Interface Controlled Organic Thin Films

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¨ Horst-Gunter Rubahn Helmut Sitter Giles Horowitz Katharina Al-Shamery Editors

Interface Controlled Organic Thin Films With 122 Figures

123

Prof. Dr. Horst-Günter Rubahn

Prof. Dr. Helmut Sitter

University of Southern Denmark NanoSYD Mads Clausen Institute Alsion 2, 6400 Sønderborg Denmark [email protected]

Universität Linz Institut für Experimentalphysik Abteilung Festkörperphysik 4040 Linz Austria [email protected]

Prof. Dr. Giles Horowitz

Prof. Dr. Katharina Al-Shamery

Université Paris VII Institut Topologie et de Dynamique Systèmes (ITODYS) 1 rue Guy-de-la-Brosse 75005 Paris France [email protected]

Universität Oldenburg Institut für Reine und Angewandte Chemie 26111 Oldenburg Germany [email protected]

ISSN 0930-8989 ISBN 978-3-540-95929-8 e-ISBN 978-3-540-95930-4 DOI 10.1007/978-3-540-95930-4 Springer Dordrecht Heidelberg London New York Library of Congress Control Number: “PCN applied for” © Springer-Verlag Berlin Heidelberg 2009 This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilm or in any other way, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer-Verlag. Violations are liable to prosecution under the German Copyright Law. The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Cover design: SPi Publisher Services Printed on acid-free paper Springer is part of Springer Science+Business Media (www.springer.com)

Preface Organic electronics is a scientific and technological field that has witnessed an enormous world-wide effort both in basic scientific research as well as in industrial development within the last decades. It is becoming increasingly clear that, if devices based on organic materials are ever going to have a significant relevance beyond being a cheap replacement for inorganic semiconductors, there will be a need to understand interface formation, film growth and functionality. A control of these aspects will allow the realisation of totally new device concepts exploiting the vast flexibility inherent in organic chemistry. The field of devicerelevant “semiconducting” organic materials has many parallels to that of inorganic semiconductors. However, the versatility of organic molecules comes at the cost of higher complexity of the materials. This rules out a 1:1 transfer of concepts established within inorganic semiconductor research to the world of organics, and makes work on organic semiconductors particularly challenging. On a world-wide scale, investigations of organic thin films focus on three main areas with different aims and with a fruitful mixture of applied and basic research: (1) the development and production of devices, (2) thin film characterization and more recently, after recognizing the importance of molecular level control (3) surface and interface science. Linking these branches together creates new synergies and has led and leads to a significant advance in the field of organic semiconductors. Eventually it will result in the development of the necessary tools for tuning device properties on a nanoscopic level. In the last 10 to 15 years a large amount of investigations of devices have been performed with a big range of active organic materials. This work has mapped out the classes of materials that proof useful for single molecule, oligomeric/molecular films and plastic electronics. In this symposium we focused on oligomeric/molecular films, because the control of molecular structures and interfaces provides unprecedently highly defined systems. This in turn allows one to study basic physics and at the same time enables one to find the important parameters necessary to improve organic devices. The E-MRS symposium conceived to bring together the leading groups, which work in the field of growth and characterisation of organic films and devices and focus them on the fabrication and characterisation of highly ordered functional organic films. The wide range of expertise of the contributing groups allowed the combination of different methodologies and aspects of physics, chemistry, and materials science for the design and understanding of well-defined organic structures. In total we received 148 contributions to the symposium in the form of invited talks, oral presentations and posters. Out of them the reviewers selected a representative amount of papers to be published in the proceedings. The main topics discussed at the symposium are reflected in the headlines of the chapters in the proceedings. Introductory review papers based on invited talks given at the symposium are followed by contributed papers. The highlights of the oral and poster presentations contributing to the same topic are summarized in the same chapter. The editors would like to thank the sponsors of the E-MRS symposium, especially the ‘Fonds der Chemischen Industrie’.

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Preface

In technical finishing the book we would like to thank Ms. Zora Milde for her extraordinary help in mastering the handling of all the electronic documents.

Sønderborg, Linz, Paris, Oldenburg January 2009

H.-G. Rubahn H. Sitter G. Horowitz K. Al-Shamery

Contents Preface .................................................................................................................... V A Thin Film Growth............................................................................................ 1 1

Toward an Ab-initio Description of Organic Thin Film Growth ............... 3 P. Puschnig, D. Nabok, and C. Ambrosch-Draxl

2

Organic Nano Fibres from PPTPP .............................................................. 11 F. Balzer, M. Schiek, A. Lützeu, and H.G. Rubahn

3

α-Sexithiophene Films Grown on Cu(110)–(2x1)O: From Monolayer to Multilayers ................................................................................................. 19 M. Oehzelt,, S. Berkebile, G. Koller, T. Haber, M. Koini, O. Werzer, R. Resel, and M.G. Ramsey

4

Para-Sexiphenyl Layers Grown on Light Sensitive Polymer Substrates ....................................................................................................... 23 G. Hernandez-Sosa, C. Simbrunner, T. Höfler, A. Moser, O. Werzer, B. Kunert, G. Trimmel, W. Kern, R. Resel and H. Sitter

5

Thermal Desorption of Organic Molecules................................................. 29 A. Winkler

6

Crystalline Stages of Rubrene Films Probed by Raman Spectroscopy................................................................................................... 37 B.A. Paez, Sh. Abd-Al-Baqi, G.H. Sosa, A. Andreev, C. Winder, F. Padinger, C. Simbrunner, and H. Sitter

7

Rubrene Thin Film Characteristics on Mica .............................................. 43 Sh.M. Abd Al-Baqi, G. Henandez-Sosa, H. Sitter, B. Th. Singh, Ph. Stadler, N.S. Sariciftci

8

Structural Properties of Rubrene Thin Films Grown on Mica ........................................................................................................ ...49 T. Djuric, H.-G. Flesch, M. Koini, Sh.M. Abd Al-Baqi, H. Sitter,and R. Resel

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Rubrene on Mica: From The Early Growth Stage To Late Crystallization................................................................................................ 55 G. Hlawacek, S. Abd-al Baqi, X. Ming He, H. Sitter, and C. Teichert

10 β-Sheeted Amyloid Fibril Based Structures for Hybrid Nanoobjects on Solid Surfaces......................................................................................... ..61 V. Bukauskas, V. Strazdienė, A. Šetkus, S. Bružytė, V. Časaitė,and R. Meškys 11 Characteristics of Vacuum Deposited Sucrose Thin Films .................... ..67 F. Ungureanu, D. Predoi, R.V. Ghita, R.A. Vatasescu-Balcan,and M. Costache 12 Electropolymerization of Polypyrrole Films in Aqueous Solution with Side-Coupler Agent to Hydrophobic Groups.................................. ..73 H.M. Alfaro-López, J.R. Aguilar-Hernandez, A. Garcia-Borquez, M.A. Hernandez-Perez,and G.S. Contreras-Puente. 13 Surface Modification of Polymer Powders by a Far Cold Remote Nitrogen Plasma in Fluidized Bed............................................................. ..79 L. Aiche, H. Vergnes, B. Despax, B. Caussat,and H. Caquineau 14 Features of Polytetrafluoroethylene Coating Growth on Activated Surfaces from Gas Phase .......................................................................... ..85 A.A. Rogachev, S. Tamulevičius, A.V. Rogachev, I. Prosycevas,and M. Andrulevičius 15 Modification of Amorphous Carbon Film Surfaces by Thermal Grafting of Alkene Molecules.................................................................... ..91 H. Sabbah, A. Zebda, S. Ababou-Girard, B. Fabre, S. Députier, A. Perrin, M. Guilloux-Viry,and F. Solal, C. Godet 16 DNA-Controlled Assemblage of Ag Nanoparticles on Solid Surfaces .. ..95 V. Bukauskas, A. Šetkus, I. Šimkienė, J. Sabataitytė, A. Kindurys, and A. Rėza, J. Babonas

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17 Characterization of Self Assembled Monolayer Formation of 11 - Mercaptoundecanoic Acid on Gold Surfaces .............................. 101 J. Stettner, P. Frank, T. Griesser, G. Trimmel, R. Schennach, R. Resel, and A. Winkler 18 SAMs of 11-MUA Grown on Polycrystalline Au-foils by Physical Vapor Deposition in UHV...................................................................................... 107 P. Frank, F. Nussbacher, J. Stettner,and A. Winkler 19 Photoreactive Self Assembled Monolayers for Tuning the Surface Polarity ........................................................................................................ 113 T. Griesser, A. Track, G. Koller, M. Ramsey W. Kern, and G. Trimmel B Traps and Defects ....................................................................................... 119 20 Spectroscopy of Defects in Epitaxially Grown Para-sexiphenyl Nanostructures............................................................................................ 121 A. Kadashchuk, S. Schols, Yu. Skryshevski, I. Beynik, C. Teichert , G. Hernandez-Sosa, H. Sitter, A. Andreev, P. Frank, and A. Winkler 21 Magnetoresistance in Poly (3-hexyl thiophene) Based Diodes and Bulk Heterojunction Solar Cells........................................................ 127 S. Majumdar, H. S. Majumdar, H. Aarnio, R. Laiho, and R. Österbacka 22 Evolution of the Bipolaron Structure in Oligo-diacetylene Films: A Semiempirical Study .............................................................................. 133 M. Ottonelli, G. Musso, and G. Dellepiane C Energy Level Alignment and Charge Transfer ....................................... 139 23 Molecular Orientation Dependence of the Ionization Energy of Pentacene in Thin Films ........................................................................ 141 G. Heimel and N. Koch 24 Charge transfer and Polarization Screening at Organic/Metal Interfaces: Single Crystalline versus Polycrystalline Gold .................... 147 H. Peisert, D. Kolacyak, A. Petershans,and T. Chassé

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25 Sensing Infrared Light with an Organic/Inorganic Hetero-junction.... 153 Gebhard J. Matt, Thomas Fromherz , Guillaume Goncalves, Christoph Lungenschmied, Dieter Meissner, and Serdar N. Sariciftci D Advanced Characterization Methods ....................................................... 159 26 Ultrafast Confocal Microscope for Functional Imaging of Organic Thin Films ................................................................................................... 161 Dario Polli, Michele Celebrano, Jenny Clark, Giulia Grancini, Tersilla Virgili, Guglielmo Lanzani, and Giulio Cerullo 27 Growth and Desorption Kinetics of Sexiphenyl Needles: An In-situ AFM/PEEM Study ..................................................................................... 167 Alexander J. Fleming, Svetlozar Surnev, Falko P. Netzer, and Michael G. Ramsey E

Organic Devices .......................................................................................... 171

