that it continues to rise at higher pressures as oxygen deficiencies are diminished.

5.3.2 Al2O3 doped ZnO

The experiments on AZO films on polymers showed that the dependency on oxygen pressure is similar. For AZO 1%, AZO 2% and AZO 3% thin films an optimum pressure is found at 0.010 mbar. This minimum resistance for AZO films grown at 0.010 mbar with different Al2O3 doping is presented in figure 5.4a. The lowest resistivity in these room temperature grown samples is obtained in AZO 1%, and is around 2 x10-3 Ωcm. Earlier results also showed an optimal oxygen pressure for PLD grown AZO films on glass at room temperature [20,21]. At an optimum of 0.007 mbar Singh et al. reported a resistivity of 7.4 x10-4 Ωcm for 2 %wt. Al2O3 doped ZnO. Increasing the Al2O3 content further however, again results in higher resistance. This behavior is determined by the carrier concentration and mobility. Figure 5.4b and c shows the dependency of these quantities on the Al2O3 content. If small amounts of Al2O3 are added to ZnO, the carrier concentration increases rapidly because Al3+ substitutes Zn2+, adding electrons to the conduction band. If the Al2O3 content increases further, the carrier concentration is not

1E-3 0.01 0.0570

75

80

85

90

95

100

Tran

smis

sion

(%)

pO2 (mbar)1E-3 0.01 0.05

1E-3

0.01

0.1

1

ρ

(Ωcm

)

pO2 (mbar)

Figure 5.3: Resistivity versus the partial oxygen deposition pressure of ZnO films deposited at PET atT=RT (a). Corresponding average optical transmission (400-700 nm) of these films in this pressure regime (b).

(a) (b)

Chapter 5 80

changing significantly. Similar to SnO2 in In2O3, the aluminum is not thermally activated and maintains electrically inactive. Besides the carrier concentration, the electron mobility has a large influence on the resulting resistivity. Although the highest carrier concentration is found at 2 %wt. Al2O3, the mobility has dropped significantly, increasing the film resistance. At first the mobility drops rapidly due to increased ionized impurity scattering for higher carrier concentrations. However, comparing the AZO 1% and AZO 3% samples, the mobility drops further whereas the carrier concentration is constant. Since neutral impurities have negligible influence (chapter 4), the mobility is likely to be affected by grain boundary scattering due to lower crystallinity. Figure 5.5a shows that there is indeed a relation between the doping content and the crystallinity of the grown films. Not only the intensity of the (002) diffraction peak drops as Al2O3 content is increased, but also the Full Width Half Maximum (FWHM) increases. The data can be fitted to Gaussian functions and an estimate of the grain size is determined using the Scherrer relationIII. The grain size decreases from ~30 nm to ~20 nm for ZnO and AZO 3% respectively. This indicates that small amounts of Al2O3 obstruct the grain grown in the polycrystalline AZO films. Analogue to SnO2 in In2O3, a larger amount of grain boundaries are formed affecting the mobility.

III The XRD data is analyzed and edited by the X’Pert Highscore plus software package

(PANalytical B.V., The Netherlands). The structural broadening used in the grain size calculation was 0.03º.

0 1 2 305

10152025303540

µ (c

m2 V

-1s-1

)

Al2O3 (%wt.)0 1 2 3

0.5

1.0

1.5

2.0

2.5

3.0

n x

1020

(cm

-3)

Al2O3 (%wt.)0 1 2 3

0.002

0.003

0.004

0.005

0.006

0.007

ρ (Ω

cm)

Al2O3 (%wt.)

Figure 5.4: Resistivity (a), carrier concentration (b) and mobility (c) versus the Al2O3 content in ZnO films deposited at PET at T=RT. The deposition is done at the optimal oxygen partial pressure of0.010 mbar for all films.

(a) (c)(b)


The optical properties of AZO thin films are summarized in figure 5.5b. The graph shows the measured optical bandgap and the average optical transmission. Comparing the obtained bandgap values with the carrier density in figure 5.4b it can be observed that they are strongly related. This indicates that the optical gap for these AZO films is affected by the BM-shift. The inset of graph 5.5b shows the energy gap versus n2/3 of the different AZO, and the straight line is drawn according to the BM-relation (formula 2.20). As discussed in chapter 4, the intrinsic bandgap is obtained from the y-axis intersection. A value of Eg0=3.21 eV is found, and corresponds to the fundamental bandgap used earlier. The difference bandgap for ZnO and AZO films is reflected on the total average transmittance. All the Al2O3 doped films are more transparent (>93%) than ZnO films (<92%) because of the increased bandgap. 5.3.3 Review and discussion Doping of TCOs can result in enhanced electrical properties. Particularly at higher deposition temperature, the lowest resistivity in AZO films is generally found around 2 to 3 %wt. Al2O3, whereas ITO is least resistive around 5 %wt. SnO2. The data of doped zinc oxide on polymeric substrates shows that there are many similarities with doped indium oxide as discussed in chapter 4. Because of the low process temperature used for deposition on polymers, the donors are not thermally

0 1 2 33.30

3.35

3.40

3.45

3.50

3.55

3.60

91

92

93

94

95

Opt

ical

gap

(eV)

Al2O3 content (%wt.)

Av.

Tra

nsm

. (%

)

34.0 34.5 35.0 35.5 36.0

ZnO

AZO 1 %wt.

AZO 2 %wt.

AZO 3 %wt.

c/

s

2θ (º)

Figure 5.5: Structural and optical properties dependant on the Al2O3 content in ZnO films deposited at PET at T=RT. The ω-2θ scans of the (002) diffraction peak and the calculated grain size is displayed in (a). The inset shows the average grain size obtained from the FWHM datacalculated by the Scherrer relation. Part (b) shows the measured optical gap () and the average optical transmission between 400-700 nm. (). The inset graph shows the bandgap-n2/3 relation of the AZO films according to the BM-theory.

(a) (b)

0 1 2 3

200

250

300

d (Å

)%wt.

0 5

3.2

3.6

Eg (e

V)

n2/3 (x1013 cm-2)

Chapter 5 82

activated to provide the necessary electrons for conduction. In contrast to Sn doping in ITO however, small amounts of Al2O3 doping in AZO does contribute to increased carrier concentrations. But this only holds for a doping concentration of 1 %wt. Al2O3, whereas the minimum doping content in ITO was 2 %wt. So a minimum resistivity in ITO 1% thin films might be possible. Since doping of room temperature deposited TCOs is not significantly effective, the electron concentration is mainly determined by the amount of incorporated oxygen vacancies. As a consequence, the carrier concentration is strongly dependant on the oxygen process pressure. The pressure for incorporation of sufficient vacancies should be controlled precisely and is relatively low, requiring vacuum techniques for manufacturing. Although the grains are very small, the structural nature of PLD deposited doped TCO systems on polymer substrates should be interpreted as polycrystalline rather than amorphous. Adding dopants results in a further decrease of the grain size. This negatively influences the electrical properties of the films. The number of grain boundaries is increased, and the effective grain dimensions are in the order of the mean free path of electrical carriers. This source of scattering lowers the intrinsic upper limit of mobility, hence increases the minimum obtainable resistance. In general the electrical end optical properties of the doped TCOs on polymers are acceptable. The properties of the undoped oxides are comparable, and even better than samples grown by other techniques. Because of the lack of thermal diffusion of oxygen during growth, most vacancies are incorporated in room temperature grown films. The actual substrate temperature in the PLD process is closer to room temperature than for instance during sputtering. 5.4 Indium Zinc Oxide

As mentioned in section 5.1, combining the advantages of both indium oxide and zinc oxide may result in a TCO that possesses mutual properties. Indium Zinc Oxide (IZO) films reported in literature are mainly deposited by sputtering techniques [31-32333435]. Many authors investigate multiple compositions to search for the ideal film stoichiometry in different applications. Sometimes combinatorial sputtering techniques are used to deposit varying compositions on one large substrate [36,37]. These setups allow the exploration of ZnO-In2O3 phase in one single run. In some cases PLD is used to grow IZO thin films [38,39], and a wide range of compositions is also used in these experiments [40,41]. Except for a few reports on sputtered IZO films on PC [42] and PET [43,44], all depositions are mainly performed on glass.


ZnO and In2O3IV powders were mixed in the atomic ratio defined as

x=[Zn]/Zn+In]. The value ranges from x=0 (pure In2O3) to x=1 (pure ZnO). The same balll-milling and target pressing procedure as for AZO targets was used. The sintering temperature was 1000 °C in atmospheric conditions. The obtained pellets have densities around 80% of the theoretical values. XRD data revealed that the targets are a ZnO-In2O3 powder mixture in all cases. The atomic ratio in the resulting targets was checked and confirmed by x-ray fluorescence (XRF) measurements. Using the deposition conditions as used for AZO thin films, IZO thin films are grown on PET at room temperature. The deposition rate decreased almost linearly from 1.0 Å/pulse to 0.75 Å/pulse on going from a composition of x=0 to x=1, respectively. The deposition times were adjusted such that all films have a thickness of 135 nm.

The stoichiometry of the grown films can be described as ZnxInyOx+1.5y with 0≤x≤1, and y=1-x. The electrical properties of IZO thin films with intermediate composition are shown in figure 5.6. Similar to pure In2O3 (x=0) and ZnO (x=1), a minimum in the resistivity is observed at optimal deposition pressure (figure 5.6a). This minimum also shifts from around 0.015 mbar for pure In2O3 down to 0.010 mbar for pure ZnO films. Furthermore, the resistivity increases for higher ZnO content, shifting the curves upwards. As for the single binary oxides, the shape of the curves can be explained by the relation between the carrier density and mobility (figure 5.6b and c).

IV ZnO (99.999% pure) and In2O3 (99.999% pure) powder is purchased from Alfa Aesar

GmbH and Aldrich respectively.

0.004 0.01 0.02 0.03 0.041E19

1E20

5E20

x= 0.75

x= 0.50

x= 0.35

x= 0.10

n (c

m-3)

pO2 (mbar)

0.004 0.01 0.02 0.03 0.04

1E-3

0.01

x= 0.75

x= 0.50

x= 0.35

x= 0.10

ρ (Ω

cm)

pO2 (mbar)

0.004 0.01 0.02 0.03 0.04

10

20

30

40

50

60

x= 0.75

x= 0.50

x= 0.35

x= 0.10

µ (c

m2 V

-1s-1

)

pO2 (mbar)

(a) (c) (b)

Figure 5.6: Electrical properties ZnxInyOx+1.5y thin films with x=0.10, 0.35, 0.50 and 0.75 on PET substrates. The resistivity (a), carrier concentration (b) and mobility (c) are plotted versus theoxygen partial deposition pressure.

Chapter 5 84

The XRD spectrum of the ZnxInyOx+1.5y films in figure 5.7 show the part of the spectrum where the prominent diffraction peaks of In2O3, ZnO and IZO phases are expected. The data shows that only pure ZnO is polycrystalline, i.e. contains crystallite grains large enough that give rise to diffracted x-ray intensity. It has been shown earlier that the regime 0.2<x<0.7 results in amorphous films even grown at elevated deposition temperatures [37], Clear intensities of Zn5In2O8, Zn3In2O6 and Zn2In2O5 phases are only observed if higher process temperatures (~300 ºC) are used [45].

Figure 5.8 shows the optical and electrical properties of IZO thin films deposited at the optimum deposition pressure versus the ZnO content. The resistivity increases almost logarithmic with the ZnO content. This is in contrast to the work of Minami et al. [31] who reported a minimum resistivity at x=0.3 for RT-sputter deposited IZO films. The carrier concentration peaked at x=0.3, which was attributed to the existence of an IZO phase that enhanced conductivity. However, the IZO films with lower ZnO content (x<0.3) has low carrier concentrations (n< 2 x1020 cm-3) increasing the resistivity in this regime, whereas the carrier concentration remains high if grown by PLD at T=RT (figure 5.8b). Sputtered In2O3 films are in general more resistive than PLD grown films. The higher effective temperature in the sputter process lowers its resistivity (chapter 4). If no additional heating is applied, the process temperature in PLD is closer to room temperature compared to sputter deposition. For this reason an optimum doping content around

Figure 5.7: The ω-2θ scans of ZnxInyOx+1.5y thin films with different ZnO content x.The arrows indicatethe positions where the mentioned reflection is expected.

30 31 32 33 34 35 36 37 38

In2O3

(004)ZnO(002)

Zn5In2O8

(0021)Zn3In2O6

(0015)Zn2In2O5

(008)In2O3

(222)

x= 1.00

x= 0.75

x= 0.50

x= 0.35

x= 0.22

x= 0.10

x= 0.00

c/s

2θ (°)


x=0.3-0.5 is always found in sputtered IZO films, with or without additional heating [37,43,45]. Sputtered IZO films with low ZnO content (x≈0.1) have similar or worse conductivity regardless of the deposition temperature [46,47,48] compared to the values in figure 5.8a. The only reported composition dependant PLD grown IZO films also show a minimum doping content of x=0.3 since they are deposited at temperatures above 110 ºC [40,41].

The evolution of resistivity versus ZnO content can be roughly divided into two regimes. From x=0-0.3 both the carrier density and mobility decrease slowly, whereas from x=0.3-1 the carrier density decreases sharply and the mobility remains constant. The carrier concentration is directly related to the

0.0 0.2 0.4 0.6 0.8 1.03.2

3.3

3.4

3.5

93

94

95

96

97

Eg (e

V)

%at. Zn

Tr. 4

50 n

m (%

)

0.0 0.2 0.4 0.6 0.8 1.05E19

1E20

2E20

3E20

4E20

5E20

0

10

20

30

40

50

60

carr

ier d

ensi

ty (c

m-3)

%at. Zn

mob

ility

(cm

2 V-1s-1

)

0.0 0.2 0.4 0.6 0.8 1.03E-4

5E-4

1E-3

0.002

0.003

80

85

90

95

100

Res

istiv

ity (Ω

cm)

%at. Zn

av.

Tr (

%)

Figure 5.8: Electrical and optical properties of ZnxInyOx+1.5y thin films with different ZnO content x. The resistivity () and average optical transmittance () between 400-700 nm (a), Carrier concentration () and mobility () (b) and the bandgap energy () and transmittance () at450 nm are shown (c).

(a)

(b) (c)

Chapter 5 86

accommodation of oxygen vacancies. In the first ‘indium-rich’ regime more vacancies can be created compared to the ‘zinc-rich’ material.

The average optical transmission is constant and as high as 92% over the complete range of compositions. However, the bandgap is varying as the composition changes. In the ‘indium-rich’ regime, the bandgap is higher than the ‘zinc-rich’ regime since not only the fundamental bandgap of In2O3 is somewhat higher; the optical gap is also enhanced by the higher carrier concentration. Similar to these observations, a minimum bandgap value of about 3.25 eV was reported earlier for ZnxInyOx+1.5y films with x=0.7 [40].

The different onset of absorption should influence the total average transmission as in ITO or INO films. However, the average transmittance in figure 5.8a remains constant despite the composition related varying bandgap. The most plausible explanation for this is that the absolute transmission in the visible region is higher for the ‘zinc-rich’ material. On average the higher total transmission compensates the decreasing bandgap on the addition of zinc. To illustrate this, the bandgap and total transmission at 450 nm of the films dependant on compositions is shown in figure 5.8c, From the figure it is immediately clear that the as the onset of absorption moves towards the visible wavelengths, the total transmission in the remainder of the optical window between 400 and 700 nm is increasing. The relation between the ZnO content in PLD deposited IZO on polymers and the electrical performance allows certain tune-ability of the resulting TCO thin film. As the optical transmission remains constant over the complete range, the electrical conductivity can be changed. Applications where a resistivity in the order of 10-3 Ωcm is sufficient can benefit from reduced indium content. Furthermore the etching rate is directly related to the ZnO content in IZO films [33]. Thus IZO thin films on polymers can therefore be easily tailored to specific applications, and are useful candidates for plastic-TCO systems. 5.5 Concluding remarks ZnO is seen as the most prominent candidate to replace ITO as a TCO in large scale applications, including those on polymers. Knowledge of the properties and growth conditions of this material on heat resistive substrates is therefore required. The PLD growth of ZnO on PET substrates at room temperature is therefore investigated. It was found that after deposition, ZnO films on PET substrates were cracked because of the strained film. Atomic ZnO species in the ablation process are highly energetic, and due to subsequent impingement on the substrate the temperature rises. Upon cooling the ZnO film on the substrate cracks because of different expansion coefficients. Smooth ZnO films on polymer


substrates can only be grown if the PLD deposition parameters, being the deposition pressure as the most important, are adjusted.

Doping ZnO with Al can enhance the electrical conductivity and chemical stability. In contrast to the addition of Sn in In2O3 films by PLD grown at room temperature, a small amount of Al in ZnO does enhance the electrical properties. But a similar trend of decreasing grain size upon doping is observed. The crystallinity decrease on the addition of Al2O3, lowering the mobility and consequently the conductivity. Excellent optical properties of ZnO and AZO films are observed. In contrast to AZO thin films, the doping concentration of undoped ZnO is too low to enhance the bandgap.

Growing doped TCOs on polymer substrates is subject to precise control of the deposition parameters, especially the deposition pressure. While the substrate is at room temperature to prevent thermal stress, it is possible to accurately control and optimize the process conditions by PLD. Although the low process temperature reduces the effect of doping, good quality films can still be grown.

The ZnO-In2O3 (IZO) system combines the properties of both individual oxides. Although different IZO phases do exist in high temperature grown films, in RT-grown films on polymers by PLD only amorphous states are detected, except for pure ZnO. Although a varying bandgap is observed, the visible transmission is constant regardless of the ZnO content in IZO. The resistivity decreases logarithmic from ~10-4 to ~10-3 Ωcm as doping ranges from 0 to 100% ZnO content. This makes IZO a material that is easy to adapt in order to meet specific conditions for applications on polymers.

Chapter 5 88

5.6 References 1. Z.-C. Jin, L. Hamberg & C.G. Granqvist, J. Appl. Phys. 64, 5117 (1988). 2. S. Major, S. Kumar, M. Bhatnagar & K.L. Chopra, Appl. Phys. Lett. 49, 394

(1986). 3. M. Chen, Z. L. Pei, X. Wang, C. Sun & L.S. Wen, J. Vac. Sci. Technol. A 19,

963 (2001). 4. T. Minami, S. Suzuki, T. Miyata, Thin Solid Films 398-399, 53 (2001). 5. H.L. Hartnagel, A.L. Dawar, A.K. Jain & C. Jagadish, Semiconducting

Transparent Thin Films, Institute of Physics Publishing, Bristol and Philadelphia (1995).

6. A. Zunger, Appl. Phys. Lett. 83, 57 (2003). 7. H. Sato, T. Minami & S. Takata, J. Vac. Sci. Technol. A 11, 2975 (1993). 8. T. Minami, T. Yamamoto & T. Miyata, Thin Solid Films 366, 63 (2000). 9. J. Hu & R.G. Gordon, Sol. Cells 30, 437 (1991). 10. T.J. Coutts, D.L. Young & X. Li, MRS Bull. 25, 58 (2000). 11. A.J. Freeman, K.R. Poeppelmeier, T.O. Mason, R.P.H. Chang & T.J. Marks,

MRS Bull. 25, 45 (2000). 12. S.J. Jung, Y.H. Han, B.M. Koo, J.J. Lee & J.H. Joo, Thin Solid Films 475, 275

(2005). 13. J. Yoo, J. Lee, S. Kim, K. Yoon, I.J. Park, S.K. Dhungel, B. Karunagaran, D.

Mangalaraj & J. Yi, Thin Solid Films 480-481, 213 (2005). 14. B.-Y. Oh, M.-C. Jeong, W. Lee & J.-M. Myoung, J. Cryst. Growth 274, 453

(2005). 15. K. Ellmer, F. Kudella, R. Mientus, R. Schieck & S. Fiechter, Thin Solid Films

247, 15 (1994). 16. B. Szyszka, Thin Solid Films 351, 164 (1999). 17. Z.Y. Ning , S.H. Cheng, S.B. Ge, Y. Chao, Z.Q. Gang, Y.X. Zhang & Z.G. Liu,

Thin Solid Films 307, 50 (1997). 18. H. Tanaka, K. Ihara, T. Miyata, H. Sato & T. Minami, J. Vac. Sci. Technol. A

22, 1757 (2004). 19. H. Kim, U,A. Pique, J.S. Horwitz, H. Murata, Z.H. Kafafi, C.M. Gilmore & D.B.

Chrisey, Thin Solid Films 377-378, 798 (2000). 20. A.V. Singh, R.M. Mehra, N. Buthrath, A. Wakahara & A. Yoshida, J. Appl.

Phys. 90, 5661 (2001). 21. H. Kim, C.M. Gilmore, J.S. Horwitz, A. Pique, H. Murata, G.P. Kushto, R.

Schlaf, Z.H. Kafafi, & D. B. Chrisey, Appl. Phys. Lett. 76, 259 (2000). 22. S.-M. Park, T. Ikegami, K. Ebihara & P.-K. Shin, Appl. Surf. Sci. 253, 1522

(2006).


23. T.L. Yang, D.H. Zhang, J. Ma, H.L. Ma & Y. Chen, Thin Solid Films 326, 60

(1998). 24. E. Fortunato, P. Nunes, D. Costa, D. Brida, I. Ferreira & R. Martins, Vacuum

64, 233 (2002). 25. Z.L. Pei, X.B. Zhang, G.P. Zhang, J. Gong, C. Sun, R.F. Huang & L.S. Wen,

Thin Solid Films 497, 20 (2006). 26. X. Xu, J. Appl. Phys. 77, 6715 (1995). 27. K. Ellmer, J. Phys. D: Appl. Phys. 34, 3097 (2001). 28. H.S. Kim & H.S. Kwok, Appl. Phys. Lett. 61, 2234 (1992). 29. J. Gonzalo, R. Gómez San Román, J. Perrière, C.N. Afonso & R. Pérez

Casero, Appl. Phys. A 66, 487 (1998). 30. S.J. Pearton, D.P. Norton, K. Ip, Y.W. Heo & T. Steiner, Superlattices

Microstruct. 34, 3 (2003). 31. T. Minami, H. Sonohara, T. Kakumu & S. Takata, Jpn. J. Appl. Phys. 34, L971

(1995). 32. K. Zhang, F. Zhu, C.H.A. Huan, A.T.S. Wee & T. Osipowicz, Surf. Interface

Anal. 28, 271 (1999). 33. T. Minami, T. Kakumu & S. Takata, J. Vac. Sci. Technol. A 14, 1704 (1996). 34. Y.S. Song, J.K. Park, T.W. Kim & C.W. Chung, Thin Solid Films 467, 117

(2004). 35. H.-C. Pan, M.-H. Shiao, C.-Y. Su & C.-N. Hsaio, J. Vac. Sci. Technol. A 23,

1187 (2005). 36. M.P. Taylor, D.W. Readey, C.W. Teplin, M. van Hest, J.L. Alleman, M.S.

Dabney, L.M. Gedvilas, B.M. Keyes, B. To, P.A. Parilla, J.D. Perkins & D.S. Ginley, Macromol. Rapid Commun. 25, 344 (2004).

37. M.P. Taylor, D.W. Readey, C.W. Teplin, M. van Hest, J.L. Alleman, M.S. Dabney, L.M. Gedvilas, B.M. Keyes, B. To, J.D. Perkins & D.S. Ginley, Meas. Sci. Technol. 16, 90 (2005).

38. N. Naghavi, A. Rougier, C. Marcel, C. Guéry, J.B. Leriche & J.M. Tarascon, Thin Solid Films 360, 233 (2000).

39. K. Ramamoorthy, K. Kumar, R. Chandramohan, K. Sankaranarayanan, R. Saravanan, I.V. Kityk & P. Ramasamy, Opt. Commun. 262, 91 (2006).

40. N. Naghavi, C. Marcel, L. Dupont, A. Rougier, J.-B. Leriche & C. Guéry, J. Mater. Chem. 10, 2315 (2000).

41. M. Mikawa, T. Moriga, Y. Sakakibara, Y. Misaki, K.-I. Murai, I. Nakabayashi & K. Tominaga, Mater. Res. Bull. 40, 1052 (2005).

42. W.J. Lee, Y.-K. Fang, J.-J. Ho, C.-Y. Chen, L.-H. Chiou, S.J. Wang, F. Dai, T. Hsieh, R.-Y. Tsai, D. Huang & F.C. Ho, Solid-State Electron. 46, 477 (2002).

Chapter 5 90

43. T. Minami, T. Kakumu, Y. Takeda & S, Takata, Thin Solid Films 290-291, 1

(1996). 44. H.-M. Kim, S.-K. Jung, J.-S. Ahn, Y.-J. Kang & K.-C. Je, Jpn. J. Appl. Phys.

42, 223 (2003). 45. T. Moriga, T. Okamoto, K. Hiruta, A. Fujiwara & I. Nakabayashi, J. Solid State

Chem. 155, 312 (2000). 46. Y.S. Jung, J.Y. Seo, D.W. Lee & D.Y. Jeon, Thin Solid Films 445, 63 (2003). 47. Y.S. Song, J.K. Park, T.W. Kim & C.W. Chung, Thin Solid Films 467, 117

(2004). 48. N. Ito, Y. Sato, P.K. Song, A. Kaijio, K. Inoue & Y. Shigesato, Thin Solid Films

496, 99 (2006).

Chapter 6 P-type conducting spinel ZnM2O4

Abstract

The research on p-type TCOs is of major importance for future applications in transparent electronics. For applications, a material that is less dependent on the film structure is desirable. In certain materials transparency can be generated by the so-called ligand field splitting in the cationic d-bands in octahedral geometry. The electrical and optical properties of epitaxial and polycrystalline spinel ZnM2O4 are investigated, where M=Co, Rh and Ir. Photoemission spectroscopy revealed that the valence band maximum is composed of occupied t2g

6 states. The observed bandgap is increasing for higher quantum numbers, being as large as ~3 eV for ZnIr2O4, which is expected from theoretical predictions. P-type conductivity is confirmed in all compounds by the position of the Fermi level and Seebeck measurements. Grown in polycrystalline phase, films of these materials still display high conductivity, well above 2 Scm-1.

Chapter 6 92

6.1 Introduction At present, TCOs are used in numerous devices in our daily lives. These

applications are mostly found in display technology, thin-film photovoltaics and ‘smart windows’. In all of these cases however, the TCOs are used as passive materials, i.e. transparent electrodes. For this purpose mostly n-type TCOs are used, because of the excellent optical properties and electrical controllability. In contrast to passive electrodes, the use of TCOs as active components in electronics is still limited and mostly restricted to the field of research. However, it is believed that transparent electronic circuits will serve as the basis for new optoelectronic devices in the next decades [1,2]. Although transparent thin-film field effect transistors (FETs) composed of only n-type TCOs have been fabricated [3,4,5,6] forming the first transparent circuits [7], the market of transparent electronics has not yet been developed due to the lack of appropriate p-type TCOs [8]. The research on p-type TCOs is of major importance, since the p-n junction is an essential building block in all semiconducting electronics.

Recent developments in p-type transparent oxide semiconductors (TOSs) have led to new applications as bleu and near-UV light emitting diodes (LEDs) [9,10,11]. LEDs based on p-type TCOs are demonstrated by several authors [12,13]. An overview of the developments of devices based on crystalline transparent conducting oxide thin films is reported by Banerjee et al. [14]. The materials in this review either consist of cuprates, mixed- or binary oxides as p-ZnO.

The design of a transparent p-type conducting semiconductors is normally subjected to two issues [8]. Due to the ionicity of metallic oxides, the 2p levels of oxygen atoms are in general far lower in energy than the valence band orbitals of metallic atoms. Positive holes, introduced by substitutional doping, are localized on the oxygen atoms and cannot migrate within the valence band. This can be solved by introduction of covalency in the metal-oxygen bond to create an extended valence band. The other issue is that the cationic metal should have a closed d- shell to avoid coloration resulting from d-d transitions. These conditions are best fulfilled by copper, and it is therefore that cuprates as CuAlO2 [15] and SrCu2O2 [16] still belong to the best performing p-type TOSs. Relying on the O-Cu-O dumbbell structure in these crystals, O 2p- and Cu 3d-states are hybridized to form delocalized valence bands. Because of these copper-oxide chains however, the electrical properties are very dependant on the crystalline nature.

Another approach is to dope existing n-type TCOs to produce p-type materials. In the last decade, tremendous effort has been made to dope zinc oxide as a p-type semiconductor. ZnO has some advantages above cuprates as large exciton binding energy and abundant mineral resources [17,18]. In contrast to n-type doping however (chapter 5), p-type doping of ZnO is difficult, due to the low

P-type conducting spinel ZnM2O4 93

solubility of the dopant and high compensation effect upon doping [19,20]. Nitrogen as a dopant seemed the best candidate, resulting in the first p-type conducting ZnO films [21]. This ignited an extensive research on p-ZnO including arsenic and phosphor as a dopant [22-,23,24,25]. However, only moderate carrier density can be obtained (order 1017), which is still inferior to n-type TOSs. Moreover, the growth of p-ZnO is submitted to delicate film growth. A high temperature deposition, combined with a controlled growth mode is required to ensure the p-type properties of these materials.

The mentioned cuprates and zinc-oxides have in common that reasonable p-type conductivity can only be obtained in a crystalline structure. The high deposition temperature makes them therefore not suitable for deposition on polymer substrates. Although moderate temperature deposition of nanocrystalline CuAlO2 films have been demonstrated [26,27], and the processing temperature of SrCu2O2 can be as low as 200 ºC [28], also these materials are still not appropriate for deposition on heat resistive substrates. Materials that can be deposited at room temperature, but still having this p-type behavior, would have a great potential.

Mizoguchi et al. [29] reported on spinel ZnRh2O4 as a new class of p-type TOSs. The conduction mechanism of this material is different from other crystalline p-type TOSs. It has been shown that no long range order is required for hole conductivity in this material [30]. The bandgap in this material is determined by the so-called ligand field splitting of the Rh 4d-bands as a result of the geometrical surrounding of oxygen anions.

In this chapter, the properties of ZnM2O4 films (where M=Co, Rh and Ir) as TOSs are investigated. The spinel structures containing these three d6- transition metals are supposed to be p-type TOSs and the bandgap is determined by the magnitude of the ligand field splitting. From theory the bandgap is expected to increase on going from Co to Rh and to Ir. This mechanism to obtain transparency is different from that in earlier found p-type TCOs. In section 6.2 these mechanisms and the theory involved will be discussed. From this one can predict the optical properties of the three ZnM2O4 compounds. The next section will deal with the bulk synthesis and analysis of these compounds for PLD target manufacturing. Subsequently, the thin film deposition as well as structural and compositional analysis will be described. In section 6.3 and 6.4, the optical and electrical properties are investigated, respectively.

6.2 Band phenomena in spinel structures 6.2.1 Ligand field theory

For interpreting the chemistry of coordination compounds the established valence band theory [31] was replaced by the crystal field theory, proposed by the

Chapter 6 94

physicist Hans Bethe in 1929 [32]. Subsequent modifications were proposed by J.H. Van Vleck in 1935 to allow for some covalency in the interactions. These modifications are often referred to as The Ligand Field TheoryI [33]. Pure crystal field theory assumes that the interactions between the metal ion and the ligands are purely electrostatic (ionic). If a spherically symmetric field of negative charges (i.e. the ligand) is placed around the metal, the d- orbitals remain degenerate, but are raised in energy as a result of the electrostatic repulsion. In ligand field theory, the surrounding ligands are regarded as point charges. Rather than a spherical field, these discrete point charges are allowed to interact with the metal. Dependant on the geometrical surrounding, not all d-orbitals will interact to the same extent with the point charges. Some orbitals will be destabilized more than others, resulting in an energy splitting of the d- bands. The number of split levels, and the relative energy between these levels can be calculated for all possible geometrical surrounding of ligands [34].

On one hand the energy split is dependant on the number, geometry and the nature of the ligand. On the other hand the split is determined by the oxidation state and the nature of the metal. Jørgensen has proposed an empirical formula which allows to estimate the ligand field splitting; ∆=f•g, where f is a function of the ligand and g is a function of the metal [35]. The value for g is expressed in electron volts whereas f is a pure number relative to the hexa-aqua ion as a ligand. Based on data for a wide variety of complexes, Jørgensen listed the ligands and metals in order of increasing field strength in the so-called spectrochemical series, the values for f and g. 6.2.2 d-orbital splitting in octahedral structures

The largest energetic split is found in octahedral geometries. Spinel phases are of the form AB2O4, where A and B are divalent and tetravalent metals respectively. The crystal lattice contains mostly tetrahedral AO4 and octahedral BO6 coordinated structures. This site preference is dependant on charge and ionic radius [36]. Higher charge on the metal is better counterbalanced by six oxygen anions, whereas smaller ionic radius fits better in the tetrahedral site. The spinel lattice of AB2O4 containing the tetrahedral and octahedral lattice sites is schematically drawn in figure 6.1aII. The six oxygen point charges are located on the +x, -x, +y, -y, +z and -z axes respectively around the central B3+ ion as

I For a review on the evolution of bonding models see C. J. Ballhausen, J. Chem. Ed.

56 194-197, 215-218, 357-361 (1979). II Picture from: C. Renkenberger, in Neue oxidische Rhodiumverbindungen:

Syntheseversuche und strukturelle untersuchungen, p.41, PhD. Thesis, Heidelberg, Germany (2004).


illustrated in figure 6.1b. The d- orbitals which lie along the diagonals of the octahedron (i.e. dx2-y2, dz2) will be destabilized more that the other orbitals (i.e. dxy, dxz, dyz) and are raised in energy (figure 6.1c and d). The latter are referred to as the t2g set, whereas the first two belong to the higher energetic eg set. The magnitude of the energetic field splitting in between the two sub-bands is denoted as ∆o. These two sub-bands can be filled by electrons in two ways in case of d4- d7- ions. In case of a weak field situation; i.e. small energy split, all electrons are unpaired because of Hund’s rule and the coulumbic repulsion between electrons when forced in the same orbital. So one or two electrons will occupy the high energy eg set, the so called high spin-state. In the strong field situation, the energy

Figure 6.1: Schematic representation of the AB2O4 spinel lattice showing the octahedral (grey) andtetrahedral structures (black). For clarity, not all polyhedrons are highlighted (a) in which themetallic cation B is surrounded by oxygen ligands in octahedral geometry (b). Thegeometrical nature of the five d-orbitals (c) and the corresponding energy split ∆ due to theligand field are shown (d).