28 Temperature Dependence of the Charge Transport in a C60 based Organic Field Effect Transistor ................................................................ 173 Mujeeb Ullah, Th.B. Singh, G.J. Matt, C. Simbruner, G. Hernandz-Sosa, S.N. Sariciftci,and H. Sitter 29 The Influence of Chain Orientation in the Electric Behaviour of Polymer Diodes........................................................................................................... 179 Marta Ramos and Helder Barbosa 30 Interface Modification of Pentacene Ofet Gate Dielectrics .................... 185 Ján Jakabovič, Jaroslav Kováč, Rudolf Srnánek, Jaroslav Kováč jr., Michal Sokolský, Július Cirák, Daniel Haško, Roland Resel, and Egbert Zojer 31 Negative Differential Resistance in C60 Diodes ........................................ 189 Philipp Stadler, Anita Fuchsbauer, Günther Hesser, Thomas Fromherz, Gebhard J. Matt, Helmut Neugebauer, and Serdar N. Sariciftci 32 Performance and Transport Properties of Phthalocyanine: Fullerene Organic Solar Cells ................................................................... 195 M. Rusu, J. Gasiorowski, S. Wiesner, D. Keiper, N. Meyer, M. Heuken, K. Fostiropoulos, and M.Ch. Lux-Steiner1

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33 Organic Transistors Based on Molecular and Polymeric Dielectric Materials...................................................................................................... 199 A. Facchetti, S. DiBenedetto, C. Kim,and T.J. Marks 34 Morphology of the Metal-Organic Semiconductor Contacts: the Role of Substrate Surface Treatment ................................................ 205 A.Petrović, E. Pavlica, and G. Bratina 35 Molecular Interactions Between Alcohols and Metal Phthalocyanine Thin Films for Optical Gas Sensor Applications..................................... 211 S. Uttiya, S. Kladsomboon, O. Chamlek, W. Suwannet, T. Osotchan, T. Kerdcharoen, M. Brinkmann, and S. Pratontep 36 Organic Thin-Film Transistors with Enhanced Sensing Capabilities.................................................................................................. 217 M. Daniela Angione, F. Marinelli, A. Dell’Aquila, A. Luzio , B. Pignataro and L. Torsi 37 Photoelectric Properties of Microrelief Organic/Inorganic Semiconductor Heterojunctions................................................................ 225 N.L. Dmitruk, O.Yu. Borkovskaya, D.O. Naumenko, I.B. Mamontova, N.V. Kotova, O.S. Lytvyn, and Ya.I. Vertsimakha List of Contributors........................................................................................... 229

Toward an Ab-initio Description of Organic Thin Film Growth Peter Puschnig, Dmitrii Nabok and Claudia Ambrosch-Draxl Chair of Atomistic Modelling and Design of Materials, Montanuniversität Leoben, Franz-Josef-Straße 18, A-8700 Leoben, Austria E-mail: [email protected] Abstract. We present an overview of recent ab initio calculations towards the modeling of organic thin film growth. First, we address the intermolecular bonding properties of the oligoacene, oligophenylene and oligothiophene series by density functional theory. By including non-local correlations to account for the van der Waals interactions we achieve excellent agreement of the cohesive energies with available experimental data and obtain surface energies for various low-index planes, thereby emphasizing the importance of dispersive interactions. Second, we review the findings for the interface energy of the model organic/metal interface, thiophene/Cu(110), using the same methodology. Finally, we show how a combination of ab inito results with an empirical force field approach leads to diffusion barriers relevant for organic film growth.

1.

Introduction

In the past years considerable experimental efforts have been placed toward controlling thin film morphologies organic semiconductors by tuning thin film growth conditions [1–3]. A defect-free layer-by-layer growth with a desired orientation of the molecules on a substrate is the key for an optimized device performance. The microscopic quantities which are determining the growth morphologies are surface and interface energies, as well as kinetic parameters, e.g. intra- and interlayer diffusion barriers. However, these numbers are difficult to access from experiments, and hence a theoretical approach for the energetics governing intermolecular as well as molecule/substrate interactions proves particularly useful. Owing to the advances in computer technology first-principles density functional calculations for complex molecular crystals have come into reach. Moreover, recent advances in density functional theory (DFT) have allowed for the incorporation of non-local dispersion forces into the DFT framework [4]. This van der Waals density functional (vdWDF) functional has proven to yield reliable binding energies for a wide class of typical van-der-Waals (vdW) bound systems ranging from noble gas dimers [4], and graphite [5], to the adsorption of benzene and naphtalene on graphite [6], polyethylene crystal structure [7], and the adsorption of thiophene on a metallic surface [8]. Thereby, the viability of the ab initio DFT approach for a number of quantities relevant for organic film growth has been demonstrated. In this article, we show three examples of how state-of-the-art first-principles calculations can contribute to the understanding of organic thin film growth. First, we discuss the cohesive and surface energies of three important types of π-conjugated H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_1, © Springer-Verlag Berlin Heidelberg 2009

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P. Puschnig, D. Nabok and C. Ambrosch-Draxl

molecules, i.e., oligoacenes (nA), oligophenylenes (nP), and oligothiophenes (nT) where n denotes to the number of monomer units [9]. We make use of the vdWDF [4] and obtain cohesive energies in excellent agreement with experimental data and provide a first-principles prediction of surface energies. The latter allows us to estimate the equilibrium crystal shapes for all studied compounds. Second, we show results for the interface energy of a model organic/metal junction, thiophene/Cu(110) [8], and demonstrate its implications for resulting growth modes. Finally, we review recent findings for the step-edge barrier observed in growth of para-sexiphenyl [10].

2.

Computational Details

Density Functional Calculations. The total-energy and force calculations are performed within the framework of DFT as implemented in the plane-wave package PWSCF [11] using ultrasoft pseudo-potentials [12] with a plane-wave energy cutoff of 40 Ry. Bulk energies are computed by adopting the experimentally known space groups and lattice constants and relaxing the internal atomic positions. Isolated molecules are treated by supercells, while surface energies are calculated by the repeated-slab approach. Only one molecular layer turned out to be sufficient to converge surface energies to within 10 meV. Vacuum distances of 10 Å were found to be adequate to prevent interaction between translational images for isolated molecules as well as slab geometries. We use three different approximations for the exchange-correlation energy. Apart from the standard functionals like the local density approximation (LDA) [13] and the generalized gradient approximation (GGA) [14] we also employ the van der Waals density functional (vdWDF) in a non-selfconsistent approach [4]. The molecular cohesive energy Ecoh is defined as the energy reduction upon forming a crystal from isolated molecules, i.e., Ecoh = Emol – Ebulk/2, where Ebulk and Emol denote the total energies of the bulk and the isolated molecule, respectively. The factor 2 takes into account the number of molecules in the unit cell. By this definition the cohesive energy is positive for any stable crystal. The surface energy is defined as the energy required for cleaving a surface from a bulk material. The surface energy can be specified either per surface unit cell as Esurf = ½(Eslab – Ebulk), or per surface area A as γ = ½(Eslab – Ebulk)/A. Here, Eslab denotes the total energy of the slab configuration, and the factor ½ stands for the fact that the slab contains two surfaces.

3.

Results

Cohesive Energies. The bulk phases of the nA, nP, and nT series are characterized by the so-called herringbone packing where two inequivalent molecules in the ab plane form layers perpendicular to the crystallographic c direction. At room temperature and ambient pressure the space group of the considered molecular crystals is monoclinic, except for tetracene (4A) and pentacene (5A) which crystallize in a triclinic space group.

Toward an Ab-initio Description of Organic Thin Film Growth

5

Cohesive energies for the oligoacene series are presented in Figure 1 where DFT results obtained within the LDA, the GGA and the vdWDF are compared to measured data. For the whole oligoacene series we find excellent agreement between experimental and vdWDF values. The LDA results systematically underestimate the measurement by almost 25% while GGA yields practically zero cohesive energy. This demonstrates the inherent problem of local (LDA) or semilocal functionals (GGA) when applied to weakly bound systems, namely, that these standard approximations do not capture the essential physics of van der Waals interactions requiring a truly non-local functional. Even though LDA reproduces the experimental trend in a reasonable manner [15] it should not be used for predicting intermolecular binding properties since it relies on the wrong physical picture. Only through its general overbinding effect does it mimic a compensation for missing dispersive interactions. It is noteworthy that the cohesive energy is an almost linear function of molecular length which appears reasonable owing to the same type of crystalline packing for the whole oligomer series. The situation is similar for the oligophenylenes, but somewhat more complicated. Here we have to consider the non-planar nature of the isolated nP molecules, which exhibit twisted interring bonds in the gas phase. Allowing for such a conformation of the isolated molecule leads to a smaller Ecoh compared to the assumption of planar molecules. For instance, the effect on the cohesive energy of biphenyl is found to be 0.1 eV, where the torsion angle corresponding to the energy minimum is about 37°, and we obtain cohesive energies of 0.98, 1.44, 1.90, and 2.82 eV for 2P, 3P, 4P, and 6P, respectively. Finally, also the oligothiophenes show a similar trend with respect to oligomer length as the nA and the nP series. Since there is no experimental data available for the thiophene oligomers, our vdW-DF data – 0.95, 1.75, and 2.78 eV for 2T, 4T, and 6T, respectively – can be viewed as an accurate theoretical prediction.

Fig. 1. Cohesive energies of the oligoacene series from napthtalene (2A) to pentacene (5A). Theoretical values obtained within the van-der-Waals density functional (vdW-DF), the local densiy approximation (LDA), and the generalized gradient approximation (GGA) are compared to experimental data.

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Relating the cohesive energy to the number of carbon atoms per molecule one finds very similar results for the acenes (85, 79, 79, and 76 meV/C atom) and phenylenes (82, 80, 79, 79 meV/C atom), while the presence of sulfur atoms in the thiophenes increases the cohesive energies by roughly 10%. Surface Energies. Concerning organic thin film morphologies, the anisotropy in the surface energies is an important quantity controlling the orientation and shape of crystals during the growth. The (100), (010), (001), and (110) surface energies for the complete nA, nP, and nT series are given in Figure 2. A common feature for all compounds is that the (001) plane exhibits the lowest surface energy. This can be understood in terms of the crystal packing in that direction which is governed by rather weak H—H interlayer interactions. Hence, on substrates with comparably small substrate/molecule interactions thin films are expected to wet the substrate and to be preferentially (001) oriented since in that way the total surface free energy is minimized.

Fig. 2. Surface energies within the vdWDF for the oligoacene (nA), oligophenylene (nP), and the oligothiophene (nT) series. Values for the (100), (010), (001), and (110) planes are given and the corresponding surface terminations are depicted for tetracene (4A).

This fact is also illustrated by the equilibrium crystal shapes (ECS) as exemplified in Figure 3 for 3A where the (001) crystal faces have the largest areas. These are based on Wulff’s construction using the surface energies as summarized in Figure 2. Compared to experiment we find excellent agreement of our equilibrium crystal shape with a recent investigation on the growth of anthracene on graphite [16]. In particular the (001) orientation of the crystal parallel to the substrate and the appearance of approximately hexagonally shaped facets of (100) and (110) planes and a small (010) facet is in line with our calculated ECS.