(b) (a)

(c) (d)

y

Z

y

Z

y

Z

y

Z

y

Z

dx2-y2 dz2

dxz dxy dyz

eg

t2g

∆0

y

Z

O B A

Chapter 6 96

split becomes larger than the pairing energy, and the electrons only occupy the low energetic t2g subset. This is known as the low spin-state. This pairing energy becomes smaller as one goes down a period. Therefore, as a rule, 4d- and 5d-transitionmetals are always low spin. 6.2.3 Bandgap engineering in ZnM2O4

As mentioned, Mizoguchi et al. [29] reported on spinel ZnRh2O4 as a new transparent p-type semiconductor. Due to the ligand field 4d- orbitals are split and the six available electrons of the Rh3+ ion completely occupy the t2g

6 subset, whereas the eg

0 levels remain empty. This low-spin state acts as a “quasi closed” shell configuration as sketched by the six electrons in figure 6.1d. The t2g

6 level behaves as a (fully occupied) valence band, whereas the eg

0 level is the (unoccupied) conduction band. The bandgap is situated in between these levels, and doping introduces positive electronic carriers (holes). This p-type conduction behavior was earlier found in LuRhO3 [37], but also in ZnCo2O4 with Co3+ in octahedral coordination [38]. The energy split in these materials however, was too low to form a bandgap that is large enough for visible transparency.

Based on the results on ZnCo2O4 and ZnRh2O4 the same behavior for compounds containing other d6-cations is expected. Besides Co 3d6 and Rh 4d6, the only candidate that allows field splitting with a “quasi-closed” shell is Ir 5d6. Moreover, the splitting between the t2g and eg levels is supposed to increase on going from 3d-<4d-<5d-metals, i.e. going down in group IX of the periodic table form cobalt, rhodium to iridium. According to Jørgensens empirical formula, the ligand field splitting can be estimated from ∆=f•g [35]. The values for g increase from 2.26 eV to 3.35 eV and 3.97 eV in the case of Co(III) to Rh(III) and to Ir(III) respectively. However, oxygen as a ligand can produce deviations from Jorgensen formula. From a rough estimation of the position of oxygen within the spectrochemical series, the value of f should be at least 0.9, but variations up to 0.37 are observed [39]. It was suggested that different bonding effect due to octahedral distortion and varying oxygen-cation bond length is responsible for this. However, the value of f is supposed to be fairly equal in all three cases of ZnM2O4. The numbers suggest an increase of ∆ with 50% and 25% from 3d- to 4d- and to 5d-. Considering the value of f to be 0.9 eV, the ligand field splitting in ZnCo2O4, ZnRh2O4 and ZnIr2O4 would be 2.03 eV, 3.01 eV and 3.57 eV, correspondingly.

According to these numbers, ZnIr2O4 would be a good candidate as p-type TCO when besides the large bandgap, reasonable transmittance and conductivity can be obtained. This material and the dependence of the bandgap on the nature of the d6-cation will therefore be investigated. Films of the three spinel ZnM2O4 compounds (where M=Co, Rh and Ir) are deposited by PLD on different substrates.


This allows examination of the electrical properties in relation to the crystallinity of these thin films. 6.3 Bulk ZnM2O4 synthesis

Commercially available powdersIII of zinc oxide and the transition metal oxides CoO, Rh2O3 and IrO2 were used to prepare the spinel ZnM2O4 phase by solid-state synthesis. The stoichiometric mixtures were ball-milled for 24 hours in ethanol. After drying, the powders were pressed to pellets; uniaxially pre-pressed up to 2000 bar, and subsequently isostatically pressed at 4000 bar. The 16 mm diameter pellets were sintered at temperatures ranging from 700 °C to 1050 °C in atmospheric conditions. Subsequently, the targets were crushed to fine powders again, and the procedure from ball-milling to sintering was repeated. Up to three of these cycles were used in order to obtain phase pure material. Finally, the resulting pressed pellets are used as targets for thin film deposition.

Prior to thin film fabrication, the phase of the targets was checked by XRD measurements (figure 6.2). All peaks of the ZnO-(CoO)2 and ZnO-Rh2O3 mixtures

III Powders were obtained from Alfa Aesar GmbH and Sigma-Aldrich Co. Purities; ZnO:

99.999%, Co(II)O: 99.99%, Rh2O3: 99.9%, Ir(IV)O2: 99.9%

Figure 6.2: XRD spectra of ZnM2O4 deposition targets, where M is the transition metal as indicated in thegraph.

10 20 30 40 50 60 70

ZnO

(002

)

IrO2 (

110)

M=Ir

M=Rh

M=Co

(311

)(3

11)

(222

)(2

22)

(220

)(2

20)

(111

)(1

11)

c/s

2θ (º)

Chapter 6 98

can be indexed and are identified as spinel ZnCo2O4 and ZnRh2O4 respectively [40,41]. However, bulk ZnIr2O4 could not be synthesized by solid-state synthesis. The peaks of both binary oxides are still visible in the diffraction spectrum, indicating a ZnO and IrO2 mixture in this target. Further attempts to induce a reaction via chemical complexation routes were madeIV. However, also these routes for single phase synthesis were unsuccessful. For the thin film fabrication of ZnIr2O4 this multi-component ZnO:(IrO2)2 target will be used.

The densities of the obtained targets appeared to be quite different. Whereas ZnCo2O4 target was rather compact (77%), the density of ZnRh2O4 and ZnO:(IrO2)2 targets was as low as 51% and 57% respectivelyV. SEM micrographs of the target surface reveal an open and porous structure containing loosely bound particles as large as 1 µm (figure 6.3b). SEM backscattered imaging revealed that in the case of ZnO:(IrO2)2 these particles mainly consist of different materials, i.e. ZnO and IrO2.

IV Bulk preparation of ZnIr2O4 was also attempted by wet chemical synthesis using EDTA

and citric acid complex agents. V Percentages of actual densities were calculated from the theoretical densities. Used

values; ZnCo2O4: 6.19 g/cm3, ZnRh2O4: 7.23 g/cm3 and ZnO:(IrO2)2: 9.62 g/cm3.

×20,000 1µm 10kV ×450 50µm 10kV

Figure 6.3: SEM micrographs with of the ZnO:IrO2 target surface showing the open porous structurecontaining particles up to 1 µm. The magnifications used are 450x (a) and 20.000x (b) respectively.

(a) (b)


6.4 ZnM2O4 growth and analysis 6.4.1 Thin film deposition

As stated in chapter 4, the laser fluency is an important parameter considering ablation from a multi-component target as the ZnO:(IrO2)2 pellet. The different oxides have different ablation rates, resulting in non-stoichiometric ablation [42]. The laser fluency should be chosen high enough in order to ablate both oxides. Although the ZnRh2O4 target can be considered as single-phase, ablation at low fluency also leads to non-stoichiometric ablation. Selective ablation occurs at the target surface leaving rhodium enriched particles. These particles provide shielding to the underlying target material leading to the formation of pillars (figure 6.4a). After several laser pulses, more target material is blocked from laser irradiation until the ablation process finally stops. Increasing the laser fluency to a higher value of 6 J/cm2 caused the particles to evaporate or to eject from the target surface (figure 6.4b, c and d).

For all three compounds the fluency was kept at 6 J/cm2 at a laser repetition rate of 10 Hz. High partial oxygen pressures are used to completely oxidize the material in order to avoid the formation of oxygen vacancies. The oxygen vacancies are so-called “killer-defects”, annihilating the positive charge carriers required for p-type conductivity [19]. Moreover, the incorporation of excess oxygen as interstitials enhances p-type conductivity as these defects add positive

×1000 10µm 18kV ×1000 10µm 18kV

×1000 10µm 18kV ×1000 10µm 18kV

Figure 6.4: SEM micrographs of the ZnRh2O4 target surface after 200 laser pulses at a fluency of3 J/cm2 (a), 4 J/cm2 (b), 5 J/cm2 (c) and 6 J/cm2 (d).

(a) (b)

(c) (d)

Chapter 6 100

holes to the valence band. The oxygen partial pressure is 0.250 mbar, and is mixed with Argon to a total background pressure of 0.300 mbar. The target-substrate distance was kept at 5 mm in all experiments.

Films with typical thickness of 100-300 nm are deposited on Al2O3 (0001) and quartz substrates. The substrate temperature was kept at 973 K for ZnRh2O4 and ZnIr2O4 films, and a lower temperature of 773 K was used for ZnCo2O4 films. For ZnCo2O4 films grown at higher temperature, additional high intensity diffraction peaks of ZnO are identified. This suggests the loss of cobalt due to the high volatility of this material during deposition, causing non-stoichiometry. Reducing the deposition temperature to 773 K solved this problem. 6.4.2 Structural analysis The diffraction spectra of ZnM2O4 films on Al2O3 (0001) are shown in figure 6.5a. The (111), (222) and (333) peaks of the spinel-structure are visible for all compounds, indicating (111) oriented coherent growth. Quartz glass was used to grow films without any epitaxial relation to the substrate. Figure 6.5b shows the polycrystalline nature of ZnM2O4 films deposited on quartz. All peaks in these films correspond to those found in the target (figure 6.2). These samples are measured in 1.5º grazing incidence geometry. Although a single phase ZnIr2O4 target could not be obtained, both epitaxial and polycrystalline thin films grown by PLD consist of the correct phase.

Table 6.1: Calculated average lattice parameters of ZnM2O4 films compared to literature values and ionic radii.

ZnM2O4 Al2O3 a (Ǻ)

quartz a (Ǻ)

Lit. value a (Ǻ)

Ionic radius (pm)

ZnCo2O4 8.079 8.104 8.108 68.5 ZnRh2O4 8.482 8.489 8.510 80.5 ZnIr2O4 8.503 8.507 n.a. 82.0

The obtained lattice parameters for ZnCo2O4 and ZnRh2O4 correspond well

(table 6.1) with the values found in literature [40,41]. The lattice spacing of ZnIr2O4 almost equals that of ZnRh2O4. No literature data is present for comparison; however, this can be expected as the size of the ionic radii of rhodium and iridium trivalent cations in a 6-coordinated octahedral environment is approximately the same. These are the first spinel ZnIr2O4 films synthesized.


10 20 30 40 50 60 70A

l 2O3

Al 2O

3Al

2O3

(333

)

(333

)(3

33)

(222

)(2

22)

(222

)

(111

)(1

11)

(111

)

M=Co

M=Rh

M=Ir

c/s

2θ (º)

10 20 30 40 50 60 70

(222

)(2

22)

(311

)(3

11)

(220

)(2

20)

(111

)(1

11)

(222

)(311

)

(220

)(111

)

M=Co

M=Rh

M=Ir

c/s

2θ (º)

Figure 6.5: XRD of with of (111)- oriented epitaxial ZnM2O4 films deposited on Al2O3 (0001) substrates (a), and polycrystalline films deposited on quartz (b). All the peaks can be indexed as spinelphase. M corresponds to the transition metal as indicated in the graph. Polycrystalline filmsare measured in 1.5º grazing incidence geometry. The halo peak around 20º originates fromthe amorphous quartz substrate. The very small intensity to the left of the (222) diffractionpeaks originates from ZnO (002), indicating slightly off-stoichiometric growth.

(a)

(b)

Chapter 6 102

6.4.3 Morphology and stoichiometry The AFM measurements (figure 6.6a) of ZnM2O4 films show smooth films

with a RMS surface roughness of about 2 nm. Although the film grows epitaxial, the imprint of the stepped sapphire substrate is no longer visible. Grains as large as 80 nm are detected on these films from cross section analysis. However, during AFM imaging sometimes very large particles are detected on the surface. These features are too large for clear imaging by AFM. Therefore SEM measurements are used to reveal the natures of these particles. These experiments show that droplets as large as 1 µm are present on the film surface, especially on rhodium and iridium containing films. These particles are likely to be originating from the porous nature of the targets of these materials.

XPS is used to determine the atomic ratios between the Zn and M cations. In figure 6.7a, a typical survey scan of ZnIr2O4 is shown in which each element gives rise to a characteristic set of peaks. From precise scans of the different elements, a small off-stoichiometry is detected. The found Zn:M ratio is sometimes found to be as low as 1.7 (instead of 2). Additionally, small ZnO (002) diffraction peak intensity can sometimes be seen in the XRD spectra of these films. This indicates that the particles mainly consist of Co-, Rh- and Ir-metal-oxides.

×2000 10µm 10kV

(a) (b)

Figure 6.6: AFM micrograph of a ZnRh2O4 film imaging an 1x1 µm2 area of the film surface which is clean from droplets (a). SEM image of ZnIr2O4 film surface revealing these droplets (b). Both films are deposited on a sapphire substrate.


6.4.4 Electronic band structure Besides compositional analysis, XPS can be used to probe the valence

band levels. The profile close to the Fermi level (0 eV) in figure 6.7a provides insight on the density of electronic states (occupied states). Low energetic radiation is used that is only capable of ionizing electrons from the outermost levels of atoms, i.e. the valence levels. In figure 6.7b the results on the electronic structure of the spinel ZnM2O4 polycrystalline thin films are shown. Ohmic contact to the film surface and argon ion gun was used for neutralizationVI. The valence band structures of the three different compounds look similar and are comparable to the measurements on ZnRh2O4 [29]. From measurements on ZnO, the band near 10 eV can be identified as Zn 3d10. The diffuse region around 5 eV originates from O 2p and transition metal band mixing. The high intensity at 1 eV is composed of the t2g

6 states due to the ligand field splitting of d6-bands of the transition metals. The Fermi level is located at the edge of this valence band maximum. This is an

VI The C 1s peak appeared at 285eV as expected. This indicates that the loading of the

sample surface is avoided.

1200 1000 800 600 400 200 0

O K

LL

Zn L

MM

1 Ir 4f

Zn 3

pZn

3s

C 1

sIr

4d5

Ir 4d

3

Ir 4p

3O 1

s

Zn 2

p3

c/s

Binding energy (eV)

Zn 2

p1

Figure 6.7: XPS survey scan of ZnIr2O4 polycrystalline sample. Aluminum irradiation of 224 eV is used as excitation source (a). XPS spectra of ZnM2O4 polycrystalline films, where M is the transition metal as indicated in the graph. Low energetic (55 eV) x-rays are used to probe the valence level. Scans are averaged over 5 cycles with a resolution of 0.1 eV (b).

12 10 8 6 4 2 0 -2 -4

M=Ir

M=Rh

t62g

O 2p

Zn 3d10

N

orm

aliz

ed in

tens

ity (c

/s)

Binding energy (eV)

Ef

M=Co

(a) (b)

Chapter 6 104

indication for p-type conductivity as the Fermi level shifts downward as acceptor states are introduced in the bandgap close to the valence band edge.

A small shift in the position of Zn 3d10 to lower binding energies is observed for ZnCo2O4. This can be attributed to the shorter Zn-O bonding length in this material, raising the Zn 3d10 electrons in energy. Also a sharper valence band is observed for this compound, indicating that the Co 3d6 electrons are more localized than Rh 4d6 and Ir 5d6. 6.5 Optical and electrical properties 6.5.1 Optical analysis

The optical data in figure 6.8 shows that the polycrystalline samples containing rhodium and iridium are reasonably transparent in the visible region (400-700 nm). The average transmission of these samples on quartz is close to 60%, whereas that of ZnCo2O4 is below 30%. This compound can therefore not be regarded as TCO. The onset of absorption of ZnM2O4 films is less sharp compared to the n-type TCOs discussed in chapter 4 and 5. As the wavelength decreases, the transmission drops to zero over more than 100 nm. This is attributed to the broad shape of the t2g valence band. Electrons can be promoted from the top, as well as the bottom of this band, giving rise to a larger range of wavelengths of absorption.

The optical gap of the three compounds is listed in table 6.2. These values are estimated from extrapolation of the absorption spectrum (figure 6.8b) [43]. The

1.5 2.0 2.5 3.0 3.5 4.00

1

2

3

4

5

(3)(2)(1)

α2 (x

1010

cm

-1)

E (eV)

200 400 600 800 1000 12000

10

20

30

40

50

60

70

(3)(2)

%T

wavelength (nm)

(1)

Figure 6.8 Optical transmission measurements of ZnCo2O4 (1), ZnRh2O4 (2) and ZnIr2O4 (3) polycrystalline films deposited on quartz (a). Photon energy dependence of α2 and theextrapolation of the curve for deducing the bandgap for the same compounds (b).

(a) (b)


primary gap of ZnCo2O4 is located around 2.26 eV, which is somewhat lower than 2.63 eV reported by Kim et al. [38]. It is possible that these films were not completely stoichiometric since the bandgap increases as the concentration of cobalt is lowered [44]. A similar absorption feature at lower energy (~1.5 eV) is detected. This absorption band is attributed to Co inter-atomic d-d transitions associated with a trigonal ligand field splitting [45]. This indicates that a small part of the cobalt ions have interchanged with the tetrahedrally coordinated zinc ions. The cobalt ion can also exist in this coordination with similar ionic radius of zinc ions [46].

The gap of ZnRh2O4 of 2.74 eV is larger than earlier observations (~2.1 eV) [29]. The slope of the absorption spectrum, which is different in our case, can be responsible for this deviation. The obtained optical gap for ZnIr2O4 is 2.97 eV, making this material most transparent in the visible region. The obtained bandgap values scale reasonable with the increase of ∆ as predicted from the empirical relation described in section 6.2. This prediction is indicated in the table as 0.9•g, in the last column the percentage of deviation between the measured optical gap and predicted gap is listed.

Table 6.2: Optical properties of ZnM2O4 polycrystalline films

ZnM2O4 %T (400-700 nm)

E gap (eV)

0.9•g (eV)

Dev. (%)

ZnCo2O4 26.1 2.26 2.03 -11 ZnRh2O4 54.8 2.74 3.01 +9 ZnIr2O4 58.8 2.97 3.57 +17

6.5.2 Electrical conduction behavior

The electrical properties of the deposited films are listed in table 6.3. All our samples were found to be highly conductive, in the order of ~1 Scm-1. Results of ZnCo2O4 and ZnRh2O4 are comparable to the values earlier found by others [38,29]. The conductivity of polycrystalline ZnIr2O4 is also found in this range, which indicates that these three systems are electronically equivalent. Furthermore, the conductivity of epitaxial and polycrystalline samples is almost equal, indicating that the structural properties of the film does not influence the electrical properties. In ZnCo2O4, a relation between the conduction type (p or n) and the oxygen deposition pressure was shown [38,47]. Considering the high oxygen deposition pressure used here, doping is introduced by excess oxygen. Additionally, cation vacancies can act as acceptors in the system.

Chapter 6 106

Table 6.3: Electrical properties of ZnM2O4 films polycrystalline films. The conductivity refers to polycrystalline as well as epitaxial films.

ZnM2O4 σ (poly- /epi-) (Scm-1)

Seebeck (µV/K)

Ea

(meV) ZnCo2O4 0.39 / 0.61 +131.4 41.7 ZnRh2O4 2.75 / 2.83 +63.4 22.1 ZnIr2O4 3.39 / 2.09 +53.9 46.9

Although the conductivity at room temperature is rather large, no significant

data concerning the carrier concentration can be extracted from Hall measurements. The slope of the line that relates the Hall voltage VH to the magnetic field B becomes considerably low. From equation 3.1 it becomes evident that this is due to a very high doping concentration. The average slope fitted to the data of the samples indicates a carrier concentration in the order of 1019 cm-3. But the uncertainty in these values is significant, more than one order of magnitude. The value for the corresponding Hall mobility is in the order of 10-1 cm2V-1s-1. Despite the uncertainty in the mobility, this value is quite common for (poly) crystalline p-type TOSs. However, one should be careful drawing conclusions from Hall

0 5 10 150

1

2

+54µV/K

+63µV/K

V (m

V)

∆T (K)

+131µV/K

4 6 8 10-3

-2

-1

0

1100K150K200K

log

σ (S

cm

-1)

1000/T (K-1)

300K

Figure 6.9 Thermo power measurements of ZnCo2O4 (), ZnRh2O4 () and ZnIr2O4 () thin films fitted to a straight line for deducing the Seebeck coefficient (a). Temperature dependantconductivity measurements of the same films

(a) (b)


measurements and their interpretation on semiconductors with low and relatively high carrier concentration [48]. Mizoguchi et al. reported that no significant Hall data for the ZnRh2O4 samples could be obtained, whereas Kim et al. did report on carrier concentration and mobility for ZnCo2O4, 2.81 x1020 cm-3 and 0.2 cm2V-1s-1 respectively [29,38].

The positive sign of the slope in the Hall measurements indicated p-type conductivity, as earlier found by PES measurements. Considering the fact that the Hall data may be doubtful, thermo-power measurements are an outcome to obtain conclusive data on conduction type behavior. This is confirmed by the low values found for the Seebeck coefficient, induced by higher hole concentrations, compared to other p-type oxide semiconductors as cuprates [49,50]. The Seebeck coefficients measurements indicate p-type conduction behavior in all cases (figure 6.9a).

Temperature dependant measurements show an Arrhenius type behavior (figure 3b). The thermal activation energy is estimated from the high temperature region (above 200 K) and was found to be smaller than 50 meV (table 6.3). So the Fermi level is located close to the valence band in which the dominant carriers (holes) are transported. Furthermore, the low activation energies are an indication that high carrier concentrations can indeed exist in these materials.

6.5.3 Discussion

From the experiments it appears that the bandgap of ZnM2O4 thin films is dependant on the nature of the d6- cationic species M. For M= Rh and Ir, the films are optically transparent and a high p-type conductivity is observed at room temperature [51].

The exact bandgap value is determined by M, and can be related to the splitting of the d-orbitals due to the anionic ligand field. The obtained values differ within 20% of the prediction from ligand field theory. Some deviation can be expected as the optical gaps are merely estimation from the absorption spectra. In these materials the onset of absorption is less steep compared to those of ITO or ZnO. The t2g and eg states are rather smeared out in energy, giving rise to a less pronounced effective bandgap. The extrapolation of the Tauc plot is therefore less accurate. Furthermore, measured optical gaps are in many cases different from the electronic gap. Defect levels within the gap enhance the absorption of lower energetic light.

Also the earlier mentioned uncertainty of f for oxygen ligands can be responsible for incorrect predictions. The used value of 0.9 is merely estimation of the position of oxygen within the spectrochemical series. Moreover, different bonding lengths due to octahedral distortion can have an effect on the exact value [39].

Chapter 6 108

The Hall data did not give any conclusive results on the hole carrier density or mobility. However, besides the electronic band structure, also the Seebeck measurements confirmed p-type conductivity. Room temperature conductivity up to 3 Scm-1 is observed, which is very high for p-type TOSs (table 2.2). An important aspect is that the electrical properties are independent of the structural nature of the deposited films. No significant difference in conductivity can be observed between epitaxial or polycrystalline samples. This means that the electrical properties are at least not affected by the grain boundaries. This partly confirms the earlier statement that these materials are less affected by structural disorder.

The electronic and optical properties of spinel ZnM2O4 is only determined by the MO6-octahedra. The valence band consists of the t2g sub-bands which are extended and rather isotropic and establish the hole conduction [30,52]. The octahedra form a network of overlapping d-bands despite the lower crystalline nature of the thin film. This is very promising with regard to deposition of these materials on polymers at low temperatures, which consequently will lead to very low crystalline thin films. These depositions, resulting in grain sizes far smaller than obtained here, should be performed to investigate if this material is still highly conducting in nanocrystalline phase. 6.6 Concluding remarks

The success of transparent electronics is dependant on the scientific improvement on p-type TCOs. Cuprates as CuAlO2 or SrCu2O2 and p-ZnO are well-known p-type TCOs, and are used to build p-n junctions. The properties of these materials however, are very dependant on their structural nature. A large process temperature is generally required to deposit highly crystalline thin films. Obviously, these conditions are not suited for deposition on polymers.

The degree of structural ordering in spinel ZnM2O4 (M=Co, Rh and Ir) can be quite low, whereas high p-type conductivity is maintained. The d6- bands of the transition metals in octahedral coordination are split due to the field strength of the oxygen ligands. This splitting mainly determines the bandgap. From the ligand field theory, an estimation of the bandgap can be obtained.

Films of spinel are prepared by PLD at 773 K and 973 K on quartz and sapphire substrates. Although no phase pure ZnIr2O4 could be obtained, XRD measurements showed that the phase of both the polycrystalline and epitaxial films was spinel for all compounds. From PES measurements, the electronic band structure is analyzed. The valence band maximum consists of the transition metal t2g subset and is located close to the Fermi level, indicating p-type conductivity.

An expected increase of this gap for the heavier ions was confirmed by optical measurements and can be related to ligand field theory. The gap increases


on going from Co to Rh and to Ir. The largest observed gap of ZnIr2O4 is ~3 eV, whereas the average optical transmission is around 60%. The electrical conductivity seems to be independent on the degree of crystallinity of the thin films. ZnIr2O4 room temperature conductivity of polycrystalline samples was above 2 Scm-1. Furthermore, the Seebeck coefficients of all samples is positive, confirming hole conduction. The conductivity in these materials is high because of a high carrier concentration.

From the electrical and optical properties of ZnRh2O4 and ZnIr2O4 it can be concluded that p-type thin film TOSs of good quality are deposited. In particular for applications that require a low process temperature these materials are excellent candidates.

Chapter 6 110

6.7 References 1. G. Thomas, Nature 389, 907 (1997). 2. J.F. Wager, Science 300, 1245 (2003). 3. K. Nomura, H. Ohta, K. Ueda, T. Kamiya, M. Hirano & H. Hosono, Science

300, 1269 (2003). 4. S. Masuda, K. Kitamura, Y. Okumura, S. Miyatake, H. Tabata & T. Kawai, J.

Appl. Phys. 93, 1624 (2003). 5. R.L. Hoffman, B.J. Norris & J.F. Wager, Appl. Phys. Lett. 82, 733 (2003). 6. P.F. Carcia, R.S. McLean, M.H. Reilly & G.Nunes, Appl. Phys. lett. 82, 1117

(2003). 7. R.E. Presley, D. Hong, H.Q. Chiang, C.M. Hung, R.L. Hoffman & J.F. Wager,

Solid-State Electron. 50, 500 (2006). 8. H. Kawazoe, H. Yanagi, K. Ueda & H. Hosono, MRS Bull. 25, 28 (2000). 9. H. Ohta, K.-I. Kawamura, M. Orita, M. Hirano, N. Sarukura, & H. Hosono,


239, 496 (2002). 12. A. Tsukazaki, A. Ohtomo, T. Onuma, M. Ohtani, T Makino, M. Sumiya, K.

Ohtani, S. F. Chichibu, S. Fuke, Y. Segawa, H. Ohno, H. Koinuma & M. Kawasaki, Nat. Mater. 4, 42 (2005).

13. S.J. Jiao, Z.Z. Zhang, Y.M. Lu, D.Z. Shen, B. Yao, J.Y. Zhang, B.H. Li, D.X. Zhao, X.W. Fan, & Z.K. Tang, Appl. Phys. Lett. 88, 031911 (2006).

14. A.N. Banerjee & K.K. Chattopadhyay, Prog. Cryst. Growth Charact. Mater. 50, 52 (2005).

15. H. Kawazoe, M. Yasukawa, H. Hyodo, M. Kurita, H. Yanagi & H. Hosono, Nature 389, 939 (1997).

16. A. Kudo, H. Yanagi, H. Hosono & H. Kawazoe, Appl. Phys. Lett. 73, 220 (1998).

17. D.M. Bagnall, Y.F. Chen, Z. Zhu, Y. Yao, S. Koyama, M.Y. Shen & T. Goto, Appl. Phys. Lett. 70, 2230 (1997).

18. P. Zu, Z.K. Tang, G.K.L. Wong, M. Kawasaki, A. Ohtomo, H. Koinuma & Y. Segawa, Solid State Commun. 103, 459 (1997).

19. A. Zunger, Appl. Phys. Lett. 83, 57 (2003). 20. A. Kobayashi & O.F. Sankey, J.D. Dow, Phys. Rev. B 28, 946 (1983). 21. K. Minegishi, Y. Koiwai, Y. Kikuchi & K. Yano, Jpn. J. Appl. Phys. 36, L1453

(1997). 22. Y.R. Ryu, S. Zhu, D.C. Look, J.M. Wrobel, H.M. Joeng & H.W. White, J.

Cryst. Growth 216, 330 (2000).


23. M. Joseph, H. Tabata, H. Saeki, K. Ueda & T. Kawai, Physica B 302-303, 140

(2001). 24. X.-L. Guo, H. Tabata & T. Kawai, J. Cryst. Growth 223, 135 (2001). 25. K.-K. Kim, H.-S. Kim, D.-K. Hwang, J.-H. Lim & S.-J. Park, Appl. Phys. Lett,

83, 63 (2003). 26. H. Gong, Y. Wang & Y. Luo, Appl. Phys. Lett. 76, 3959 (2000). 27. S. Gao, Y. Zhao, P. Gou, N. Chen & Y. Xie, Nanotechnology 14, 538 (2003). 28. D.S. Ginley & C. Bright, MRS Bull. 25, 15 (2000). 29. H. Mizoguchi, M. Hirano, S. Fujitsu, T. Takeuchi, K. Ueda & H. Hosono, Appl.

Phys. Lett. 80, 1207 (2002). 30. T. Kamiya, S. Narushima, H. Mizoguchi, K.-I. Shimizu, K. Ueda, H. Ohta, M.

Hirano & H. Hosono, Adv. Funct. Mater. 15, 968 (2005). 31. L. Pauling, in The Nature of the Chemical Bond, Cornell University Press:

Ithaca, 3rd ed. 1960. 32. H. Bethe, Ann. Physik 3, 135 (1929). 33. C.J. Ballhausen, In Introduction to ligand field theory, McGraw-Hill Book

Company, USA (1962). 34. R. Krishnamurthy, W.B. Schaap & B. Ward, J. Chem. Educ. 46, 799 (1969). 35. C.K. Jørgensen, in Modern aspects of ligand field theory, 1st edition, North-

Holland Publishing Company, Amsterdam, Netherlands (1971), Ch. 26. 36. B.N. Figgis & M.A. Hitchman, in Ligand field theory and its applications, Chap

7, Wiley-VCH, New York (2000). 37. H.S. Jarret, A.W. Sleight, H.H. Kung, & J.L. Gibson, J. Appl. Phys. 51, 3916

(1980). 38. H.J. Kim, I. C. Song, J. H. Sim, H. Kim, D. Kim, Y. E. Ihm & W. K. Choo, J.

Appl. Phys. 95, 7387 (2004). 39. C. E. Schäffer, J. Inorg. Nucl. Chem, 8, 149 (1958). 40. Natl. Bur. Stand. (U.S.) Monogr. 25 10, 60 (1972). 41. Schulz, Eysel, Mineral.-Petrographisches Institut der Univ. Heidelberg,

Germany. ICDD Grant-in-Aid (1989). 42. L.M. Doeswijk, in Pulsed laser deposition of oxides on silicon: Exploring their

passivating qualities, Chap.2, PhD thesis, University of Twente, The Netherlands (2002).

43. J. Tauc & R. Grigorovici, A. Vancu, Phys. Status Solidi 15, 627 (1966). 44. K. Samanta, P. Bhattacharya, R.S. Katiyar, W. Iwamoto, R.R. Urbano, P.G.

Pagliuso & C. Rettori, Mater. Res. Soc. Symp. 891, 0891-EE10-09.1 (2006). 45. K. Samanta, P. Bhattacharya & R.S. Katiyar, Appl. Phys. Lett, 87, 101903

(2005). 46. R.D. Shannon & C.T. Prewitt, Acta crystallogr. B 25, 925 (1969).

Chapter 6 112

47. H.J. Kim, I.C. Song, J.H. Sim, H. Kim, D. Kim, Y.E. Ihm & W.K. Choo, Phys.

Stat. Sol. b 241, 1553 (2004). 48. D.C. Look, D.C. Reynolds, C.W. Litton, R.L. Jones, D.B. Eason & G.

Cantwell, Appl. Phys. Lett. 81, 1830 (2002). 49. H. Kawazoe, M. Yasukawa, H. Hyodo, M. Kurita, H. Yanagi & H. Hosono,

Nature 389, 939 (1997). 50. A. Kudo, H. Yanagi, H. Hosono & H. Kawazoe, Appl. Phys. Lett. 73, 220

(1998). 51. M. Dekkers, G. Rijnders & D.H.A. Blank, Appl. Phys. Lett. 90, 021903 (2007). 52. S. Narushima, H. Mizoguchi, H. Ohta, M. Hirano, K.-I. Shimizu, K. Ueda, T.

Kamiya, & H. Hosono, Mat. Res. Soc. Symp. Proc. 747, V2.2.1 (2003).

Chapter 7

Amorphous p-type TOS on polymer substrates

Abstract

P-type TCOs are essential for electronic devices based on wide bandgap semiconductors. For applications on polymeric substrates, the low process temperature implies amorphous materials. ZnO:x·IrO2 is a p-type transparent semiconductor with excellent properties even in a nanocrystalline phase. The majority of iridium in the film appears in a trivalent valency. The film is transparent up to 70%, and the p-type conductivity is as large as 2 Scm-1. The film is used as p-layer in a junction device, and diode rectifying behavior is observed. Furthermore, the applicability of this material in all-amorphous devices on polymer substrates is demonstrated.