Toward an Ab-initio Description of Organic Thin Film Growth

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Fig. 3. Calculated equilibrium crystal shape of anthracene (filled octagon) compared to a photograph of the epitaxial growth of an anthracene single crystal on a graphite (0001) substrate taken from Ref. [16]. Interface Energies. In the previous section we have learned that characteristic surface energies of organic molecular crystals are in the range of 80–160 mJ/m2. For growth on inorganic substrates these numbers for the adsorbate surface energy γa are to be compared to typical surface energies of the inorganic substrate γs which are typically an order of magnitude larger. This is of course due to the fact that in inorganic materials strong covalent bonds have to be broken at the surface while in organic molecular crystals only comparably weak van der Waals forces have to be overcome to create a surface from the bulk. Close to thermodynamic equilibrium the magnitude of the substrate surface energy γs, the adsorbate surface energy γa, and the interface energy γi determine the resulting growth mode by the requirement of minimizing the total free energy. Therefore, it is important to calculate the organic/inorganic interface energy by employing a reliable and accurate ab-initio approach. As a model system we have studied the interaction between a thiophene molecule and the Cu(110) surface [8]. By employing vdWDF as previously for the investigation of the surface energies we have found van der Waals interactions to play a crucial role also for this organic/metal junction. This is illustrated in Figure 4 where we compare DFT results obtained within the LDA and GGA to the vdWDF results. We find an adsorption energy of Ea = 0.5 eV which is almost twice as large as the corresponding GGA value, and more than two times smaller than the adsorption energy predicted by LDA. We should emphasize the fact that our vdWDF value is in good agreement with the experimental Ea reported for thiophene on Cu(100) [17]. Using the definition, γi = γs + γa – Ea/A, for the interface energy as sum of the substrate and adsorbate surface energies minus the adsorption energy per area A we arrive at a value of γi = 1.55 J/m2 for the thiophene/Cu(110) interface. For this finding we have taken into account the surface energy of Cu(110) of γs = 1.70 mJ/m2 and approximated the surface energy of thiophene γa by the (100) value of bithiophene which is 0.15 mJ/m2 according to Figure 2.

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Fig. 4. Left: Adsorption energy of a thiophene ring on a Cu(110) surface calculated within the LDA, the GGA, and the van der Waals density functional [8]. Right: Schematic representation of the thiophene/Cu(110) adsorption.

In inorganic film growth the difference Δγ = γa + γi – γs is known to be a good indicator for the type of growth mode. A negative value for Δγ leads to layerby-layer (Frank van der Merve) growth since it is energetically favorable to cover the substrate completely by the adsorbate, while a positive value points towards three-dimensional island (Volmer-Weber) growth [18]. Evaluating Δγ for the thiophene/Cu(110) model system we arrive at a value close to 0. This indicates the intermediate case referred to as Stranski-Krastanov growth mode in which layerby-layer growth is succeeded by three-dimensional island growth after a critical adsorbate thickness. While in inorganic film growth the critical layer thickness depends on the strain built up in the deposited film, the situation in organic growth is more complicated due to the anisotropy and the internal degrees of freedom of the organic building blocks. For instance, wetting layers having various molecular conformation or orientation may be formed compared to the molecular orientation at later growth stages. Also, taking into account values for Δγ ≅ 0 as observed for the thiophene/Cu(110) model system we can expect a diversity of growth morphologies as is indeed observed in organic film growth. Kinetic Barriers. The discussion above assumes that the energetically lowest configuration can indeed be reached during the thin film growth process. However, kinetic barriers often limit the mobility of the deposited molecules and give rise to additional phenomena. In this section we exemplify the consequence of a step edge-barrier on the resulting growth morphology for the growth of 6P on a modified mica surface. According to Figure 2, the crystal plane exhibiting the lowest surface energy in 6P is the (001) plane resulting in γa(001) = 0.11 eV. Since the interaction between a (001) terminated 6P layer and the mica substrate is expected to be small one can assume the interface energy for that system to be smaller than the surface energy of the mica substrates. This would result in a layer-by-layer growth of upright standing 6P molecules. Instead one observes

Toward an Ab-initio Description of Organic Thin Film Growth

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the formation of terraced mounds which can be explained by the existence of an additional step-edge barrier (Ehrlich-Schwoebel barrier). An analysis of these growth mounds as a function of coverage by atomic force microscopy revealed a sizeable Ehrlich-Schwoebel barrier (ESB) of 0.67 eV [10]. By combining firstprinciples density functional calculations with an empirical force field method we were able to calculate the total activation barrier to be 0.63 eV. By taking into account the computed on-terrace diffusion barrier of only 0.02 eV, the ESB results in 0.61 eV which is in excellent agreement with the experimental value [10]. Moreover, the calculation of the transition state provided further insight into the nature of the step-edge diffusion barrier, by showing that the 6P molecule diffuses toward the [100] edge with its long axis perpendicular to the edge and gradually slides down the (100) plane by bending over the edge between (001) and (100) facets. This is illustrated in Figure 5 which depicts four snap shots of the diffusion path for the step-edge crossing.

Fig. 5. Four snapshots of the diffusion path for step-edge crossing of sexiphenyl on a (001) 6P surface. The energy of each configuration is indicated, where the second structure corresponds to the transition state with an activation barrier of 0.63 eV.

4.

Conclusions

In summary, we have presented an ab-initio study of bonding properties in molecular crystals by consistently taking into account non-local van der Waals interactions. We reliably produce correct equilibrium layer distances to within 0.2 Å. This new level of precision enables us to obtain surface energies with an estimated accuracy of 5% allowing for accurate predictions of equilibrium crystal shapes and thin film morphologies where subtle energy differences play an important role. Moreover, our results open the perspective towards future investigations aiming at kinetic parameters for organic thin film growth. Acknowledgements. The work was supported by the Austrian Science Fund, Project No. S9714 within the National Science Network “Interface Controlled and Functionalised Organic Films”.

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References 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18

F.-J. M. zu Heringdorf, M. C. Reuter, and R. M. Tromp, Nature 412, 51, 2001. M. A. Loi, E. da Como, F. Dinelli, M. Murgia, R. Zamboni, F. Biscarini, and M. Muccini, Nat. Mat. 4, 81, 2005. G. Hlawacek, Q. Shen, C. Teichert, R. Resel, and D. M. Smilgies, Surf. Sci. 601, 2584, 2007. M. Dion, H. Rydberg, E. Schröder, D. C. Langreth, and B. I. Lundqvist, Phys. Rev. Lett. 92, 246401, 2004. H. Rydberg, M. Dion, N. Jacobson, E. Schröder, P. Hyldgaard, S. I. Simak, D. C. Langreth, and B. I. Lundqvist, Phys. Rev. Lett. 91, 126402, 2003. S. D. Chakarova-Käck, E. Schröder, B. I. Lundqvist, and D. C. Langreth, Phys. Rev. Lett. 96, 146107, 2006. J. Kleis, B. I. Lundqvist, D. C. Langreth, and E. Schröder, Phys. Rev. B 76, 100201(R), 2007. P. Sony, P. Puschnig, D. Nabok, and C. Ambrosch-Draxl, Phys. Rev. Lett. 99, 176401, 2007. D. Nabok, P. Puschnig, and C. Ambrosch-Draxl, Phys. Rev. B 77, 245316, 2008. G. Hlawacek, P. Puschnig, P. Frank, A. Winkler, C. Ambrosch-Draxl, and C. Teichert, Science 321, 108, 2008. S. Baroni, A. D. Corso, S. de Gironcoli, P. Giannozzi, C. Cavazzoni, G. Ballabio, S. Scandolo, G. Chiarotti, P. Focher, and A. Pasquarello, http://www.pwscf.org/, 2007. D. Vanderbilt, Phys. Rev. B 41, 7892, 1990. J. P. Perdew and A. Zunger, Phys. Rev. B 23, 5048, 1981. J. P. Perdew, K. Burke, and M. Ernzerhof, Phys. Rev. Lett. 77, 3865, 1996). J. E. Northrup, M. L. Tiago, and S. G. Louie, Phys. Rev. B 66, 121404(R), 2002. S. Jo, H. Yoshikawa, A. Fujii, and M. Takenaga, Surf. Sci. 592, 37, 2005. B. A. Sexton, Surf. Sci. 163, 99, 1985. A. Groß, Theoretical Surface Science – A Microscopic Perspective, Springer, Berlin, 2003.

Organic Nanofibers from PPTPP Frank Balzer1, Manuela Schiek1, Arne Lützeu2 and Horst-Günter Rubahn1 1

Syddansk Universitet, Mads Clausen Institute, NanoSYD, Alsion 2, DK-6400 Sønderborg, Denmark E-mail: [email protected] 2 University of Bonn, Kekulé-Institute of Organic Chemistry and Biochemistry, Gerhard-Domagk-Str. 1, D-53121 Bonn, Germany Abstract. The growth of 2,5-Di-4-biphenyl-thiophene (PPTPP) on the dielectric substrates NaCl, KCl, KAP, muscovite mica, and phlogopite mica is investigated by atomic force microscopy (AFM) and fluorescence microscopy. In all cases fibers are formed with several ten nanometers height and several hundred nanometers width, respectively. Only for PPTPP on muscovite mica the fibers are mutually parallel aligned along a single substrate direction, i.e. along muscovite 〈110〉. This uniaxial growth is explained by an electrostatic interaction between the molecules and surface electric fields in combination with epitaxy. The various growth directions on other substrates are dictated by epitaxy alone.

1.

Introduction

Nanofibers from organic conjugated molecules such as bare and functionalized para-phenylenes [1,2], α-thiophenes [3], phenylene-thiophene co-oligomers [4,5], or coumarin derivatives [6] have a high application potential in future optoelectronic devices [7]. Waveguiding [8–10], tunable light emission [11], gain narrowing and lasing [12–14], as well as frequency doubling [15,16] have already been observed. In the past the growth of the blue-light emitting para-hexaphenylene (p-6P) molecules on muscovite mica and on KCl has been investigated in detail [17–20]. For p-6P on muscovite mutually parallel aligned needles along a single 〈110〉 substrate direction have been observed, their growth being steered by the anisotropic electric surface fields [21]. This 〈110〉 direction corresponds to the direction of grooves on the muscovite surface, which alternate by 120° in between consecutive cleavage planes. This direction is denoted as 〈110〉g, the non-grooved one as 〈110〉ng. The long molecule axis is oriented at ±76° with respect to muscovite 〈110〉g, resulting in a single energetically favorable needle orientation. Low energy electron diffraction (LEED) and thermal desorption spectroscopy (TDS) have shown, that the first growth step is the formation of a wetting layer from lying molecules. Then clusters grow, and finally these clusters aggregate into needles. On KCl no such wetting layer and almost no clusters have been detected [22]. Needles grow along the two 〈110〉 substrate directions simultaneously [23]. For the investigated rod-like molecules so far the possible needle directions on muscovite depend on the epitaxial orientation of the molecules on the substrate, and on the packing of these molecules within a needle [24]. Therefore parafunctionalized para-phenylenes show a very similar growth behavior compared to p-6P. The α-thiophenes quaterthiophene and sexithiophene, however, align with their long molecular axes along the two muscovite high symmetry directions without H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_2, © Springer-Verlag Berlin Heidelberg 2009

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a groove, i.e. along 〈110〉ng, and [100], and therefore three needle orientations are accomplished simultaneously. That way fortunate optical properties of an uniaxially aligned needle film such as polarized light absorption and emission are lost. To preserve some of the phenylene growth properties but, e.g., change the emission color substantially we therefore have investigated the growth of a thiophene phenylene co-oligomer: 2,5-Di-4-biphenyl-thiophene (PPTPP), Fig. 1. Here we show first results for growth experiments on different substrate surfaces under similar growth conditions, i.e. similar substrate temperatures Ts, deposition rates, and nominal film thicknesses.