Chapter 7 114

7.1 Introduction With regard to processing, the main advantage of amorphous

semiconductors is that low temperatures can be employed. Consequently, the manufacturing cost is decreased, and heat resistant substrates such as polymers can be utilized. Amorphous (or nanocrystalline) n-type TCOs are demonstrated to be suitable for this purpose as described in chapter 4 and 5 (and the references therein). However, not many devices based entirely on amorphous oxide semiconductors (AOS) are reported.

Mizoguchi et al. [1] was the first to show p-type conductivity in ZnRh2O4 thin films. No long range order is required for hole conductivity in this material [2]. Based on the knowledge of ZnRh2O4 polycrystalline films, this material is expected to be a transparent hole conductor even in the nanocrystalline/amorphous phase. These ZnO:Rh2O3 films grown at room temperature were x-ray amorphous, but still displayed p-type conductivity up to 1 Scm-1. The existence of RhO6 octahedral networks was claimed to form conduction paths within the amorphous phase. Using this material, diode rectifying behavior has been observed in the only all-amorphous p-n junctions reported so far [3,4].

Spinel ZnIr2O4 is related to ZnRh2O4 as illustrated in chapter 6, and the optical properties are even shown to be better. Because of its similarity with the rhodium based material, the iridium containing compound is expected to have the same electrical behavior when grown at room temperature. Moreover, the higher bandgap of polycrystalline ZnIr2O4 can contribute to a higher visible transmission in the amorphous state as well.

In this chapter, experiments on ZnIr2O4 thin films deposited on polymer substrates in order to obtain a p-type TCO is described. A special deposition technique in order to obtain a smooth film surface without droplet formation is discussed in section 7.2. Next, the structure and composition of these films is explained by the valency of iridium cations. In section 7.4 the electrical and optical properties of these films are presented. Finally, the material is implemented as a transparent p-layer in amorphous p-n heterojunctions deposited on n-type silicon and polymer substrates. 7.2 Thin film deposition by “eclipse” PLD

Targets of ZnO:(IrO2)2 powder mixtures, prepared as described in section 6.2 were used in the experiments. Except for the substrate temperature, the parameters for room temperature deposition were kept equal to the high temperature experiments. That means a fluency of 6 J/cm2 at a laser repetition rate of 10 Hz and a target substrate distance of 50 mm. Unless stated otherwise the oxygen partial pressure was 0.250 mbar, whereas the total pressure was 0.300

Amorphous p-type TOS on polymer substrates 115

×7,500 2µm 10kV

Figure 7.1: SEM micrograph of the ZnO:IrO2 thin film surface deposited on silicon substrate at roomtemperature.

mbar. For the experiments, PET and quartz are used as substrates. Prior to deposition the substrates are ultrasonically cleaned in ethanol, and subsequently blown dry by nitrogen. The surface of the as deposited films was checked by SEM. The resulting SEM micrographs reveal particles on the film surface ranging from tenths of nanometers up to dimensions as large as 1 µm (figure 1a). Although some droplet formation was also detected on high temperature grown films, a substantial amount of particles has formed on this room temperature grown film. Depositions performed at lower total pressure, 0.050 mbar instead of 0.300 mbar, did not decrease the droplet formation. This indicates that the particles are not formed in the plasma, which is possible in high background pressure [5,6], but they are more likely to originate directly from the target. Because of the micro-structural nature of the of ZnO:(IrO2)2 target as stated in chapter 6 (figure 6.3), single particles as large as 1 micron can be ejected from the target instead of being completely evaporated. However, it is not clear why the number of droplets on the film surface is substantially higher for room temperature grown films compared to high temperature grown films. One possibility is the difference in sticking coefficient at high and low temperature. But small droplets may also desintegrate at a temperature of 700 ºC, since the melting point of IrO2 is only 1070 ºC.

The formation of droplets is a known issue and can be an obstacle for the use of PLD in commercial applications [7]. A large number of particles are not desired, especially when the film is part of a stacked multilayered structure as in a p-n junction. Other than changing the target properties to one single phase dense pellet, which proved impossible for ZnO:(IrO2)2, Kinoshita et al. [8] developed a

Chapter 7 116

deposition method to reduce droplet formation. By simply placing a mask between the target material and substrate, they eliminated the particle problem peculiar to the excimer laser deposition of YBa2Cu3O y thin film deposition. This technique is tentatively called the “eclipse method”, referring to the mask shielding the substrate from the irradiated area of the target from which the plasma plume originates. The major advantage of the eclipse configuration is that it is much simpler than using particle velocity filters [9] or the “crossed fluxes technique” [10] serving the same purpose, and can therefore be easily implemented in the PLD system. A schematic representation and working principle of “eclipse PLD” is illustrated in figure 7.2:

The actual deposition takes place only by diffusion. Due to the pressure gradient between the plasma plume and the area behind the mask, light atomic species and clusters can be scattered to the substrate. The substrate is directly shielded from the large and more heavy particles that either are blocked by the mask, or are not scattered to the substrate because of its larger weight. The drawbacks of “eclipse PLD” are a large decrease in deposition rate, but also thickness homogeneity and stoichiometry can be affected [11,12], especially when heavy atomic species are involved.

A mask was placed in the system at 3/5 of the target substrate distance of 50 mm, i.e. at 20 mm in front of the substrate. According to earlier observations this distance would give the best results for pressures higher than 0.1 mbar [11]. The mask size has the dimensions of 6x6 mm2, since it than exactly shadows the 10x10 mm2 substrate. In order to check thickness homogeneity, a grid structure was patterned on a large silicon substrate by photolithographic lift off technique. Subsequently, the step height of the grid pattern is measured at different positions on the substrate. The resulting thickness profile is shown in figure 7.3. As expected,

Figure 7.2: Schematic representation of the setup and working principle of the eclipse PLD method toreduce the droplet formation.

Target

Substrate

Mask

Laser beam


a minimum thickness is observed in the center of the substrate. The thickness profile at the 5x5 mm2 area of interest behind the mask is rather uniform (~15%), and is comparable to the gradual drop off in conventional PLD [11]. The deposition rate has dropped to 25% of the value obtained from conventional PLD.

Besides the thickness uniformity, the first aim of “eclipse PLD” was the reduction of droplets on the film surface. The SEM micrograph in figure 7.4a shows again the film surface earlier shown in figure 7.1 of the zinc iridium oxide layer deposited by conventional PLD, whereas figure 7.4b shows the surface of a thin film

-6 -4 -2 0 2 4 60

100

200

300

400

500

600

700

Thic

knes

s (n

m)

Position (mm)

Figure 7.3: Thickness profile of a zinc iridium oxide film grown on silicon at room temperature by “eclipsePLD”. The position is indicated as distance from the centre of the substrate. In the graph, thesquares indicate the horizontal thickness profile, whereas the open circles denote thethickness in vertical direction from the substrate centre. The accuracy is within 10%.

×7,500 2µm 10kV ×7,500 2µm 10kV

Figure 7.4: SEM micrographs of the thin film surface after ablation from ZnO:IrO2 target without (a), and with using the PLD eclipse method (b).

(a) (b)

Chapter 7 118

grown using the “eclipse” technique. The deposition time during the “eclipse” deposition process was four times as large, resulting in films of equal thickness. From the results it is obvious that “eclipse PLD” is a useful tool to increase the surface smoothness. The number of droplets has decreased significantly leaving a flat surface. However, it is very likely that the total amount of material in the droplets do not consist of the same zinc-iridium ratio as detected in the target. It is therefore possible that a flat surface by “eclipse PLD” can only be obtained on the price of a disturbed film composition if the target stoichiometry is not compensated for. 7.3 Structural and chemical composition 7.3.1 Structural analysis

XRD measurements are performed in order to check the structural disorder of the room temperature grown zinc iridium oxide films on PET substrates. The crystalline nature of these films is expected to be very limited, or not present; i.e. the film is completely amorphous. Hence the films are measured using grazing incidence geometry at very low scan speeds. In order to give a clue on the nature of the particles on the film surface, also samples deposited by conventional PLD are measured. The results of these measurements are presented in figure 7.5.

10 20 30 40 50 60 70

Iridium

Eclipse PLD

Normaldeposition

c/s

2θ (°)

Figure 7.5: XRD spectra of ZnM2O4 thin films deposited at room temperature on PET substrates withoutand with using the PLD eclipse method. Measurements are done in 1.5 º grazing incidencegeometry; the scan speed is 0.15 º/min and the resolution is 0.01º. The halo peak around 20 originates from the PET substrate.


The film deposited using eclipse PLD do not show any observable sharp peak features, indicating that this film is (x-ray) amorphous. This result is similar to earlier observations of Narushima et al. [3] that room temperature grown zinc rhodium oxide films are amorphous. The XRD spectrum of the film grown by conventional PLD shows three low intensity peaks. These peaks are identified as the fingerprint for a metallic iridium phase. This indicates that the composition of the droplets is iridium-rich. Off-stoichiometry in the thin film grown by “eclipse PLD” can therefore be expected. 7.3.2 Stoichiometry of “eclipse deposited” films

Position dependant XPS measurements are performed to determine the chemical composition of the film behind the shadow mask. The result is shown in figure 7.6a. In the middle of the sample behind the eclipse mask the iridium-zinc ratio is only 1.2, instead of the stoichiometric ratio of 2 in the target. As the XPS measurement is performed closer to the edge of the “eclipsed” sample area, this ratio increases. However, a difference is observed between the ratios measured in horizontal and vertical position from the middle of the sample surface. This can be explained by the finite dimension of the spot size. The laser spot on the target should not be considered as a point source, but is the projection of a rectangular mask and is about 1x3 mm2. Whereas a square “eclipse” mask blocks exactly a

-4 -2 0 2 40.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8

2.0

[Ir] /

[Zn]

Position (mm)

(a)

(b)

Figure 7.6: Position dependant XPS measurements of ZnO:IrO2 film deposited by the eclipse method. The atomic ration [Ir]/[Zn] dependant on the position measured vertically () and horizontally () from the middle of the sample is shown (a), as well as a schematic representation of theplaces of XPS measurements on the sample (b).The square area indicates the eclipsed areaof the sample surface “seen” from a point source, whereas the rectangular dotted shaperepresents the area blocked by the actual laser spot.

maskspot

particle trajectrory

Chapter 7 120

square area on the sample in case of a point source, the shielded area is rectangular in case of an elongated laser spot. The particle trajectory takes of at a slightly smaller angle because of the material which is ejected from the target at the edge of the illuminated area. A schematically representation is drawn in figure 7.6b.

Because of the rectangular shape the points measured vertical from the middle are closer to the edge of the shielded area than the horizontal measured ones. The zinc-iridium ratio is therefore higher in the first case. It is not straightforward to measure the composition of the unshielded thin film, since here the iridium-rich particles are still present on the surface. The XPS measurement spot is too large for probing in between these particles, artificially increasing the iridium-zinc ratio. However, the measured ratio relatively close to the edge of the shielded region is expected to be almost equal to the unshielded thin film composition; i.e. a ratio of 1.6.

So the films deposited by the eclipse method are off-stoichiometric and should therefore be indicated as ZnIr2-δO4 instead of ZnIr2O4. The value of δ in ZnIr2-δO4 varies between about 0.3 and 0.8. The films deposited for the electrical and optical characterization are structured by a shadow mask as described in section 3.3.1. The sampled area of these patterned films is 5x5 mm2, which was located in the middle of the shielded area during deposition. This means that δ will be close to 0.8. On the other hand, the composition is reasonably homogeneous in this area. Obtained results can therefore be considered as that of a ZnIr1.3O4 film with a uniform composition. 7.3.3 Chemical composition

The room temperature deposited ZnIr2-δO4 films were found to be x-ray amorphous. Although the film can be nanocrystalline, it is not expected that the small grains consist of the spinel phase because of the non-stoichiometry. Besides the compositional ratio of zinc and iridium in the sample, XPS measurements can also be used to determine the nature of the chemical bonds within the sample. From the binding energy it is possible to investigate if the oxidation state of iridium is indeed trivalent, as is the case is spinel ZnIr2O4. Different samples are grown and analyzed to distinguish between the different oxidation states. Besides room temperature deposited ZnIr2-δO4, amorphous IrO2 and spinel ZnIr2O4 samples were used as reference. The spinel ZnIr2O4 film is deposited as described in section 6.4.1, and the IrO2 sample was deposited at the same conditions as ZnIr2-δO4 from an IrO2 target. This target was fabricated from pure IrO2 powder as described in section 6.3.


The Ir 4f level (as in figure 6.7a) is actually split into two lines (Ir 4f7/2 and Ir 4f5/2) located at 60.9 eV and 63.9 eV for metallic iridium, respectively. For oxygen bound iridium these lines are located at higher binding energies, and this value increases at higher oxidation states, i.e. on going form Ir(III) to Ir(IV) [13]. Generally, in case of coordination with more electronegative ligands the core level electrons of a central atom appear at higher binding energy. This so-called chemical shift may influence the energetic position of the emission peak up to several electronvolts, which can easily be resolved by the XPS equipment. The three different samples are all measured with respect to a fixed reference (531 eV of O 1s).

The measured value Ir 4f5/2 in the IrO2 film is found at 62.3 eV, which is consistent with observations by others [14,15]. On the other hand, the spectrum of the ZnIr2O4 film shows a combination of four peaks. These peaks can be deconvoluted in two sets of Ir 4f peaks that can be related to different oxidation states (figure 7.7b). The lowest observed value of Ir 4f5/2 is located at 61.5 eV, which is higher than the binding energy for metallic iridium, and smaller than the energies related to tetravalent Ir(IV) cations in IrO2. This intensity can therefore be associated with trivalent Ir(III) cations, which is the case in spinel ZnIr2O4. However, the largest portion (~75%) of intensity could still be ascribed to the presence of tetravalent iridium. This can probably be attributed to IrO2 or hydroxides normally present at the film surface. No etching to remove the top layer on these samples was used, as earlier observations showed that ion etching of the top layer resulted in decomposition of IrO2 [14].

Figure 7.7: XPS measurements of the 4f peaks of crystalline ZnIr2O4 deposited at 700 ºC, amorphous ZnO:x·IrO2 and amorphous IrO2 deposited at room temperature by the “eclipse” method(a).The peak profile of ZnIr2O4 can be deconvoluted into 4 peaks which can be attributed totwo different oxidation states (b)

72 70 68 66 64 62 60 58

Nor

m. i

nt. (

c/s)

Binding energy (eV)

76 74 72 70 68 66 64 62 60 58 56

Nor

m. i

nt. (

c/s)

ZnIr2O4

Ir 4f7/2Ir 4f5/2

Binding energy (eV)

ZnO:x·IrO2

IrO2

(a) (b)

Chapter 7 122

The XPS data of ZnIr2-δO4 shows that no intensity can be related to trivalent iridium. The spectrum of the material points out that it is similar to the chemical environment of IrO2. As a result this sample should not be considered as single phase, but merely as a composition of ZnO and IrO2. Therefore this compound should be indicated as ZnO:x·IrO2 instead of ZnIr2-δO4, in which x is the iridium zinc atomic ratio. Looking closer to figure 7.7a it might even be that the peaks of IrO2 and ZnO:x·IrO2 consist of both Ir(III) and Ir (IV). The peaks are rather broad, and the binding energy of Ir(III) in ZnIr2O4 is located in the tail of this intensity profile.

Additional information on structure and composition can be obtained from the specific binding energy of the oxygen. Amorphous IrO2 thin films are in general far from stoichiometric materials. These films mostly consist of mixtures of Ir, O, Ir(IV)O2, H2O and Ir(III)(OH)3 [16]. This can also be expected for the ZnO:x·IrO2 compound. From the O 1s XPS signal one can observe to which extent these mixtures are present, i.e. the IrO2 is hydrated [15]. The O 1s XPS intensity is build up from contributions of three different species, viz. lattice oxide (530 ± 0.5 eV), hydroxide (531.5 ± 0.5 eV), and water (533 ± 1 eV) [17]. These contributions are generally found in many transition metal oxides [18]. It can be observed that the O 1s spectra of the amorphous samples are rather broad (figure 7.8a and b). Water and a relative large amount of hydroxides are present in these films, which is responsible for the observation of trivalent iridium. The spectrum of spinel ZnIr2O4 mainly consists of lattice oxide. The hydroxide and water here, is likely to originate from the surface. These results also confirm that the ZnO:x·IrO2 film structure, as well as the composition is not related to spinel ZnIr2O4.

538 536 534 532 530 528 526

c/s

Binding energy (eV)

538 536 534 532 530 528 526

c/s

Binding energy (eV)

538 536 534 532 530 528 526

c/s

Binding energy (eV)

O

H2O

OH

(a) (b) (c)

Figure 7.8: XPS spectra of the O 1s oxygen level in IrO2 (a), ZnO:x·IrO2 (b) and ZnIr2O4 (c) thin films. The peaks are deconvoluted to lattice oxide O (~530 eV), hydroxide OH (~531.5 eV) andwater H2O (~533 eV).


7.3.4 Iridium valency Although the phase of ZnO:x·IrO2 could not be related to spinel, the

octahedral coordination of the oxygen ligands in IrO2 induces the same electronic band behavior as in spinel ZnIr2O4. In other words, because of the ligand field splitting, the band structure of IrO2 is also composed of t2g and eg states. However, because of the tetravalent Ir(IV), the t2g state is only occupied by five electrons, hence is not completely filled . The film is colored because of Ir d→d and O p→Ir d transitions [19,20]. To investigate this, the density of states of the valence levels is probed for Zn and IrO2 separately, and ZnO:x·IrO2. The data is presented in figure 7.9a, where the settings and conditions are similar to those used for the measurements presented in section 6.4.4.

The DOS of ZnO show that this material is truly n-type, the Fermi level is located far from the top of the valence band formed by the O 2p-levels. On the other hand, the profile of amorphous IrO2 shows that the onset of the valence levels is close to the Fermi level. This is different from observations by Kötz et al. who showed that the Fermi level moves into the valance band by about 1 eV, which is caused by the not fully occupied t2g

5 band in IrO2 [21].

12 10 8 6 4 2 0 -2 -4

ZnO

IrO2

t 62g

Zn 3d10

Nor

mal

ized

inte

nsity

(c/s

)

Binding energy (eV)

Ef

ZnO:x·IrO2

200 400 600 800 10000

10

20

30

40

50

%T

wavelength (nm)

T=700 °C

T=RT

Figure 7.9: XPS spectra of the valence states of ZnO, IrO2 and ZnO:IrO2 films grown at room temperature. The difference in transmission between as-deposited bleached Ir(III)(O/OH)3

and colored Ir(IV)O2 annealed at 700 °C (b), the XRD spectrum of both spectrum is also shown (c). The diffraction peaks are identified as IrO2 (110) and (101).

(a) (b)

10 20 30 40

T=RT

T=700 °C

c/s

2θ (°)

(c)

Chapter 7 124

IrO2 is known for its electrochromic properties; i.e. the material can change from colored to bleached, which is believed to be a result of oxidation from an electrolyte, chancing the valency of Ir(III) to Ir(IV) [22]. The as deposited amorphous IrO2 grown film described above is indeed hydrated as was observed from XPS measurements. In this case both valencies are frozen in by the presence of Ir(IV)O2 and Ir(III)(OH)3. This can be confirmed by the transmission of this IrO2 film. Figure 6.7b shows that this sample is somewhat transparent in the visible regime (~25%). Subsequently, the sample is annealed at 700 °C in air. The transmission has dropped drastically resulting in a blackened film. Figure 7.9c shows the XRD spectrum of this film before and after annealing. The as-deposited film is highly amorphous because of hydration. After annealing, most hydroxides are removed, and diffraction peaks related to IrO2 become visible, confirming that the Ir(IV) is responsible for coloration. Furthermore, the film is highly electronic conducting since the Fermi level moves in the valence band, resulting in metallic-like behavior [23,24]. The valence levels for Ir(IV) and Ir(III) are schematically drawn in figure 7.10a and b respectively, showing the difference in conduction behavior. It is obvious that only the “quasi-closed” d-shell leads to semiconducting behavior. The presence of trivalent Ir(III) in hydrated amorphous IrO2 is prerequisite for visible transparency and, when doped appropriately, p-type semiconducting behavior [25, 26].

stat

es

energy (eV)

O 2p

Ir t2g5

Ir eg0

EF

stat

es

O 2p

energy (eV)

Ir t2g6

Ir eg0

EF

(b) Ir(IV) Ir(III)

(a)

Figure 7.10: Schematic representation of the density of states of Ir(IV) (a) and Ir (III) (b) close to the Fermilevel. Grey areas denote occupied states, whereas white areas are unfilled. For Ir(IV) theFermi level is situated in the partly filled t2g

5 band, which acts as a conduction band. The film display highly “metal-like” conduction behavior. The bandgap is determined b the top of the O2p valence band and the Fermi level. For Ir(III) the t2g

6 levels are filled which now acts as a valence band. The Fermi level is located in between this band and the empty eg

0 conduction band, giving rise to semiconducting behavior.


7.4 Performance of amorphous films 7.4.1 Optical behavior

The transparency measurement of the ZnO:x·IrO2 films deposited on PET is presented in figure 7.11a (dashed curve). A high transmission is observed down to 300 nm. At this wavelength the light is absorbed by the PET substrate (see figure 3.2). In order to examine the optical properties over the complete, thin films are deposited on quartz substrates using similar conditions. The optical transmission of this sample is designated in figure 7.11a by a solid line. For comparison, the transmission curve of polycrystalline spinel ZnIr2O4 from chapter 6 is indicated in the graph as well. The measurements show that amorphous ZnO:x·IrO2 is transparent even in the UV-regime up to 50%. For decreasing wavelength the absorption starts around 450 nm but remains relatively weak down to 250 nm. The slope of absorption is even more fainted compared to ZnIr2O4.

The absorption versus the bandgap energy of the thin film on quartz is shown in figure 7.11b. From this graph two bandgap values were obtained; 2.2 eV and 5.5 eV respectively. However, especially for the high bandgap value, the determination by extrapolation is not very accurate. The absorbance line is not

1 2 3 4 5 6 70

2

4

6

8

10

12

α2 (x

1013

cm

-1)

E (eV)

200 400 600 800 10000

10

20

30

40

50

60

70

80

90

100

%T

wavelength (nm)

1 2 3 4 50.0

0.4

0.8

1.2

(a) (b)

Figure 7.11: Transmission spectrum of room temperature amorphous ZnO:x·IrO2 on PET (dashed line) and on quartz (solid line) substrate, as well as polycrystalline ZnIr2O4 deposited on quartz at 700 ºC (dotted line) (a). Absorption spectrum of ZnO:x·IrO2 on quartz for bandgap determination (b). The inset shows a magnification.

Chapter 7 126

straight. Furthermore, from figure 3.2 it can be seen that at this high energies the quartz substrate can somewhat influence the transmission.

No value was obtained that could be related to the spinel ZnIr2O4, ZnO or crystalline IrO2 phase; i.e. ~3.0 eV, ~3.2 eV, and ~3.5 eV respectively [26], Absorption starts from about 2.3 eV, and is relatively weak. This value is similar to the absorption energy obtained by Kamiya et al. on amorphous ZnO:Rh2O3 prepared in a 1:1 ratio [2]. This stoichiometry is close to that obtained in the ZnO:x·IrO2 (x≥ 1.3) described in this chapter. Apparently, the bandgap in amorphous films is lower with respect to the crystalline Rh(III) and Ir(III) containing material. The highly disordered structure can distort the octahedral coordination of these cations. The effective gap between the split d-band by the field of the ligands will therefore be lowered. Furthermore, in the case of amorphous hydrated ZnO:x·IrO2 some of the ligands consist of OH groups. From the spectrochemical series it appears that this ligand gives rise to a lower split compared to oxygen. In other words, the value of f in Jørgensen formula ∆=f•g is lower [27].

It is not exactly clear what causes the stronger absorption at around 5 eV. However, some authors suggest that this may be related to promotion of electrons from the deeper lying O 2p levels to the conduction band [19]. 7.4.2 Electrical conduction behavior Conductivity and Seebeck measurements are performed to determine the electrical properties of the amorphous films on PET. Also the dependency on the partial oxygen deposition pressure is investigated. The results are plotted in figure 7.12. The graph also shows the average transmission of the corresponding thin

0.09 0.1 0.2 0.3

1.5

2.0

2.5

3.0

3.5

55

60

65

70

75

cond

uctiv

ity (S

cm-1)

pO2 (mbar)

av.

Tr.

(%)

Figure 7.12: Conductivity () and average transmission versus the partial oxygen pressure () ofZnO:x·IrO2 thin films deposited on PET substrates.


films. The conductivity is found to vary between 1.5 Scm-1 and 3.3 Scm-1 over the oxygen pressure range investigated. The conduction type of the samples was determined by Seebeck measurements which confirmed p-type in all cases. Such high p-type conductivity is even difficult to obtain in crystalline TCOs as cuprates and p-ZnO [28,29]. Additionally, the conductivity shows some dependence on the oxygen pressure. The highest conductivity is obtained at pO2=0.250 mbar and is as high as 3.3 Scm-1. However, the best optical performance is not obtained in this sample. The transmission seems to be inversely related to the conductivity, possibly due to the increase of absorption on free carriers. The oxygen dependence implies that excess oxygen is one of the mechanisms for generation of positive hole carriers. On the other hand, acceptor doping is also possible. The film contains large amounts of zinc, but XRD results showed no presence of polycrystalline ZnO. It is plausible however, that the mixed state of the material significantly decreased the grainsize of ZnO similar to the observations in chapter 5 making them undetectable in diffraction experiments. But also amorphous ZnO absorbs light around 3.2 eV. And as stated earlier, no clear absorption features around this energy is observed. It can be concluded that at least some amount of the zinc is incorporated in the film in a different way. Excessive Zn2+ can therefore act as acceptors in the ZnO:x·IrO2 matrix. Similar to the spinel films in chapter 6, no meaningful Hall measurements could be performed. But the Seebeck coefficients of all samples are low (around +15 µV/K) indicating high hole carrier concentrations. From temperature dependant conductivity measurements, activation energy of 19 meV was obtained. Such a low value confirms that indeed high carrier concentrations can be obtained in this material.

7.4.3 Conduction paths Spinel ZnM2O4 thin films as presented in chapter 6, showed electrical properties that are independent on the crystallinity. This was ascribed to the edge sharing MO6-octahedra, forming conduction paths through the material [2]. In x-ray amorphous ZnRh2O4 deposited at room temperature, Narushima et al. [3,30] observed nanocrystals with dimensions of 2-3 nm by using TEM analysis. The high conductivity was explained from the fact that edge sharing octahedra have short M-M cationic distances. Together with the strong delocalization of the d-orbitals a large overlap can create conduction paths across grain boundaries, although the long range order is totally lost. In order to check the nanocrystallinity, TEM experiments with amorphous films were performedI. In figure 7.13a the TEM image of the film and the

I Thin films of 25 nm thickness were grown at room temperature on carbon covered

copper grids.

Chapter 7 128

corresponding diffraction pattern is presented. The high resolution image of the same sample is shown in figure 7.13b.

The layer presents a uniform distribution of very small grains. The diffraction pattern from the layer consists of several hazy rings which are formed by amorphous or to very small size crystal grains of this compound. The successive rings formed can be indexed as follows: 3.379 Å (110), 2.721 Å (101), 2.297 Å (200), 1.813 Å (211) and 1.674 Å (220). These values are close to IrO2, but measured distances do not correspond exactlyII. The obtained errors are around 10%. This can be an indication that the grains do indeed consist of octahedral structures, which are distorted by the highly disordered film, resulting in deviation of the structural distances.

From the high resolution images the existence of grains with size around 3 nm is proved. These grains are crystalline with a random orientation. Also from this structural analysis it was not possible to extract any information regarding the ZnO presence. These observations correspond to the result of earlier work on amorphous ZnO:Rh2O3. Here it was also mentioned that the excess Zn2+ ions act as network formers connecting the octahedron containing grains [2]

II JCPD standard (15-0870).

Figure 7.13: TEM image of the ZnO:x·IrO2 thin film and the corresponding diffraction pattern (a), and thehigh resolution image of the same sample (b)

(a) (b)

100 nm 5 nm


7.5 Transparent devices 7.5.1 P-N junctions

The discovery of p-type TCOs has opened up a new field in optoelectronics device technology [31] as now p-n junction could be fabricated, the building block in electronics. Since both n- and p-type materials are wide bandgap semiconductors, it became possible to design all-transparent devices [32,33] and blue or UV-light-emitting-diodes [34,35,36].

By using the material as p-layer in a p-n junction, the type of conductivity can be proved. The easiest way is to deposit the layer directly onto an n-type conducting substrate [37]. Amorphous ZnO:x·IrO2 thin films are deposited on n-type silicon at room temperature. A deposition mask partly covered the substrate defining the junction area. Ag contacts were applied on top of the film and on both sides of the substrate. A schematic representation of the junction is illustrated in figure 7.14a. The ohmic contact resistance is tested by measuring the I-V characteristic using both contacts on the substrate. A straight line confirmed the ohmic behavior. The same result was obtained for two contacts on top of an unstructured ZnO:x·IrO2 thin film. Subsequently, the I-V characteristic of the junction was measured. The results are shown in figure 7.14b.

(a)

Figure 7.14: Schematic representation of the p-ZnO:x·IrO2 (300 nm) / n-Si (575 µm) heterojunction, the junction area is 12 mm2 (a). I-V characteristic of the junction showing diode rectifying behavior (b).Schematic band diagram of the heterojunction under zero bias (left) and forwardbias (right) (c).

(c)

Ag contact

n- Si

ZnO:x·IrO2 Wiring

1.1

eV EF

EC

EV

~ 5

eV electrons

holes

~1 eV

P

N -4 -2 0 2 40.0

0.5

1.0

1.5

Cur

rent

den

sity

(mA/

cm2 )

Voltage (V)

(b)

Chapter 7 130

From the measurement it can be seen that diode rectifying behavior was obtained for this junction. The threshold voltage is around 1 eV, which is quite similar to the bandgap of silicon (1.1 eV). The bandgap energy is different for the n- and p-type section of the junction, so the barrier region is not symmetrical. If both sections are brought into contact, the bands will shift in order to equalize the Fermi level as can be seen from the schematic band diagramIII of figure 7.14c [38]. In forward bias, the holes can diffuse into the n-type section if the applied potential is equal or higher than the bandgap of silicon.

So the observed threshold voltage corresponds with the expectations for a p-n junction using silicon. Furthermore, a small leakage current is observed in the reverse bias region. The forward-reverse bias current ratio was about 80 in the - 4.0 to + 4.0 voltage range.

Although the diode rectifying behavior proves that amorphous ZnO:x·IrO2 is truly a p-type semiconductor, the diode performance is not optimal. At low applied forward bias voltages only the barrier for holes is eliminated. Only the holes can diffuse to the n-type region, since the barrier for electrons is still high. Although the hole concentration is high, their low mobility limits the diffusion, thus the diode performance. Figure 7.14c shows that if the bias voltage is increased sufficiently, the barrier for electrons will also disappear. Above this threshold voltage both holes

III In this schematic representation the band bending effects are ignored.

Figure 7.15: The I-V characteristic of the junction on silicon substrate at higher bias current. The inset shows the logarithmic graph to illustrate the diode forward-reverse bias current ratio.

-10 -8 -6 -4 -2 0 2 4 6 8 10

10

20

30

40

Cur

rent

Den

sity

(mA/

cm2 )

Voltage (V)

-8 -4 0 4 8

0.01

0.1

1

10


and electrons will diffuse to the neighboring, resulting in higher junction current. This effect becomes visible as the applied bias voltage range of the same junction is increased. The corresponding diode characteristic is presented in figure 7.15. After the current increases around I V due to the bandgap of silicon, another sharp increase is observed starting around 6 V. This value can be related to the sharp absorption in figure 7.11, corresponding to a large bandgap. Apparently, the earlier observed bandgap around 2.3 eV has negligible effect compared to that of 5 eV. The logarithmic graph shows that the bias current ratio has significantly increased at higher applied voltages.

7.5.2 Transparent electronics on plastic

The use of a wide bandgap n-type semiconductor instead of silicon not only results in transparent devices [39], but the bandgap of the materials is tuned. The junction is more symmetric, and the threshold voltage corresponds to the lowest bandgap value. Furthermore, wide bandgap materials should be utilized if the junction should operate as an LED. The emitted light corresponds to the bandgap of the material in which electron-hole pair recombination takes place.

Many reports on the fabrication of p-n junctions consisting of wide bandgap crystalline oxides on heat resistant substrates can be found in literature (references in section 6.1 and 7.5.1). On the other hand the realization of junctions consisting of all-TCOs on polymers is only limited to p-ZnO:Rh2O3 [3]. The authors demonstrated diode rectifying behavior in a device using this material on flexible plastic substrate. From the results presented in this chapter it is likely that the same could be achieved by using p-ZnO:x·IrO2.

In order to fabricate the p-n junction, several multilayer stacks are deposited on PET substrates (250 µm Melinex®) at room temperature. Except for the top electrode, all subsequent layers are deposited in-situ by changing the target material. The bottom electrode is 200 nm ITO 5%, deposited at optimum conditions found in chapter 4. The 250 nm thick n-ZnO layer is deposited at oxygen pressure higher than optimum conditions (pO2=0.040 mbar) to establish semiconducting behavior. The p-ZnO:x·IrO2 is 200 nm thick and is deposited at pO2=0.250 mbar. After completion of this layer, the sample was taken out of the PLD system and the 150 nm gold top-electrode was sputtered.

On top of the multilayer, squared patterns of photo-resist were created on the surface by a standard photolithographic process. After determination of the sputter rate, the sample was argon etched down to the bottom electrode as schematically drawn in figure 7.16a. After removal of the remaining photo-resist, several junctions with different area dimensions are left. A top view photograph of the sample showing some junctions is also shown in figure 7.16a. The squares are the gold electrodes, whereas the grey area in between consist of only ITO on PET.