2.

Experimental Methods

2,5-Di-4-biphenyl-thiophene has been synthesized in a two-fold Suzuki crosscoupling reaction from commercially available 2,5-dibromo-thiophene and 4-biphenyl boronic acid, using 5 mol% tetrakis(triphenylphosphino)palladium as catalyst together with cesium fluoride as base in dry tetrahydrofurane. The desired product has been obtained in yields of 80% after refluxing for 50 h. The final product precipitated from the reaction mixture and was washed with water and organic solvents repeatedly for purification. By outgassing in vacuo residual organic solvents are removed to give the desired compounds in high purity. Note that PPTPP has been synthesized previously using very similar approaches [25,26]. HO Br Br

+ 2

S

B HO

Pd(PPh3)4 + CsF THF, Δ, 50 h

Suzuki Cross Coupling Conditions

S 80 % yield

PPTPP citreous amorphous solid

Fig. 1. Schematics of the synthesis of 2,5-Di-4-biphenyl-thiophene (PPTPP).

As substrate materials five different single crystals have been chosen: potassium chloride (KCl), sodium chloride (NaCl), potassium acid phthalate (KAP) [27], muscovite mica [28], and phlogopite mica. All are cleaved in air and are transferred immediately into a high vacuum system (base pressure 2 × 10–8 mbar). Either right away or after annealing a clear low energy electron diffraction (MCPLEED, Omicron) pattern is observed for each of them. Organic molecules are deposited from an effusion cell. The nominal deposited film thickness is estimated by a water-cooled quartz microbalance located next to the substrate.

Organic Nanofibers from PPTPP

13

After deposition the samples are characterized in situ by LEED, ex situ by atomic force microscopy (AFM, JPK NanoWizard) and fluorescence microscopy (excitation wavelength λexc = 365 nm from a high-pressure mercury lamp).

3.

Results and Discussion

Figure 2 shows 150 × 150 µm2 fluorescence microscope images of typical samples deposited at Ts = 350 K – 380 K with a deposition rate of 0.1 Å/s – 0.2 Å/s and nominal thicknesses up to 5 nm. On all of those substrates needle-like structures from PPTPP form. All emit blue-green light after normal incidence UV irradiation. The emitted fluorescence from the needles is strongly polarized, the polarization vector being oriented perpendicular to the local needle direction. Similar to the case of, e.g., p-6P [21,29,30] this points to fibers made from lying organic molecules, the long molecular axis being perpendicular to the long needle axis. In addition for all substrates except for muscovite mica a green light emitting background is visible. This can be clearly seen in Fig. 2(c), where the border between the bare substrate and the deposition area is imaged. Typical fluorescence spectra are presented in Fig. 3. A well resolved vibronic progression is visible between 400 nm (a)

(b)

(d)

(c)

(e)

Fig. 2. Fluorescence microscope images, 150 × 150 µm2, of PPTPP on (a) NaCl, (b) KCl, (c) KAP, (d) muscovite mica, and (e) phlogopite mica. The nominal thickness of all samples is between 2 nm and 5 nm, the surface temperature during deposition varies between Ts = 350 K – 380 K.

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Fig. 3. Unpolarized measured fluorescence spectra of PPTPP needle films on different substrates after excitation with unpolarized λexc = 365 nm light under normal incidence. The dotted vertical lines mark the peak positions at 437 nm, 463 nm, 493 nm, 529 nm, and 570 nm.

and 600 nm with an energy spacing of approximately 1300 cm–1, being most prominent in the case of muscovite and phlogopite mica. Slightly different colors in the microscopy images stem from different relative intensities of the excitonic transitions. The reason for the distribution of intensity between the fluorescence peaks is still unknown, but might be related to different molecule–molecule interactions for the molecules forming the patches of upright molecules, see below, and for molecules forming needles [31]. Although needle-like structures are formed in all five cases, the realized needle directions depend very much on the substrate. On NaCl, KCl, and KAP two needle directions evolve, with different angles in between. On NaCl and KCl the angle is 90°, and the needles grow along the two substrate 〈110〉 directions simultaneously. On KAP the angle is 68°, corresponding to the angle between the two KAP 〈101〉 directions [32]. On muscovite mica all of the needles grow along a single 〈110〉 direction, with altogether two orientational domains being present on the whole sample. On phlogopite mica three different growth directions exist, most pronounced at the very early growth stages and for deposition at room temperature. For larger coverages and higher substrate temperatures the needles tend to curl and to form rings. To a lesser extent this curling is also observed for PPTPP on NaCl and muscovite, Figs. 2(a) and (d), the tendency increasing with the overall needle length. Corresponding AFM images (see Fig. 4) provide widths and heights of the needles, but also show in more detail the greenish background from Fig. 3. This background results from patches of one or two multiples of approximately 2.2 nm height, suggesting that they are formed from upright molecules [25]. This background appears for all substrates except for muscovite mica, where clusters appear instead. These clusters are also found on phlogopite mica. They do not exist directly besides the needles (denuded zones).

Organic Nanofibers from PPTPP (a)

(b)

(c)

(d)

(e)

(f)

15

(g)

Fig. 4. Atomic force microscope images, 40 × 40 µm2, of PPTPP on (a) NaCl, (b) KCl, (c) KAP, (d) muscovite mica, and (e) phlogopite mica, corresponding to the fluorescence microscope images from Fig. 2. The height scales are (a) – (c) 100 nm, and (d) – (e) 50 nm. Arrows emphasize substrate directions. In the 5 × 2.5 µm2 AFM images for muscovite (f) and phlogopite (g) the clusters from lying molecules as well as layers from upright ones are clearly visible.

4.

Conclusions and Outlook

Obviously the growth directions of the needles depend strongly on the crystal structure of the underlying substrates. The symmetry of the substrate is retained in the needles’ growth directions. Similar to the case of p-6P we attribute this observation to an epitaxial relationship of the organic molecules with the corresponding substrate. On NaCl and on phlogopite the alignment is not as perfect as for KCl and muscovite mica, respectively. The uniaxial growth on muscovite mica is explained by an electrostatic interaction between the molecules and surface electric fields in combination with epitaxy. Epitaxy leads to an alignment of the molecules along muscovite high symmetry directions, whereas the electric fields choose the energetically most favorable of the three possible growth directions: needles grow along 〈110〉g. Muscovite and phlogopite mica exhibit almost identical lattice constants and surface compositions, but differ in that phlogopite is a trioctahedral mica, whereas muscovite is a dioctahedral one. This leads to the already describe grooves along a single 〈110〉 direction on muscovite, which are missing on phlogopite [33]. Therefore on phlogopite three simultaneous needle directions exist, on muscovite only one.

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Not only the realized needle directions, but also the growth mechanism is similar to the case of p-6P. The observed clusters on muscovite and phlogopite are remnants from the initial growth stage. Only when the cluster number density reaches a critical value, needles start growing, mainly by agglomeration of the clusters. A LEED pattern from a wetting layer of lying molecules is observed for the case of muscovite mica. For all other cases no such diffraction pattern has been detected. As a conclusion the growth mechanism of PPTPP and the realized needle directions on the different substrates are similar to that of p-6P. Understanding such basic growth principles allows one to predict qualitatively nanowire surface growth from other conjugated molecules and thus allows for a sophisticated design of new devices. Adding, for example, another thiophene ring next to the existing one to form 5,5´-Di-4-biphenyl-2,2´-bithiophene (PPTTPP) does not alter the basic growth mode, but changes the epitaxial alignment on muscovite. From that two simultaneous needle directions on muscovite are predicted and are actually observed [24]. Acknowledgements. We are grateful to the Danish research agencies FNU and FTP as well as the Danish Advanced Technologies Trust for supporting this work by various grants. MS and AL thank the German research foundation DFG for financial support.

References 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15

F. Balzer and H.-G. Rubahn, Adv. Funct. Mater., 15, 17, 2005. M. Schiek, F. Balzer, K. Al-Shamery, A. Lützeu, and H.-G. Rubahn, Soft Matter, 4, 277, 2008. F. Balzer, L. Kankate, H. Niehus, and H.-G. Rubahn, Proc. SPIE, 5724, 285, 2005. H. Yanagi, T. Morikawa, S. Hotta, and K. Yase, Adv. Mater. 13, 313, 2001. F. Balzer, M. Schiek, A. Lützeu, K. Al-Shamery, and H.-G. Rubahn, Proc. SPIE 6470, 647006, 2007. M. Mille, J.-F. Lamere, F. Rodrigues, and S. Fery-Forgues, Langmuir 24, 2671, 2008. K. Al-Shamery, H.-G. Rubahn, and H. Sitter, editors. Organic Nanostructures for Next Generation Devices, Vol. 101 of Springer Series in Materials Science, Berlin 2008. K. Takazawa, Chem. Phys. Lett. 452, 168, 2008. H. Yanagi and T. Morikawa, Appl. Phys. Lett. 75, 187, 1999. F. Balzer, V.G. Bordo, A.C. Simonsen, and H.-G. Rubahn, Phys. Rev. B 67, 115408, 2003. Y.S. Zhao, H. Fu, F. Hu, A. Peng, W. Yang, and J. Yao, Adv. Mater. 20, 79, 2008. F. Quochi, F. Cordella, A. Mura, G. Bongiovanni, F. Balzer, and H.-G. Rubahn, J. Phys. Chem. B 109, 21690, 2005. F. Quochi, F. Cordella, R. Orru, J.E. Communal, P. Verzeroli, A. Mura, G. Bongiovanni, A. Andreev, H. Sitter, and N.S. Sariciftci, Appl. Phys. Lett. 84, 4454, 2004. H. Yanagi, T. Ohara, and T. Morikawa, Adv. Mater. 13, 1452, 2001. F. Balzer, J. Brewer, J. Kjelstrup-Hansen, M. Madsen, M. Schiek, K. Al-Shamery, A. Lützeu, and H.-G. Rubahn, Proc. SPIE 6779, 67790I, 2007.