Chapter 7 132

The dimensions of the junctions range from 500x500 µm2 down to 5x5 µm2. In order to measure the junction characteristics, contact was made to bottom and top electrodes by using a probe station. It appeared that several junctions were cracked. Probably the heating step (1 minute at 90 ºC) in the lithographic process caused thermal stress at the polymer–thin-film interface. In addition, some junctions that seemed to be intact did not show rectifying behavior. Pinholes, shorts on the side of the junction or other defects are able to negatively affect the diode characteristics. Only few junctions showed I-V characteristics than can be associated with p-n interfaces. The best obtained result is presented in figure 7.16b, but still this forward bias current is considerably low. Although the reverse bias current is relatively large, diode rectifying behavior can be observed. The threshold voltage is higher than for junctions on silicon. A small increase is observed around 4 V, which is likely to be related to the bandgap of ZnO (3.2 eV). However, also in this junction the sharp increase starts around 6 V.

From the above results one can conclude that the fabrication of the p-n junction on PET is far from optimal. This results in decreased performance.

PET

ITO

ZnO

Au

ZnO:x·IrO2

-10 -8 -6 -4 -2 0 2 4 6 8 10

0.5

1.0

1.5

2.0

2.5

Cur

rent

(x10

-6 A

)

Voltage (V)

(b) (a)

Figure 7.16: Schematic representation of the stacked layered heterojunction on PET substrate. Thesubsequent layers are ITO electrode (200 nm), n-ZnO (250 nm), p-ZnO:x·IrO2 (200 nm) and gold top electrode (150 nm). The photograph below shows a microscope image of the top view structured junctions. The largest area is 500x500 µm2 (a). Typical I-V characteristic of the heterojunction on PET (c).


Furthermore, the reproducibility of manufacturing these devices with the process used here is low. However, the results also show that besides p-ZnO:Rh2O3, it is also possible to create devices with p-ZnO:x·IrO2 on heat sensitive substrates, since these two materials posses the unique property of p-type conductivity in an (nearly) amorphous structure. Diode rectifying behavior of a room temperature deposited all-amorphous oxide multilayer could therefore be demonstrated. Although much improvement on the fabrication process of the devices on polymers is required, the method used here is sufficient for a proof-of-principle. 7.6 Concluding remarks

In general, p-type TCOs are crystalline materials requiring high process temperatures. These materials are not suitable for transparent electronic applications on flexible polymeric substrates. ZnO:x·IrO2 however, is a transparent material that possesses p-type conductivity in a nanocrystalline phase if grown at room temperature. The observed grainsize in this material is only 3 nm.

The material is deposited by the PLD “eclipse method” to reduce iridium droplet formation on the surface. This method results in a smooth surface, but also non-stoichiometric films. For ZnO:x·IrO2, a value ranging from 1.2 to 1.8 seems not to influence the electrical and optical properties significantly.

Similar to ZnIr2O4, the bandgap in ZnO:x·IrO2 is induced by the ligand field splitting of the iridium d-levels. From experiments, no traces of the spinel ZnIr2O4 phase were found. The IrO6-octahedra however, are still present in the material, forming a network within the thin film. The valency of the majority of iridium in the hydrated material is found to be threefold. This gives rise to transparency and p-type semiconducting behavior when doped appropriately, because of the “quasi-closed” d-shell. Depending on the oxygen pressure the average optical transmission is as high as 70%, whereas the p-type conductivity can be above 3 Scm-1. Most light is absorbed at energies higher than 5 eV, making the material reasonable transparent in the UV-regime.

Amorphous ZnO:x·IrO2 is used as p-layer in junction devices. Rectifying diode behavior was observed. Furthermore, the material is used in junctions consisting of all-amorphous semiconductors on PET substrates. These experiments demonstrated that ZnO:x·IrO2 is a serious candidate for applications as LEDs or (transparent) electronics in the near future.

Chapter 7 134

7.7 References 1. H. Mizoguchi, M. Hirano, S. Fujitsu, T. Takeuchi, K. Ueda & H. Hosono, Appl.

Phys. Lett. 80, 1207 (2002). 2. T. Kamiya, S. Narushima, H. Mizoguchi, K.-I. Shimizu, K. Ueda, H. Ohta, M.

Hirano & H. Hosono, Adv. Funct. Mater. 15, 968 (2005). 3. S. Narushima, H. Mizoguchi, K.-I. Shimizu, K. Ueda, H. Ohta, M. Hirano, T.

Kamiya & H. Hosono, Adv. Mater. 15, 1409 (2003). 4. H. Ohta, H. Mizoguchi, M. Hirano, S. Narushima, T. Kamiya & H. Hosono,

Appl. Phys. Lett. 82, 823 (2003). 5. P.M. te Riele, in Growth of nanoparticles in the pulsed laser deposition

process, Master thesis, University of Twente, The Netherlands (2004). 6. M.S. Tillack, D.W. Blair & S.S. Harilal, Nanotechnology 15, 390 (2004). 7. P.R. Willmott & J.R. Huber, Rev. Mod. Phys. 72, 315 (2000). 8. K. Kinoshita, H. Ishibashi & T. Kobayashi, Jpn. J. Appl. Phys. 33, L417

(1994). 9. D.B. Chrisey & G.K. Hubler, In Pulsed Laser Deposition of Thin Films, John

Wiley & Sons Inc., New York (1994) 10. M.D. Strikovsky, E.B. Klyuenkov, S.V. Gaponov, J. Schubert & A. Copetti,

Appl. Phys. Lett. 63, 1146 (1993). 11. Z. Trajanovic, S. Choopun, R. P. Sharma & T. Venkatesan, Appl. Phys. Lett.

70, 3461 (1997). 12. A. Marcu, C. Grigoriu, W. Jiang & K. Yatsui, Thin Solid Films 360, 166 (2000). 13. C. Wagner, W. Riggs, L. Davis, & J. Moulder, in Handbook of x-ray

photoelectron spectroscopy, edited by G.E. Muilenberg, Perkin-Elmer Corporation, Eden Prairie, Minnesota, (1979).

14. R.H. Horng, D.S. Wuu, L.H. Wu & M.K. Lee, Thin Solid Films 373, 231 (2000).

15. S. Yao, M. Wang & M. Madou, J. Electrochem. Soc. 148, H29 (2001). 16. P.S. Patil, P.S. Chigare, S.B. Sadale, T. Seth, D.P. Amalnerkar & R.K. Kawar,

Mater. Chem. Phys. 80, 667 (2003). 17. R. Sanjinés, A. Aruchamy & F. Lévy, J. Electrochem. Soc. 136, 1740 (1989) 18. M. Peuckert, J. Electroanal. Chem. 185, 379 (1985). 19. A.K. Goel, G. Skorinko & F.H. Pollak, Phys. Rev. B 24, 7342 (1981). 20. J.D.E. McIntyre, W.F. Peck & S. Nakahara, J. Electrochem. Soc. 127, 1264

(1980). 21. E.R. Kötz & H. Neff, Surf. Sci. 160, 517 (1985). 22. S. Hackwood, A.H. Dayem & G. Beni, Phys. Rev. B 26, 471 (1982). 23. S. Gottesfeld, J. Electrochem. Soc. 127, 1922 (1980). 24. R. Sanjinés, A. Aruchamy & F. Lévy, Solid State Commun. 64, 645 (1987).


25. J. Kukkonen, O. Försen, J. Aromnaa & S. Ylasaari, Mat. Sci. Forum 192-194,

213 (1995). 26. T.M. Silva, A.M.P. Simões, M.G.S. Ferreira, M. Walls & M. Da Cunha Belo, J.

Electroanal. Chem. 441, 5 (1998). 27. C.K. Jørgensen, in Modern aspects of ligand field theory, 1st edition, North-

Holland Publishing Company, Amsterdam, Netherlands (1971), Ch. 26. 28. H. Gong, Y. Wang & Y. Luo, Appl. Phys. Lett. 76, 3959 (2000). 29. K.-K. Kim, H.-S. Kim, D.-K. Hwang, J.-H. Lim & S.-J. Park, Appl. Phys. Lett,

83, 63 (2003). 30. S. Narushima, H. Mizoguchi, H. Ohta, M. Hirano, K.-I. Shimizu, K. Ueda, T.

Kamiya, & H. Hosono, Mat. Res. Soc. Symp. Proc. 747, V2.2.1 (2003). 31. A.N. Banerjee & K.K. Chattopadhyay, Prog. Cryst. Growth Charact. Mater.

50, 52 (2005). 32. G. Thomas, Nature 389, 907 (1997). 33. J.F. Wager, Science 300, 1245 (2003). 34. H. Ohta, K.-I. Kawamura, M. Orita, M. Hirano, N. Sarukura, & H. Hosono,


239, 496 (2002). 37. P. Wang, N.F. Chen & Z.G. Yin, Appl. Phys. Lett. 88, 152102 (2006). 38. H. Ibach & H. Luth, In Solid-Stae Physics, An introduction tp principles of

Materials Science, chapter 12, 2nd edition, Springer, Berlin (1996). 39. A. Kudo, H. Yanagi, K. Ueda, H. Hosono, H. Kawazoe & Y. Yano, Appl. Phys.

Lett. 75, 2815 (1999).

Chapter 7 136

Summary

This thesis describes the research on thin films of transparent conducting oxides (TCOs) on polymeric substrates manufactured by pulsed laser deposition (PLD). TCOs are an indispensable part in optoelectronic applications such as displays, solar cells, light-emitting diodes, etc. At present, in many of these applications there is an increasing need for the use of flexible, cheap and light-weight substrates. Such polymer substrates however, limit the deposition temperature of thin films which results in deteriorated properties of the TCO. A profound understanding of the fundamental aspects of transparent conductors is required in order to improve either the properties of existing materials, or design new types of TCOs. These insights are of great scientific importance for the realization of high performance TCOs on polymer substrates.

This research focuses on thin film growth by PLD. This technique is a powerful tool for thin film research. A large freedom of choice in independently controllable deposition parameters allows one to quickly obtain results on the exploration and optimization of existing and new materials. The ablated species can be tuned over a large energetic range, enabling optimum conditions at lower substrate temperatures normally used for high performance TCO materials. This makes the deposition of TCOs on heat resistive substrates possible. The substrates utilized in this research are all commercial available and commonly used materials such as translucent polyethylene terephthalate (PET). Analysis of the PLD grown films is done by a variety of tools to obtain information on the electrical, optical and structural properties as well as the thin film composition.

Different TCO materials for deposition on polymers are investigated:

SnO2-doped In2O3 (ITO), pure and Al2O3-doped ZnO (AZO) and the In2O3-ZnO compound (IZO). These materials are all n-type wide bandgap semiconductors (Eg>3 eV). Their electrical and optical performance is in a large extent determined

Summary 138

by fundamental properties as structure and composition, which can be tuned and manipulated in the deposition process.

ITO is a commonly used TCO in optoelectronic applications. In this work the deposition process of ITO on polymeric substrates is optimized. By careful control of the PLD parameters, films with high optical transmittance (>85%) and low resistivity (~5 x10-4 Ωcm) are grown at room temperature on PET substrates. The influence of Sn-doping in In2O3 thin films on conductance, transmission, and granular structure is studied. It is found that the Sn-dopant is not thermally activated and does not contribute to enhanced conductivity in room temperature deposited films. The electrical properties in these ITO films are therefore governed by the oxygen vacancies, resulting in a strong dependence on the oxygen deposition pressure. Only in a narrow deposition pressure regime (0.012-0.017 mbar) high quality ITO thin films on polymer substrates can be grown. Somewhat higher pressure, in the order of only a few thousandth of a millibar compared to the optimum, results in decreased amount of oxygen vacancies, hence lower carrier density. The resistivity increases and the bandgap narrows due to the Burstein-Moss shift. On the other hand, a lower pressure results in higher carrier density. However, the increased ionized impurity scattering lowers the free carrier mobility, also causing a higher resistivity.

The resistivity of room temperature grown thin films ablated from Sn-doped targets is higher compared to samples of pure In2O3. A relation between the structural nature of the thin film and the amount of Sn doping is observed. The size of the formed grains during growth decreases at higher doping content. The obtained grain sizes are in the order of the mean free path of the charge carriers. The grain boundaries of these nanocrystalline films influence their electrical properties. Due to increased scattering, the charge carrier mobility is lowered. Sn-doped films with a higher grain boundary density compared to undoped samples are therefore more resistive.

TCOs composed of (doped) ZnO and Zn containing compounds are of major interest, since these materials are regarded as replacement for ITO. In this work undoped ZnO is deposited onto PET substrates. In order to grow smooth films of ZnO by PLD on polymers, the process parameters are different to those used for ITO deposition. The pressure is increased (0.050 mbar), whereas the fluency is lowered (1.5 J/cm2). This prevents excessive substrate heating by impingement of high kinetic energy of ablated Zn-species. Otherwise, it is this heating and the large difference in expansion coefficient between film and substrate that causes cracking of the thin film.

Similar to ITO, the properties of ZnO also showed to be strongly dependant on the oxygen partial deposition pressure. Optimized films show excellent optical properties (T~90%). The electrical resistivity is still a factor of ten higher compared

Summary 139

to ITO thin films on PET. A small amount of Al2O3 doping (AZO) enhances the electrical properties of the room temperature grown films. However, analogous to SnO2 in In2O3, addition of Al2O3 in ZnO decreases the grain size. This increases the grain boundary scattering in AZO films with higher doping content. The optimum Al2O3 doping content in room temperature deposited films on PET is therefore found at 1 %wt.

The benefits of both ZnO and In2O3 are combined in one compound system. The properties of this material are dependant on the In2O3 to ZnO ratio. The optical transmission of films from the ZnO-In2O3 compound system on PET substrates is high and constant over the whole composition range from pure In2O3 to pure ZnO. The electrical resistivity can be tuned by the ZnO content as it increases one order of magnitude on going from In2O3 to ZnO.

In general n-type TCOs are used as electrodes, but in combination with

p-type TCOs active components can be realized. After all, the p-n junction is an essential building block in all semiconducting electronics. The lack of appropriate p-type wide bandgap semiconductors however, hampers the development of new optoelectronic devices based on polymeric substrates. Since the use of polymer substrates requires low temperature deposition, this research focuses on materials that are less dependent on the crystalline structure compared to existing p-type TCOs. These experiments resulted in the synthesis of a new p-type TCO; ZnIr2O4. This material is also suitable for deposition at room temperature on polymeric substrates.

The closed d10-shell in existing p-type TCOs as cuprates, avoid coloration of the material. However, transparency can also be achieved by using d6-transition-metal cations in octahedral geometry surrounded by oxygen ligands. A “quasi-closed” shell is induced since the ligands cause a splitting of the d-bands in a completely filled t2g

6 and empty eg0 level. The bandgap of the material is determined

by the energy split in between these levels. This phenomenon is observed in spinel ZnM2O4 thin films, where element M is a d6- transition metal.

ZnM2O4 thin films, where M=Co, Rh and Ir, are pulsed laser deposited on quartz and crystalline Al2O3 (0001) substrates at temperatures between 773 and 973 K. The material is ablated from targets obtained by solid state synthesis from binary oxide powders. The crystal structure is found to be spinel of both polycrystalline films on quartz and epitaxial films on Al2O3. The electronic band structure of the grown materials reveals that the valence band maximum is composed of occupied t2g

6 states. The observed bandgap is increasing for higher quantum number of element M, being as large as ~3 eV for ZnIr2O4. This increasing bandgap is expected from ligand field theory and scales with the theoretical predictions. P-type conductivity is confirmed in all compounds by positive Seebeck

Summary 140

coefficients and the position of the Fermi level with respect to the valence level. The conductivity of ZnIr2O4 is well above 2 Scm-1, whereas the optical transmission is around 60%.

Room temperature deposition of ZnIr2O4 on PET substrates results in amorphous films. A special deposition method referred to as “eclipse PLD” is used to avoid droplet formation on the film surface. The material cannot be identified as single phase spinel, but merely as a non-stoichiometric ZnO:x·IrO2 compound, where 1.2<x<1.7. However, the film is hydrated and the majority of iridium in the film appears in a trivalent valency. Nanocrystals of about a few nanometers containing trivalent iridium in octahedral surrounding are dispersed in the Zn2+-matrix. The network of octahedrons forms a conduction path for the positive holes generated by excess oxygen and cationic vacancies. The p-type conductivity of these films deposited on PET substrates is as large as 2 Scm-1. The film is transparent up to 70% in the visible, and a 50% transparency is observed in the UV regime.

The film is used as p-layer in a junction device based on silicon. Diode rectifying behavior is observed and is in accordance with the different bandgap values of the separate p- and n-layer. A device deposited on a polymer substrate using ZnO as n-layer is also tested. Although the reproducibility of this device is low due to the critical manufacturing process, diode rectifying behavior is also observed in these junctions. This demonstrates the applicability of this material in devices containing all-amorphous wide bandgap semiconductors on polymeric substrates.