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16 J. Brewer, M. Schiek, A. Lützeu, K. Al-Shamery, and H.-G. Rubahn, Nano Lett. 6, 2656, 2006. 17 L. Kankate, F. Balzer, H. Niehus, and H.-G. Rubahn, J. Chem. Phys. 128, 084709, 2008. 18 A. Andreev, T. Haber, D.-M. Smilgies, R. Resel, H. Sitter, N.S. Sariciftci, and L. Valek, J. Cryst. Growth 275, e2037, 2005. 19 A.Y. Andreev, C. Teichert, G. Hlawacek, H. Hoppe, R. Resel, D.-M. Smilgies, H. Sitter, and N.S. Sariciftci, Org. Electron. 5, 23, 2004. 20 P. Frank, G. Hlawacek, O. Lengyel, A. Satka, C. Teichert, R. Resel, and A. Winkler, Surf. Sci. 601, 2152, 2007. 21 F. Balzer and H.-G. Rubahn, Appl. Phys. Lett. 79, 3860, 2001. 22 P. Frank, G. Hernandez-Sosa, H. Sitter, and A. Winkler, Thin Solid Films 516, 2939, 2008. 23 T. Mikami and H. Yanagi, Appl. Phys. Lett. 73, 563, 1998. 24 F. Balzer, M. Schiek, K. Al-Shamery, A. Lützeu, and H.-G. Rubahn, J. Vac. Sci. Technol. B 26, 2008. In print. 25 T.J. Dingemans, N.S. Murthy, and E.T. Samulski, J. Phys. Chem. B 105, 8845, 2001. 26 S. Hotta, H. Kimura, S.A. Lee, and T. Tamaki, J. Heterocycl. Chem. 37, 281, 2000. 27 M. Campione, A. Sassella, M. Moret, A. Papagni, S. Trabattoni, R. Resel, O. Lengyel, V. Marcon, and G. Raos, J. Am. Chem. Soc. 128, 13378, 2006. 28 E.W. Radoslovich, Acta Cryst. 13, 919, 1960. 29 A. Andreev, G. Matt, C.J. Brabec, H. Sitter, D. Badt, H. Seyringer, and N.S. Sariciftci, Adv. Mater. 12, 629, 2000. 30 A. Niko, E. Zojer, F. Meghdadi, C. Ambrosch-Draxl, and G. Leising, Synth. Met. 101, 662, 1999. 31 E. Da Como, M.A. Loi, M. Murgia, R. Zamboni, and M. Muccini, J. Am. Chem. Soc. 128, 4277, 2006. 32 A.V. Alex and J. Philip, J. Appl. Phys. 88, 2349, 2000. 33 Y. Kuwahara, Phys. Chem. Miner. 28, 1, 2001.

α-Sexithiophene Films Grown on Cu(110)-(2x1)O: From Monolayer to Multilayers Martin Oehzelt1,2, Stephen Berkebile2, Georg Koller2, Thomas Haber2, Markus Koini3, Oliver Werzer3, Roland Resel3, and Michael G. Ramsey2 1

Institute of Experimental Physics, Johannes Kepler University Linz, Altenbergerstraße 69, A-4040 Linz, Austria E-mail: [email protected] 2 Institute of Physics, Karl-Franzens University Graz, Universitätsplatz 5, A-8010 Graz, Austria 3 Institute of Solid State Physics, Graz University of Technology, Petersgasse 16, A-8010 Graz, Austria Abstract. The growth of α-sexithiophene (6T) on copper (110) and oxygen reconstructed Cu(110) is studied by multiple techniques such as STM (scanning tunnelling microscopy), XRD (X-ray diffraction), XPS (X-ray photoelectron spectroscopy) and NEXAFS (near edge X-ray absorption fine structure). Selected data will be presented here and we will show that the long axes of the molecules on Cu(110) and Cu(110)-(2x1)O (CuO) are aligned along the valleys of the surface corrugations, i.e. along [1–10] and [001], respectively. With GIXD (grazing incidence X-ray diffraction) measurements the monolayer structure of 6T on Cu-O could be determined. Thicker films were studied by the X-ray diffraction pole figure technique. On all surfaces the (010) net planes of the bulk crystal structure are parallel to the surface i.e. the films grow exclusively (on Cu-O) or pre- dominantly (on Cu) in the (010) orientation. In the thick film the long molecular axes of the 6T molecules are found to be parallel to those of the monolayer. To study the transition from the monolayer to the multilayer structures NEXAFS measurements were carried out.

1.

Introduction

Organic semiconductors have attracted considerable interest in the last decades and as a result remarkable progress in the field of organic electronics has been made [1]. Especially their potential use in low-cost devices such as organic solar cells, light emitting diodes or thin film transistors have encouraged many investigations including basic research studies [2]. Here we present data on the growth of α-sexithiophene (6T), as this molecule is one of the most prominent organic transistor materials [3] and the focus will be on structural changes of 6T on the Cu(110)-(2x1)O surface from the initial nucleation to thick films. Additional information can be found in the following papers [4–7].

2.

Results and Discussion

The STM image of Fig. 1a reveals the molecular orientation of the 6T molecules in the monolayer regime. The long molecular axis is parallel to the substrate H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_3, © Springer-Verlag Berlin Heidelberg 2009

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M. Oehzelt et al.

surface and parallel to the Cu-O rows which is the Cu[001] direction. The molecules are stacked in rows perpendicular to the long molecular axis. While the order between the stacks is less pronounced, the distance between neighboring molecules is 5.1 Å and commensurate to the Cu-O rows. This ordering is also visible in the Fourier transform of the image – sharp points for the 5.1 Å periodicity and streaks in the perpendicular direction.

Fig. 1. STM images of a monolayercoverage of 6T on Cu(110)-(2x1)O. Tunneling parameter: Vt = +2.0V, It = 1nA. The high symmetry directions are indicated and the long molecular axes are parallel to [001]. The insert is the fourier transformation of the STM image which reveals the characteristic distances between the molecules. Figure 1b is a GIXD pattern for net-planes perpendicular to the long molecular axis and reflecting the periodicity between molecular stacks in the monolayer regime of 27 Å.

For thicker films deposited on Cu-O, the molecules crystallize in their bulk crystal structure [4] with 6T(010) net-planes parallel to the substrate. The in-plane orientation is such that the long molecular axes are still aligned along the Cu-O rows [4] and therefore they clearly adopt the molecular orientation of the monolayer. Note that in the 6T(010) net-plane the molecules are still organized in stacks, but in contrast to the monolayer this surface unit cell is oblique and the stacking direction is not 90° with respect to the long molecular axis but 66.5°. The thickness where this transition from a parallel stacking to a oblique one takes place is still subject of further investigations. Nevertheless, recent grazing incidence X-ray diffraction (GIXD) experiments have shown that 6T films with bulk crystallites still show the ordered monolayer structure underneath. In Fig. 1b a GIXD graph is shown where a series of peaks originating from the monolayer are shown. The distances determine the spacing between the stacks and picking up the 27 Å periodicity of the long molecular axis. Figure 2 shows the results from a series of NEXAFS measurements. Details about the measurement geometry and the data evaluation can be found in [4] and references therein. In summery the NEXAFS measurements reveal that the molecular tilt angle in the monolayer is 4° higher than in the multilayer. This result combined with the knowledge that the monolayer structure with a spacing of 5.1 Å is reduced compared to the multilayer one leads to the following model: The spacing between adjacent molecules determined by the Cu-O rows ends up

α-Sexithiophene Films Grown on Cu(110)-(2x1)O: From Monolayer to Multilayers 21

in a higher tilt angle to keep their Van der Waals distance (d) constant (see Fig. 2). With increasing thickness the bulk structure develops. How the different stacking within the molecular rows comes into play and how rapidly the transition from the monolayer to the multilayer occurs are still open questions.

Fig. 2. NEXAFS spectra for a thick 6T layer (left) and the monolayer (right) grown on Cu-O are shown (for details see also Ref. 5). The model in the middle shows the change in the molecular packing according to the NEXAFS, STM and XRD measurements.

Acknowledgements. This work was supported by the Austrian Science Foundation (FWF) and by the European Synchrotron Research Facility (ESRF).

References 1 2 3 4 5 6 7

N. Koch in ChemPhysChem, Vol. 8, 1438, 2007. A. Facchetti, M.H. Yoon, J.T. Marks in Advanced Materials, Vol. 17, 1705, 2005. H.E. Katz in Journal of Materials Chemistry, Vol. 7, 376, 1997. M. Oehzelt, G. Koller, J. Ivanco, S. Berkebile, T. Haber, R. Resel, F.P. Netzer, M.G. Ramsey in Advanced Materials, Vol. 18, 2466, 2006. M. Oehzelt, L. Grill, S. Berkebile, G. Koller, F.P. Netzer, M.G. Ramsey in ChemPhysChem, Vol. 8, 1707, 2007. M. Oehzelt, S. Berkebile, G. Koller, J. Ivanco, S. Surnev, M.G. Ramsey, in preparation. M. Koini, T. Haber, O. Werzer, S. Berkebile, G. Koller, M. Oehzelt, M.G. Ramsey, R. Resel, in preparation.

Para-Sexiphenyl Layers Grown on Light Sensitive Polymer Substrates G. Hernandez-Sosa1, C. Simbrunner1, T. Höfler2, A. Moser3, O. Werzer3, B. Kunert3, G. Trimmel2, W. Kern2,4, R. Resel3 and H. Sitter1 1

Institute for Semiconductors and Solid State Physics, Johannes Kepler University Linz, Altenbergerstrasse 69, 4040 Linz, Austria 2 Institute for Chemistry and Technology of Materials, Graz Technical University, Stremayrgasse 16, 8010 Graz, Austria 3 Institute of Solid State Physics, Graz Technical University, 8010 Graz, Austria, Petergasse 16, 8010 Graz, Austria 4 Department of Chemistry of Polymeric Materials, Montanuniversität Leoben, Franz-Josef-Strasse 18, 8700 Leoben, Austria Corresponding author E-mail: [email protected] Abstract. In this contribution the deposition of Para-sexiphenyl (PSP) layers on poly (diphenyl bicyclo[2.2.1]hept-5-ene-2,3-dicarboxylate) (PPNB) by Hot Wall Epitaxy (HWE) is reported. It is demonstrated that pre-treating the substrate by UV-illumination induces a clear change in the morphology of the grown PSP films due to the polarity modification of the substrate surface. PPNB surface polarity increases when illuminated by UV via photo-Fries rearrangement. By detailed atomic force microscopy analysis the influence on the growth kinetics by the substrate temperature, deposition time and particularly by the UVtreatment of the substrate was investigated. A high crystalline order of the films is underlined by the observation of growth spirals and terraced islands, providing mono-layer step heights of standing PSP-molecules.

1.

Introduction

Morphology and crystalline order is determining for improving the electrical and optical properties of organic films. Therefore it is of great importance to study the growth kinetics of organic materials on substrate surfaces with well controlled properties. Para-sexiphenyl (PSP) (C36H26), a six units oligomer of para-phenylene, is a promising candidate as an electro active layer in organic LED displays due to its blue luminescence with high quantum yield. [1,2]. Moreover, it is classified as a wide gap organic semiconductor with an electronic band gap of 3.1 eV with photoluminescence in the blue visible range and polarized absorption and emission when provided in well ordered films [3]. The Hot-wall epitaxy (HWE) technique offers the advantage that it works close to thermodynamic equilibrium. Therefore it allows organic molecules to find the most suitable arrangement into the crystal lattice and as a consequence, highly ordered organic thin films can be obtained [3,4]. Organic thin films grown by HWE have shown outstanding optical and electrical properties [5,6]. Extensive morphological and structural characterization has been already performed on HWE grown

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PSP films deposited on various substrates such as KCl and muscovite mica, showing that the nature of the substrate and the growth conditions are ruling parameters for the molecular packing of the films [7]. In the presented work, an amorphous polymer poly(diphenyl bicyclo[2.2.1] hept-5-ene-2,3-dicarboxylate) (PPNB) containing photoreactive aryl-ester groups was chosen as substrate. Upon illumination with UV-light of λ < 280 nm these ester groups isomerise to the corresponding hydroxyketones in the so-called photo-Fries reaction [8,9]. Recently, we have shown that the photoreaction in PPNB yields up to 21% hydroxyketones as photoproducts [10]. The UV-induced reaction leads to a large increase of the refractive index as well as of the surface polarity. This enhanced polarity and the generated new functional groups can be used for selective post-modification reactions [11]. In the present work the influence of this change of surface polarity on the surface morphology of the PSP deposited layers is studied.

2.

Experimental Methods

The procedure to synthesise PPNB is reported by T. Höfler et al. [10]. For the substrate preparation a 10 mg/ml solution of PPNB in CHCl3 was prepared and stirred for 12 h. Then the solution was spin cast onto Si-substrates resulting in 80 nm film thickness. In order to provide equivalent growth conditions for UVexposed and non-irradiated surfaces, each substrate was divided in two regions. One half of the substrate was exposed for 20 min to UV light (254 nm) while the other half was protected from UV-illumination. The unfiltered light of an ozonefree mercury low pressure UV lamp (Heraeus Noblelight; 254 nm) was used. The illumination-process was done in inert gas atmosphere (nitrogen with a purity >99.95%) in order to avoid unwanted oxidation reactions. After the illumination process the substrates were transferred via a load lock to a HWE growth chamber working at a dynamic vacuum of 9 × 10–6 mbar. A 15 min in-situ preheating procedure was applied in order to reduce surface contaminations. The substrate temperature during preheating is chosen the same as the growth temperature in order to allow constant thermal conditions during the whole deposition process. The preheating process also removes possible adsorbed species from the surface of the substrate. A complete series of samples was prepared varying growth time (5–60 min.) and substrate temperature (100°C – 160°C) whereas the source- (240°C) and wall-temperature (260°C) were kept constant. The working principles of a HWE system can be found elsewhere [12]. Ex-situ atomic force microscopy (AFM) studies were performed using a Digital Instruments Dimension 3100 in the tapping mode. A SiC tip was used and the scanning area was 10 × 10 μm.

3.

Results and Discussion

PSP was deposited using the HWE technique on PPNB at different preparation conditions as explained in the experimental methods section. A detailed morphology analysis of the grown sample series, using AFM in tapping mode was performed.

Para-Sexiphenyl Layers Grown on Light Sensitive Polymer Substrates

25

Fig. 1. Atomic Force Microscopy images (10 μm × 10 μm) of Para-sexiphenyl on PPNB previously treated with UV light (upper half of each image) and as prepared (bottom half) for different substrate temperatures and different deposition times.

The influence of the polymer substrate on the surface properties of the grown PSP is analysed depending on substrate temperature, deposition time and UVillumination of the substrate before growth. Figure 1 depicts a chart with the AFM images for grown PSP films at, 100°C, 130°C and 160°C varying the deposition time from 5 to 60 min. The morphology of the film grown on the UV-illuminated side and on the as prepared surface are on the upper and bottom side of each image, respectively. The height scale (z0) is presented at the bottom of each image. A clear morphological difference between non-illuminated and UV-illuminated regions can be observed. Whereas on the non-illuminated side a homogeneous PSP film is formed, pre treating the substrate by UV illumination is leading to island formation and consequently to a change of the growth kinetic. This behaviour is a direct consequence of the increase in surface polarity from the substrate resulting of the photochemical reaction of PPNB described by T. Höfler et al. [10]. Changing the deposition time of PSP from 5 to 60 min is leading to an areaexpansion of the single islands. A more detailed comparison of the as prepared and UV-illuminated regions of the sample grown for 20 min at 160°C (Fig. 1-IIIb) is presented in Figure 2. On the one hand on the as prepared substrate (Fig. 2a, Fig. 1-IIIb bottom half) we can observe a closed film with spiral features which are characteristic of screw dislocations. On the other hand, on the UV illuminated side of the substrate (Fig. 2c, Fig. 1-IIIb upper half) crystallites with terraces of up to hundredths of nanometres in length can be observed, underling a change in the growth procedure.

26

G. Hernandez-Sosa et al.

Fig. 2. Atomic force microscopy scans and profiles of the surface structures formed by PSP deposited on the as prepared (a,b) and pre UV illuminated (c,d) PPNB surface. Scan size is 2 × 2 μm.

Figures 2b,d represent a detailed analysis of cross sections indicated in figures 2a,c as black solid lines. It is demonstrated that the step heights of the observed features are in good agreement with the value for one monolayer of standing PSP molecules corresponding to 2.6 nm. The formation and increment in size of these crystallites is induced by increasing the substrate temperature during PSP deposition, which is observed for both – illuminated and non-illuminated substrates – as shown in Figure 1. Similar results are found in literature for small organic molecules [13].

4.

Conclusions

High quality Para-Sexiphenyl films where successfully deposited on PPNB at different temperatures and deposition times. Careful AFM investigations show that PSP shows a different morphology and a different growth process when deposited on UV illuminated or as prepared PPNB substrates. This effect is attributed to the change in surface polarity after the UV illumination treatment and to the fact that the surface becomes more hydrophilic. Both changes are consequence of the photo-Fries rearrangement that the aryl-ester groups of PPNB undergo upon UV irradiation. On this work it is demonstrated that, the surface morphology of PSP grown on PPNB can be influenced by a simple and well controlled pre-treatment of the substrate surface and deposition conditions. This procedure could potentially influence the optical and electrical properties of the films deposited on the different parts of the substrate. Therefore the demonstrated sample structure opens new perspectives for the fabrication of devices which electrical and optical properties can be controlled by UV treatment of the substrate.

Para-Sexiphenyl Layers Grown on Light Sensitive Polymer Substrates

27

Acknowledgements. This work was supported by the Austrian Science Foundation projects NFN-S9702, NFN-S9706 and NFN-S9708. GHS wants to thank Consejo Nacional de Ciencia Tecnología (CONACYT), in México for scholarship.

References 1 2

3 4 5

6

7 8 9 10 11 12 13

S. Tasch, C. Brandstatter, F. Meghdadi, G. Leising, G. Froyer, L. Athouel, Advanced Materials, Vol. 9, 33, 1997. G. Leising, S. Tasch, W. Graupner, Fundamentals of Electroluminescence in Paraphenylene-type Conjugated Polymers and Oligomers, in Handbook of Conducting Polymers, Edited by T. Skothem, R. Elsenbaumer, J. Reynolds, 2nd ed., New York 1997. A.Y. Andreev, G. Matt, C.J. Brabec, H. Sitter, D. Badt, H. Seyringer, N.S. Sariciftci, Advanced Materials ,Vol. 12, 629, 2000. D. Stifter and H. Sitter, Applied Physics Letters, 66, 679, 1995. Th. B. Singh, N. Marjanovic, G.J. Matt, S. Gunes, N.S. Sariciftci, A.M. Ramil, A. Andreev, H. Sitter, R. Schwodiauer, and S. Bauer, Organic. Electronics. Vol. 6, 105, 2005. F. Quochi, F. Cordella, R. Orru, J.E. Communal, P. Verzeroli, A. Mura, and G. Bongiovanni A. Andreev, H. Sitter, and N.S. Sariciftci Applied Physics Letters, Vol. 84, 4454, 2004. T. Haber, A. Andreev, A. Thierry, H. Sitter, M. Oehzelt, R. Resel, Journal of Crystal Growth, Vol. 284, 209, 2005. J.C. Anderson, C.B. Reese, Proceedings of the Chemical. Society, 217, 1960. D. Bellus, Advances in Photochemistry, Vol. 8, 109, 1981. T. Höfler, T. Griesser, X. Gstrein, G. Trimmel, G. Jakopic, W. Kern, Polymer, Vol. 48, 1930, 2007. T. Griesser, T. Höfler, S. Temmel, W. Kern, G. Trimmel, Chemistry of Materials, Vol. 19, 3011, 2007 A. Lopez-Otero, Thin Solid Films, Vol. 3, 4, 1978. Th.B. Singh, N.S. Sariciftci, H. Yang, L. Yang, B. Plochberger and H. Sitter, Applied Physics Letters, Vol. 90, 213512, 2007.

Thermal Desorption of Organic Molecules Adolf Winkler Institute of Solid State Physics, Graz University of Technology, Petersgasse 16, A-8010 Graz, Austria E-mail: [email protected] Abstract. The applicability of thermal desorption spectroscopy (TDS) to differentiate between a strongly bound wetting layer and the subsequently formed multilayers of organic thin film is discussed. After describing the fundamentals of TDS, particularly for large organic molecules, several model systems are presented (p-4P on Au(111), p-6P on Au(111), mica(0001) and KCl(001)), demonstrating the importance of the substrate material, surface structure and surface contaminations on the formation of a wetting layer. The wetting layer acting as a template for the multilayer growth strongly influences the structure and morphology of the organic thin film.

1.

Introduction

Thermal desorption spectroscopy (TDS) is a powerful and well known experimental technique to investigate adsorption/desorption kinetics and ener-getics of small (inorganic) molecules on surfaces [1,2]. The application of this technique for large organic molecules and ultra-thin organic films is not that widely acknowledged. In this review I will demonstrate that TDS can be successfully used to obtain information on a number of parameters characterizing organic films in the sub-monolayer, monolayer and near multilayer regime. In particular, the frequently put question as to the existence of a wetting layer prior to the growth of island like films can be answered unambiguously. In addition, the binding energy of the first monolayer and the heat of evaporation of the multilayer can be determined. Other kinetic parameters of adsorption and desorption, like the sticking coefficient and the pre-exponential factor for desorption can be obtained as well. In this contribution I will exemplify the power of TDS on some model systems: p-quaterphenyl (p-4P) on Au(111) and p-hexaphenyl (p-6P) on Au(111), mica(0001) and KCl(001). I will particularly show that the chemical composition (carbon contamination) and the morphology (surface roughening by sputtering) of the substrate surface can have a tremendous effect on the thin film growth.

2.