Depositions at low temperature for applications on polymers involve

nanocrystalline and amorphous materials. The research topics in this thesis demonstrate that the understanding of the fundamental properties of wide bandgap semiconductors with such structures is essential. This knowledge is employed for the optimization of the deposition process, leading to the improvement of polymer-TCO systems. Comprehension of the fundamental aspects down to the atomic level resulted in design and synthesis of a new p-type TCO. The performance of this TCO is high in crystalline as well as the amorphous phase. The obtained results can contribute to the realization of transparent electronics on heat resistive substrates in the near future.

Samenvatting

Dit proefschrift beschrijft het onderzoek aan dunne lagen van transparant geleidende oxides (TCO’s) gemaakt met behulp van gepulste laser depositie (PLD) op polymere substraten. TCO’s maken een onmisbaar onderdeel uit van opto-elektronische toepassingen zoals displays, zonnecellen, LED’s, enz. In veel van deze toepassingen bestaat tegenwoordig een toenemende behoefte aan het gebruik van flexibele, lichtgewicht en goedkope substraten. Zulke polymere substraten stellen echter een limiet aan de te gebruiken depositietemperatuur, wat een verslechtering van de eigenschappen van de TCO’s tot gevolg heeft. Een diepgaand begrip van de fundamentele aspecten van TCO’s is daarom nodig om eigenschappen van bestaande materialen te verbeteren, en om nieuwe te kunnen synthetiseren. Zulke bevindingen zijn van groot wetenschappelijk belang om hoogwaardige TCO’s op polymere substraten te realiseren.

Dit onderzoek richt zich op het groeien van dunne films met behulp van PLD. Deze techniek is uitermate geschikt voor het doen van onderzoek naar dunne lagen. Vele depositieparameters zijn onafhankelijk van elkaar te controleren zodat nieuwe en bestaande materialen relatief snel onderzocht en geoptimaliseerd kunnen worden. De kinetische energie van de geableerde deeltjes is nauwkeurig te controleren. De optimale depositietemperatuur van TCO’s kan hierdoor verlaagd worden, waardoor het mogelijk wordt om substraten te gebruiken die niet hittebestendig zijn. De in dit onderzoek gebruikte substraten zijn gangbare en commercieel verkrijgbare materialen, zoals transparant polyethyleentereftalaat (PET). De elektrische, optische, structurele en compositionele eigenschappen van de gedeponeerde dunne films zijn onderzocht met behulp van verschillende analyse technieken.

Verschillende transparant geleidende materialen op polymere substraten

zijn onderzocht: SnO2 gedoteerd In2O3 (ITO), puur en Al2O3 gedoteerd ZnO (AZO) en de samenstelling In2O3-ZnO (IZO). Dit zijn alle n-type halfgeleiders met een grote waarde van de bandgap (Eg >3 eV). De elektrische en optische kenmerken

Samenvatting 142

worden in hoge mate bepaald door fundamentele eigenschappen als kristalstructuur en compositie, welke manipuleerbaar zijn in het depositieproces.

ITO is de meest gebruikte TCO in opto-elektronische toepassingen. In dit werk is het depositieproces van ITO op polymere substraten geoptimaliseerd. Zorgvuldige regulering van de PLD parameters resulteert in dunne films met hoge optische transmissie (>85%) en lage resistiviteit (~5 x10-4 Ωcm) op PET substraten. De invloed van Sn-dotering in In2O3 op de geleidbaarheid, transmissie en granulaire structuur is onderzocht. Het blijkt dat Sn niet thermisch geactiveerd is en dus niet bijdraagt aan de elektrische geleiding in deze ITO films gegroeid op kamertemperatuur. De elektrische eigenschappen zijn daarom vooral bepaald door de zuurstofvacatures in deze dunne films, en zijn dus sterk afhankelijk van de zuurstofdruk tijdens depositie. Enkel in een smal drukregime (0.012 – 0.017 mbar) kunnen films met zeer goede eigenschappen op polymeren worden gegroeid. Een paar duizendste millibar hogere druk dan het optimum vermindert drastisch het aantal zuurstofvacatures, en dus ook de ladingsdrager dichtheid. De resistiviteit neemt toe en de bandgap wordt kleiner ten gevolge van de Burstein-Moss verschuiving. Aan de andere kant resulteert een lagere druk in meer ladingsdragers. De toegenomen ladingsverstrooiing verlaagt echter de mobiliteit van elektronen, dat ook een hogere resistiviteit tot gevolg heeft.

Op kamertemperatuur gegroeide dunne films afkomstig van Sn-gedoteerde targets hebben een hogere weerstand dan films van pure In2O3 targets. Een verband kan dan ook worden waargenomen tussen de structuur en Sn-dotering van de dunne film. Naarmate de Sn-dotering toeneemt, neemt de korrelgrootte in de groeiende film af. De korrelgroottes zijn van dezelfde orde van grootte als de vrije weglengte van de ladingsdragers. De elektrische eigenschappen van deze nano-kristallijne films worden daarom mede bepaald door de invloed van korrelgrenzen. Door verstrooiing aan deze korrelgrenzen daalt de mobiliteit van de elektronen. Sn-gedoteerde films hebben een hogere dichtheid aan korrelgrenzen, en dus ook een hogere weerstand.

TCO’s die bestaan uit (gedoteerd) ZnO of zinkhoudende verbindingen staan wetenschappelijk erg in de belangstelling. Deze materialen worden gezien als vervangers van ITO in de nabije toekomst. In dit onderzoek zijn ongedoteerde ZnO dunne lagen gegroeid op PET substraten. De depositieparameters verschillen sterk ten opzichte van ITO depositie. Films die niet barsten worden alleen verkregen indien een hogere druk (0.050 mbar) en lagere energiedichtheid van de laserspot (1.5 J/cm2) wordt gebruikt. Dit voorkomt overmatige opwarming van het substraat door het bombardement van hoog energetische Zn deeltjes. Deze opwarming en het grote verschil in expansiecoëfficiënt tussen film en substraat resulteert namelijk in het barsten van de dunne laag.

Net als ITO, zijn de eigenschappen van ZnO sterk afhankelijk van de partiële zuurstofdruk tijdens depositie. Geoptimaliseerde films vertonen zeer goede optische eigenschappen (T~90%); de resistiviteit is echter een factor tien hoger dan ITO films op PET. Dotering van ZnO met kleine hoeveelheden Al2O3 verbetert de

Samenvatting 143

elektrische eigenschappen van de films gegroeid op kamertemperatuur. Echter, net als SnO2 in In2O3, vermindert toevoeging van Al2O3 in ZnO de korrelgrootte van de dunne laag. Verstrooiing aan korrelgrenzen is daarom van grotere invloed in AZO films met hogere Al2O3 concentratie. De optimale Al2O3 concentratie in ZnO films gedeponeerd bij kamertemperatuur op PET is 1 %wt.

De goede eigenschappen van ZnO en In2O3 zijn gecombineerd in een samengesteld systeem. De eigenschappen van dit materiaal zijn afhankelijk van de verhouding In2O3 en ZnO. De optische transmissie van deze samenstelling op PET substraten is hoog en constant over de gehele range variërend van puur In2O3 tot puur ZnO. De resistiviteit kan tot een orde van grootte worden veranderd door de hoeveelheid ZnO te variëren. Deze waarde daalt als de compositie verandert van puur ZnO naar puur In2O3.

Over het algemeen worden n-type TCO’s gebruikt als elektrodes. In

combinatie met p-type TCO’s, kunnen daarentegen actieve componenten worden gerealiseerd. De p-n junctie is immers een essentiële bouwsteen in alle halfgeleider elektronica. Het gebrek aan geschikte p-type materialen met grote bandgap stagneert de ontwikkeling van opto-elektronische toepassingen op polymere substraten. Omdat het gebruik van polymere substraten gepaard gaat met lage depositietemperaturen, richt dit onderzoek zich voornamelijk op materialen waarvan de eigenschappen minder afhankelijk zijn van de kristalstructuur. Deze experimenten resulteerden in de synthese van een nieuw p-type TCO; ZnIr2O4. Dit materiaal is geschikt voor depositie op polymere substraten.

De gesloten d10-schil in de meeste bestaande p-type TCO’s (zoals de koperverbindingen) voorkomt kleuring van het materiaal. Transparantie kan echter ook worden gecreëerd door d6-transitiemetalen omgeven door zuurstofliganden in octaëdrische geometrie te gebruiken. De liganden veroorzaken een splitsing van de d-banden in een gevulde t2g

6 en lege eg0 orbitalen, zodat een “quasi-gesloten” schil

ontstaat. De bandgap bevindt zich tussen deze energieniveaus en de grootte wordt bepaald door de mate van de splitsing. Dit fenomeen doet zich voor in spinel ZnM2O4 dunne films, waarin M een d6-transitiemetaal is.

ZnM2O4 dunne lagen, met M =Co, Rh en Ir, zijn met PLD gegroeid op kwartsglas en kristallijne Al2O3 (0001) substraten op een temperatuur tussen 773 en 973 K. De PLD targets zijn verkregen door vaste-stof synthese met poeders van de binaire oxiden. De gevormde poly-kristallijne en epitaxiale kristalstructuur van de films op respectievelijk kwarts en Al2O3 is in beide gevallen spinel. De elektronische bandenstructuur van de materialen bevestigt dat de valentie band inderdaad bestaat uit gevulde t2g

6 orbitalen. De waargenomen bandgap wordt groter naarmate het kwantumgetal van element M toeneemt. De grootste bandgap van ongeveer 3 eV is daarom waargenomen in ZnIr2O4. De gevonden toename van de bandgap schaalt met de theoretische voorspellingen die volgen uit de ‘ligand-veld theorie’. P-type geleiding is bevestigd in alle drie de materiaalsystemen door de positieve Seebeck- coëfficiënt en de positie van het Fermi-niveau ten opzichte van de

Samenvatting 144

valentie band. De geleiding van ZnIr2O4 is hoger dan 2 Scm-1 terwijl de gemeten transmissie rond de 60% is.

ZnIr2O4 gedeponeerd bij kamertemperatuur op PET resulteert in amorfe films. Een speciale depositiemethode, genaamd “eclips PLD”, is toegepast om de deeltjes die ontstaan op het oppervlak van de dunne laag te verminderen. In plaats van een stoichiometrische spinel fase bestaat het materiaal uit een ZnO:x·IrO2 samenstelling, waarbij 1.2<x<1.7. Echter, de laag is gehydrateerd en het iridium ion is driewaardig. Nano-kristallen (2-3 nm) bevatten dit driewaardige iridium ion in octaëdrische omringing in een Zn2+-matrix. Het netwerk van octaëders vormt een geleidingspad voor de positieve elektrongaten die zijn ontstaan door een zuurstofoverschot en vacante kationposities. De p-type geleiding van deze films gedeponeerd op PET is 2 Scm-1. De film is 70% transparant in het optische, en 50% in het UV regime van het spectrum.

Een op kamertemperatuur gegroeide ZnIr2O4 film is gebruikt als p-laag in een junctie door depositie op n-type silicium. Diode-gedrag is geconstateerd en is in overeenstemming met de waarden van de bandgap van de p- en n-laag. Een junctie gebruikmakend van n-type ZnO op een polymeer substraat is tevens getest. Hoewel door het productieproces de reproduceerbaarheid van deze junctie laag is, kan ook hier diode-gedrag geconstateerd worden. Dit demonstreert de toepasbaarheid van ZnO:x·IrO2 in applicaties bestaande uit geheel amorfe transparante halfgeleiders op polymeren.

Depositie bij lage temperaturen voor toepassingen op polymere substraten

resulteert veelal in nano-kristallijne en amorfe materialen. De onderzoeksthema’s in dit proefschrift tonen aan dat het begrip van de fundamentele eigenschappen van zulke transparante halfgeleiders van groot belang is. Deze kennis draagt namelijk bij aan de optimalisatie van het depositieproces wat resulteert in verbeterde polymeer-TCO systemen. Kennis van de fundamentele aspecten tot op atomair niveau heeft geleid tot het ontwerp en synthese van een nieuw p-type TCO. De eigenschappen van deze TCO zijn goed in zowel kristallijne als amorfe fase. De gevonden resultaten kunnen bijdragen aan de realisatie van transparante elektronica op polymere substraten in de nabije toekomst.

Dankwoord

Eigenlijk is dit de beste plek om een samenvatting van mijn werk te plaatsen, aangezien het dankwoord het meest gelezen onderdeel van een proefschrift schijnt te zijn. Dit maakt het voor mij dan ook het moeilijkste stuk om te schrijven. Een foutje in de eerste 144 pagina’s blijft (hopelijk) onopgemerkt, terwijl iemand vergeten te bedanken op deze laatste twee pagina’s desastreus is. De afgelopen jaren van onderzoek met dit boekje als gevolg zouden immers nooit mogelijk zijn geweest zonder de hulp, steun en aanwezigheid van velen.

Op de eerste plaats wil ik daarom Dave bedanken. Enthousiasme is niet voor niets het woord dat altijd met jou geassocieerd wordt. Door deze enthousiaste stimulans heb jij ervoor gezorgd dat ik het traject van mijn studie, afstudeerproject tot en met promotie als vanzelfsprekend heb ervaren. De vele ideeën die je aandroeg hebben er voor gezorgd dat ik geen moment mijn interesse in het onderzoek heb verloren. Daarnaast was er buiten werktijd altijd plaats voor gezelligheid tijdens etentjes of conferentiebezoek. Verder ben ik blij dat we onze belangstelling voor de materiaalkunde en de Formule 1 wederzijds hebben kunnen maken.

Natuurlijk is de input van Guus de afgelopen jaren onmisbaar geweest. Ondanks grote drukte blijf jij door de bomen het bos zien, en hebt bijgedragen aan ontelbare suggesties en sturing van het onderzoek. De overdracht van zowel kennis als informatie gaat jou even gemakkelijk af. En…, de ‘natte-vinger benadering’ kan uiteindelijk toch zeker leiden tot succes. De wetenschappelijke, maar ook de vele off-topic discussies tijdens lunch, koffie of borrel werden zeer gewaardeerd. Het heeft mij geleerd om alles altijd kritisch te blijven bekijken.

Alle leden van de begeleidingscommissie wil ik bedanken voor de belangrijke feedback in en om de meetings. De gelegenheid om ook eens in “andermans keuken te kijken” was uitermate interessant.

Dankwoord 146

Dick, Frank en Henk zorgden voor alle technische ondersteuning en uitleg die noodzakelijk zijn voor het bedrijven van wetenschappelijk onderzoek. Jullie laten je nooit uit het veld slaan door de hordes studenten en AIO’s die onophoudelijk vragen stellen en dingen stuk maken. Sybolt en Gerrit hebben toch een beetje een kristallograaf van mij gemaakt, iets om trots op te zijn. Ard, jouw pionierswerk naar p-type TCOs is mij erg behulpzaam geweest. Dankzij Evert kon ik nog wat last-minute juncties meten, en Beatriz hielp met de TEM experimenten. Ook de discussies met alle leden van de technische staf van de vakgroep hebben een constructieve bijdrage aan mijn onderzoek geleverd. Marion en José stonden garant voor hulp in administratie en IT.

Met mijn kantoorgenoten heb ik buitengewoon gezellige tijden beleefd. Tussen de werkzaamheden door was er altijd wel even tijd voor ontspanning. Ook buiten het kantoor was de sfeer altijd opperbest. Frank!, alias Frans, ik vond Veldhoven goed, Nice beter, maar Kreta was allerbest. Met Paul heb ik waarschijnlijk wel de meeste bierella’s, het smeermiddel van de wetenschap, genuttigd. Mijn vertrouwen in de Veen gaat zelfs zover dat ik hem tot de nieuwe captain van het voetbalteam heb benoemd.

Zaalvoetbal, studiereis, karten, skiën, borrels en conferenties waren allen lang niet zo leuk geweest zonder Molag, Arjen, Joska, Mark, Maarten, Aico, Jeroen, Martijn, Willem, Koray, Maas, Vittorio, Blok, Gerwin, Ole en Nicolas. Vele andere promovendi, post-docs en afstudeerders zijn in de loop der jaren bij IMS en LT gekomen en gegaan. Iedereen heeft op vele manieren bijgedragen aan een geweldige periode zowel op het werk als daarbuiten.

Uiteraard nog een groot woord van dank aan de jongens van het thuisfront. Niets zorgt voor een grotere verbroedering als een lange tafel vol met bruine flesjes in een voetbalkantine. In de regel waren de gemiddelde weekenden en (F1) uitstapjes zo zwaar dat werken op maandagochtend altijd weer een opluchting was.

Mijn familie, en vooral pa en ma, wil ik bedanken voor de aanhoudende interesse in “wat ik nou precies dagelijks aan het doen was”. Bovendien was ik zonder de door jullie geboden kansen nooit gekomen tot waar ik nu ben. Tot slot wil ik Gerdien in het bijzonder bedanken voor de liefde en ondersteuning de afgelopen periode. Hopelijk wordt je geduld beloond met het inhalen van de verloren “quality time” in de afgelopen periode…

Matthijn

Documents

Transparent Conducting Oxides on Polymeric Substrates · 2013-07-14 · given to illustrate the diversity of materials and deposition techniques. In chapter 3 the basis principles