Fundamentals of TDS

For thermal desorption spectroscopy (TDS) the material is first put onto the surface and thereafter the sample is heated with a constant heating rate until the material is desorbed again from the surface. The desorbing particles are detected with a mass spectrometer and the signal vs. time (or temperature) is called the TD H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_5, © Springer-Verlag Berlin Heidelberg 2009

29

30

A. Winkler

spectrum. In most cases the desorption spectrum can be successfully described by the Polanyi-Wigner equation:

dN dN x R =− =− β = ν ⋅ N ⋅ exp( − E (Θ)/ kT ) des des dt dT

(1)

here is R des : desorption rate, β: heating rate, ν: pre-exponential factor, N: number of molecules per square unit, x: desorption order, Edes: desorption energy, Θ: coverage, T: surface temperature, k: Boltzmann’s constant. The coverage Θ is usually defined (for small inorganic molecules and atoms) as the ratio of the number of adsorbed particles to the number of surface atoms per square unit. It is given in monolayers (ML), where the saturation coverage is typically 1 ML. For large organic molecules the value of the saturation coverage would be much smaller than unity. In this case the coverage is frequently defined as the ratio of the number of adsorbed molecules to the maximum number of adsorbed molecules per square unit (Nmax), which can be regarded as one physical monolayer. Then N = Nmax·Θ. The desorption order x describes the coverage dependence of the desorption rate. For organic molecules only the desorption orders x = 1 (first order) and x = 0 (zero order) are relevant. Zero order desorption takes place in the presence of multilayers. In this case the maximum number of molecules per surface unit is always available and therefore no coverage dependence exists (N0). When it comes to desorption of the last monolayer the desorption rate is proportional to the number of available molecules, i.e. proportional to N1. The desorption energy Edes is constant for multilayer desorption and is equal to the heat of evaporation. For the desorption of the final monolayer the desorption energy may be coverage dependent: E(Θ) = E0 ± ωΘ. The ± sign describes an attractive (+) or repulsive (−) lateral interaction between the molecules, with the interaction energy ω. The pre-exponential factor ν can be correlated with the attempt frequency of the adsorbed molecules to overcome the adsorption potential. For atoms and small molecules this value is typically in the order of 1013 s–1. For molecules with a large number of atoms this perception is not appropriate. The pre-exponential factor actually takes into account the change of all translational and internal degrees of freedom during desorption. As a result of transition state theory (TST) considerations [3] the pre-exponential factor can be described as:

⎛ kT ⎞ q⊕ ⎟ ⎝ h ⎠ q

ν =⎜

(2)

with h: Planck’s constant, q: partition function of the adsorbed state, q ⊕: partition function of the desorbed (free) state. For atoms and small molecules both partition functions are similar and therefore ν ≈ kT/h ≈ 1013 s–1, for typical desorption temperatures. For large molecules, however, the partition function of the free molecules is due to the many rotational and vibrational degrees of freedom much larger than for the adsorbed molecule, where only frustrated rotations and vibrations exist. Therefore ν is in this case typically by many orders of magnitude larger than 1013 s–1 [4,5].

Thermal Desorption of Organic Molecules

31

The evaluation of the desorption energy for multilayer desorption, i.e. the heat of evaporation, is straightforward. The plot of ln Rdes vs. 1/T for the leading edge of the spectrum yields a straight line, where the slope is equal to –Edes/k. If the desorption rate is known quantitatively then the intercept of this straight line with the Y-axis yields the pre-exponential factor. For first order desorption a characteristic feature of the spectra is the coverage independence of the peak maxima. According to Redhead [6] the correlation between the desorption energy Edes and the peak maximum Tm is given by:

Edes ⎛ ν =⎜ 2 kTm ⎝ β

⎞ ⎟ exp (− Edes / kTm ) ⎠

(3)

An approximate solution of this equation is:

⎛ ⎛νT ⎞ ⎞ Edes ≈ kTm ⎜ ln ⎜ m ⎟ − 3.64 ⎟ ⎝ ⎝ β ⎠ ⎠

(4)

For many inorganic adsorbates, where ν ≈ 1013 s–1, the following numerical equation holds (Redhead equation): Edes (cal / mol ) ≈ 60 ⋅ Tm ( K )

(5)

however, this approximation cannot be used for large organic molecules. A better approach is to determine first the pre-exponential factor ν for multilayer desorption and to take this value then for the desorption of the monolayer. For example, p-6P multilayer desorption from Au(111) yields a pre-exponential factor ν = 5.6 × 1025 s–1 [7]. This leads to a numerical equation (5) which contains a value of 124 instead of 60! This is one of the reasons why the extraction of desorption energies from TD spectra via the Redhead formula was sometimes performed incorrectly [8]. The shape of the desorption spectra and the shift of the peak maxima Tm with changing coverage also gives some valuable information. Zero order desorption is characterized by a sharp cutoff at the trailing edge when the coverage goes to zero (Fig. 1a). The peak maxima shift to higher temperatures with increasing coverage. However, this shape is never observed for real situations, because desorption of the last monolayer will always change to a first order reaction, which leads to a less sharp trailing edge (Fig. 1b). First order desorption is characterized by asymmetric peaks but with coverage independent peak maxima Tm (Fig. 1c). If the desorption energy is coverage dependent then a shift of the peak maxima as a function of coverage appears. In the case of repulsive interaction, as frequently observed in the sub-monolayer range for organic molecules, Tm shifts to lower temperature (Fig. 1d), whereas for attractive interaction a shift to higher temperature is observed with increasing coverage.

A. Winkler

Desorption rate / arbitary units

4.5 4

3.5

a

Desorption rate / arbitary units

32

3.5 3

2.5 2

1.5 1

0.5

Desorption rate / arbitary units

3.5 3

320

340

360

380

Temperature / K

2.5 2

1.5 1

0.5 0

250

2 1.5 1 0.5

2

c

300

350

Temperature / K

400

b

2.5

0

400

Desorption rate / arbitary units

0

3

320

340

360

380

Temperature / K

400

d

1.5

1

0.5

0 250

300

350

Temperature / K

400

Fig. 1. Calculated desorption spectra for pure zero-order desorption (a), for changing the desorption order from zero to first order at 1 monolayer (b), for pure first order desorption (c) and for first order desorption with repulsive lateral interaction (d).

3.

Experimental Technique

The technique of TDS is quite simple and straightforward. The molecules which have first been put onto the surface at sufficiently low surface temperature by gas dosing or evaporation are afterwards desorbed again by heating of the sample. Typically, linear heating rates between 1 to 10 K/s are applied. The desorbed material can be detected by a pressure gauge or by a mass spectrometer. Since the amount of material detected is quite small, TDS is generally performed under ultra-high vacuum conditions. According to Redhead [6] the desorption rate is correlated with the detected signal in the following way: Rdes = KSp / A + KVdp / dt

(6)

with S: pumping speed, p: pressure signal, V: volume of the vacuum chamber, A: surface area, K = 3.27 × 1019 molecules/l. In the case of very large pumping speed S the second term in Equ. 6 can be neglected and the desorption rate becomes directly proportional to the measured signal. Whereas this condition for small (volatile) gas molecules is not necessarily fulfilled in all cases, for condensable species like large organic molecules this is fulfilled to a great extent. In other words, condensable desorbing molecules do not

Thermal Desorption of Organic Molecules

33

contribute to an increased background pressure, but are immediately pumped away when hitting the surface of the vacuum chamber at room temperature (S→∞). But this also means that the particles have to be detected by an in-line-of-sight detector. Multiplexing of the mass spectrometer allows also to detect possible reaction products of the adsorbate. For example, the dehydrogenation of a p-6P monolayer on gold surfaces could be verified by this method [9]. On the other hand cracking of the molecules in the ionization region of the mass spectrometer has to be taken into account. In particular for large reactive organic molecules, e.g. for thiol based molecules, which are typically used for SAM preparation, the discrimination between the cracking of the molecule at the surface and in the mass spectrometer can be a challenging task [10,11]. Thermal desorption spectroscopy is a very sensitive method regarding the kinetics and energetics of the adsorbate, which in turn depends strongly on the surface conditions. Therefore, a comprehensive characterization of the surface prior to TDS is indispensable. Typically, Auger electron spectroscopy (AES) or X-ray photoelectron spectroscopy (XPS) are applied for the determination of the surface chemical composition and Low Energy Electron Diffraction (LEED) for the geometrical characterization.

4.

Wetting Layers Observed by TDS

TDS is particularly suited to differentiate between a strongly bound wetting layer of organic molecules and the more weakly bound molecules in the multilayer. In most cases rod or disk like molecules will be adsorbed in a flat-on configuration in the wetting layer. It turns out that the structural configuration of the wetting layer generally determines the structure and morphology of the subsequently formed multilayer [12,13]. The wetting layer depends strongly on the substrate material, the substrate structure and contaminations on the substrate surface. In Fig. 2 desorption spectra of p-6P on mica(0001) and KCl(001) are shown. On mica (Fig. 2a) the existence of a strongly bound wetting layer is clearly observed (ß-peak), which saturates at about 0.2 nm mean thickness [14]. With increasing coverage the less strongly bond α peak appears which does not saturate, which is representative for multilayer desorption. Thus, the needle like islands observed for this system grow on top of the wetting layer. In contrast, on the potassium chloride surface (Fig. 2b) no evidence for a strongly bound first layer can be seen [15]. Even for the smallest coverage of 0.1 nm the desorption peak is located at the leading edge of the multilayer peaks. Since it is known that for small coverage the film grows in the form of needle like islands [16,17] one can conclude that no wetting layer exists between the islands. The influence of carbon contamination on the p-6P layer growth on Au(111) is shown in Fig. 3. On the clean surface several strongly bound adsorption states (ß1, ß2, ß3) can be identified for small coverage (Fig. 3a). The ß1 and ß2 states have been explained to be due to flat lying and side tilted molecules with coverage of 0.5 monolayers each. The ß3 peak stems from a second layer which can still be energetically distinguished from the multilayer peaks denoted by α (not shown here) [7]. The contamination of the surface by carbon (e.g. by X-ray irradiation of a thick p-6P film) leads to a significant change of the wetting layer.

34

A. Winkler

Fig. 2. (a) TDS for p-6P from mica(0001) showing the wetting layer (ß-peak) and multilayer (α-peak) (After Frank et al. [14]), (b) TDS for p-6P from KCl(001) reveals that no wetting layer exists in this case (After Frank et al. [15]).

Only one rather weakly bound desorption state (β) shows up in this case before multilayer desorption starts (Fig. 3b). Also the second layer cannot be distinguished any longer from the multilayer peaks. The morphology of thicker films grown on these two substrates are totally different. Whereas in the former case needle like island are formed consisting of molecules oriented parallel to the surface, in the latter case mound like island consisting of standing molecules can be observed [18]. Similarly, on mica(0001) C-contamination or roughening of the surface by heavy sputtering leads to a removal of the wetting layer. This results in a film growth with mound like islands consisting of standing molecules [14].

Fig. 3. (a) TDS of p-6P from clean Au(111) shows three states (ß1-ß3) belonging to the wetting layer, (b) TDS of p-6P from carbon contaminated Au(111) shows only one state of a more weakly bound wetting layer (After Müllegger and Winkler [7]).

Thermal Desorption of Organic Molecules

5.

35

Determination of Desorption Energies for Wetting and Multilayers

The desorption energy of multilayers (heat of evaporation) can be easily obtained as outlined in section 2. As an example we show the multilayer desorption of p-4P from a carbon covered polycrystalline Au-foil (Fig. 4a) and the corresponding lnR vs. 1/T plot (Fig. 4b). In Fig. 4a one can clearly see the zero order desorption behavior as described above (common leading edge, sharp cutoff of trailing edge). The determination of the desorption energy from Fig. 4b yields Edes = 1.6 eV. For the pre-exponential factor one obtains ν = 5 × 1021 s–1 [19]. 4P/0.5ML C/Au(poly) 300 k

TDS - signal (arb. units)

α

film thickness 12 nm 6.0 nm 2.5 nm 1.0 nm

34

In(Rate)

a

p-4P on Au(poly) multilayer

b

32

30

m = 306

350

400 450 Temprature (K)

0.0025

500

0.0026

0.0027

0.0028

1/T

Fig. 4 (a) Series of TDS for p-4P from a polycrystalline gold surface as a function of mean thickness, (b) lnR vs. 1/T plot for determination of the desorption energy (After Müllegger et al. [19]).

The desorption energies for the monolayer states (wetting layer) can be calculated by using Equ. 4 and inserting the pre-exponential factor obtained from the multilayer peak. In Table 1 the desorption energies and pre-exponential factors for various adsorption systems are compiled. Table 1. Compilation of pre-exponential factors and desorption energies for wetting layers and multilayers of p-4P and p-6P on various substrates.

System

ν (multi) [s–1]

Edes(multi) [eV]

Edes(mono)-β2 [eV]

Edes(mono)-β1 [eV]

p-4P/Au(poly) p-4P/Au(111) p-6P/Au(111) p-6P/Mica(0001) p-6P/KCl(001)

5 × 1021 1.6 × 1021 5.6 × 1025 3.7 × 1025 3 × 1024

1.6 1.5 2.4 2.5 2.3

1.9 2.1 3.2 – –

2.9 2.6 3.6 2.9 –

36

A. Winkler

6.

Conclusion

Thermal desorption spectroscopy is a powerful method to identify and characterize wetting layers in ultra-thin organic films. Furthermore, TDS allows to determine energetic (desorption energy) and kinetic (pre-exponential factor, sticking coefficient) parameters for the individual adsorption systems. In combination with other surface analytical tools a quite comprehensive characterization of the initial stages of organic film growth can be obtained. This has been exemplified for oligophenylenes (p-4P, p-6P) on various substrates (Au(111), mica(0001) and KCl(001)). Acknowledgements. I would like to thank S. Müllegger and P. Frank for their valuable contributions to this work. The financial support by the Austrian Science Fund (FWF) is gratefully acknowledged.

References 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19

D. A. King, Surface Sci. 47, 384, 1975. A. M. de Jong, J. W. Niemantsverdriet, Surface Sci. 233, 355, 1990. V. P. Zhdanov, Surf. Sci. Rep. 12, 183, 1991. K. R. Paserba and A. J. Gellmann, Phys. Rev. Lett. 86, 4338, 2001. K. A. Fichthorn and R. A. Miron, Phys. Rev. Lett. 89, 196103, 2002. P. A. Redhead, Vacuum 12, 203, 1962. S. Müllegger and A. Winkler, Surface Sci. 600, 1290, 2006. C. B. France and B. A. Parkinson, Appl. Phys. Lett. 82, 1194, 2003. S. Müllegger and A. Winkler, Surface Sci. 600, 3982, 2006. J. Stettner, P. Frank, T. Griesser, G. Trimmel, R. Schennach, R. Resel, and A. Winkler, in this proceedings. P. Frank, J. Stettner, F. Nußbacher, and A. Winkler, in this proceedings. R. Resel, M. Oehzelt, T. Haber, G. Hlawacek, C. Teichert, S. Müllegger, and A. Winkler, J. Cryst. Growth 283, 397 2005. T. Haber, S. Müllegger, A. Winkler, and R. Resel, Phys. Rev. B74, 045419, 2006. P. Frank, G. Hlawacek, O. Lengyel, A. Satka, C. Teichert, R. Resel, and A. Winkler, Surface Sci. 601, 2152, 2007. P. Frank, G. Hernandez-Sosa, H. Sitter, and A. Winkler, Thin Solid Films 516, 2939, 2008. F. Balzer and H. G. Rubahn, Surface Sci. 548, 170, 2004. T. Haber, A. Andreev, A. Thierry, H. Sitter, M. Oehzelt, and R. Resel, J. Cryst. Growth 284, 209, 2005. S. Müllegger, G. Hlawacek, T. Haber, P. Frank, C. Teichert, R. Resel, and A. Winkler, Appl. Phys. A 87, 103, 2007. S. Müllegger, O. Stranik, E. Zojer, and A. Winkler, Appl. Surf. Sci. 221, 184, 2004.

Crystalline Stages of Rubrene Films Probed by Raman Spectroscopy B.A. Paez1, Sh. Abd-Al-Baqi2, G.H. Sosa2, A. Andreev1, C. Winder1, F. Padinger1, C. Simbrunner2 and H. Sitter 2 1

NANOIDENT Technologies AG, Untere Donaulände 21-25, A-4020 Linz, Austria E-mail: [email protected] 2 Institute of Semiconductor and Solid State Physics, University Linz, Altenbergerstr. 69, A-4040 Linz, Austria E-mail: [email protected] Abstract. We report on ex situ Raman characterization of rubrene thin films grown by hotwall epitaxy on cleaved mica substrates. Analysis of the vibrational bands revealed that at earliest growth stages the film is amorphous. In particular, a broad band at 1373 cm–1 proves the amorphous nature of the film. The rubrene molecules in amorphous phase are geometrically distorted, since the appearance of the Raman band at 1606 cm–1 is only infrared active for rubrene molecules with the C2h symmetry group. Further growth leads to seeding of spherulites in the amorphous matrix and further to their coalescence. Raman bands from isolated spherulites embedded in an amorphous matrix and from coalesced spherulites show polarization dependence (depolarization ratio < 0.6), thus demonstrating their crystalline nature. It is also found that the breathing mode (1003 cm–1) represents the rubrene fingerprint feature independent of layer crystallinity.

1.

Introduction

Understanding and engineering of molecular crystals is of great interest to achieve a tuned performance of organic-based devices. Rubrene, a tetraphenyl derivative of tetracene, has recently attracted much attention since hole mobili-ties as high as 20 cm2 V–1 s–1 for single crystals at room temperature have been reported. It is also found that the limited cofacial π-stack interactions result in rubrene in very efficient electronic coupling, which is consistent with the band regime at room temperature [1]. Raman spectroscopy plays a very important role in determining vibrational properties of the matter. The usefulness of the method is mainly due to its sensitivity and non-destructiveness, capability to provide information on chemical identity, charge states [2,3], processes at interfaces [4] and structural order [5]. Investigations of rubrene single crystals have been reported elsewhere [6]. Although similar structures as those reported by us have been used in organic field effect transistors [7], Raman investigations on thin films are still missing up to now.

H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_6, © Springer-Verlag Berlin Heidelberg 2009

37

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B.A. Paez et al.

2.

Experimental Methods

The rubrene layers were deposited by hot-wall epitaxy [8] on cleaved mica substrates. Raman investigations were carried out ex situ in a Labram Aramis spectrometer from Jobin Yvon HORIBA. The structures were analyzed in the back scattering configuration and excited with the 784 nm (1.57 eV) light from a laser diode. Throughout this report the Porto coding for light polarization is adopted. It takes into account the propagation of the incoming and scattered beams together with the polarization, i.e., z(xy)z’: means the incident light is propagating along the z axis and then back scattered (z’); xy in the brackets labels the electric field polarizations with respect to a fixed reference plane holding the sample.

3.

Results and Discussion

3.1. Surface Morphology Three growth stages were selected for the vibrational investigations. The first one was homogenous rubrene layer (further called as amorphous matrix-it is explained in section 3.3), the second were isolated spherulites embedded in amorphous matrix, and finally coalesced spherulites forming a closed rubrene layer. Corresponding optical micrographs are shown in Figure 1. More details can be found in Ref. [9]. Fig. 1. Morphology of rubrene films in different growth stages: a) isolated spherulites in amorphous matrix; b) closed rubrene layer in coalescence stage. Raman measurements points: 1,2 – the center and close to border of an isolated or coalesced spherulite; 3 – amorphous matrix.

3.2. Vibrational Identity Rubrene films are first characterized by the breathing mode at 1003 cm–1. This mode proves the presence of the rubrene molecules on the mica substrate, i.e. it represents a fingerprint for their presence on a surface. The corresponding Raman spectra for the spherulites in the amorphous matrix and for the closed rubrene layer are shown in Figure 2.

Crystalline Stages of Rubrene Films Probed by Raman Spectroscopy 39 Fig. 2. Raman spectra (breathing mode) for isolated spherulites in amorphous matrix (top), and for the closed rubrene layer (bottom). Numbering as follows: (1)-spectra from the center, (2)-close to the border of the spherulite, (3)from amorphous matrix. Number 2 is also used for the closed layer, where all spherulites are coalesced. Rubrene powder and mica substrate spectra are also shown as reference.

Breathing (Ag) 1003 cm-1 λexc = 785 nm

50

z (xy) z ˙

cts, s-1

x3.6 x3.6 x2.4

z (yy) z ˙ z (xy) z ˙ z (yy) z ˙

x6.4

z (xu) z ˙

x6.6

z (xy) z ˙

x0.4 x1.0 x0.7 x1.0

z (yy) z ˙ z (xy) z ˙ z (yy) z ˙ z (xu) z ˙

50

x1.6 x1.0

Rubrene Powder Mica substrate

1000

2. 3. 1. 1. 2. 2.

cts, s-1

z (yy) z ˙

950

1. 1. 2.

Raman shift / cm-1

2. 2.

1050

The spectra were recorded at different laser light polarizations. In order to skip strong differences on the Raman intensity, all the spectra are normalized and the multiplication factors are indicated. It is found that the profile of the breathing mode is quite similar for amorphous matrix, isolated spherulites in amorphous matrix, and for closed rubrene layer, excepting the higher signal to noise ratio in the amorphous phase. 3.3. Amorphous Matrix and Crystallinity: Raman Investigations The use of polarized light in Raman spectroscopy allows to identify amorphous phases, crystallinity, isotropy, and in general to test selection rules of the investigated structure. The Raman spectra measured from the rubrene matrix material were found to be independent of light polarization and also the scattered light was left unpolarized (z(yu)z’). That means no preferred orientation exist in this part of the rubrene layer. Therefore we called it the amorphous matrix. In Figure 3, the Raman spectra of the spherulites in amorphous matrix and of the closed layer are shown. Although, the Raman spectrum of the amorphous phase lacks of several vibrational features like those observed in rubrene powder, there are clear signatures, which proves the presence of this phase in rubrene films. As it was already mentioned, the first one (but not exclusive) is the isotropy of the Raman spectrum upon light polarization. Second one is the broad band at 1373 cm–1 with a small shoulder at 1356 cm–1 (both labeled AM*). Namely, the both bands are clearly seen for amorphous phase, but completely missing for rubrene powder and closed rubrene layers. Isolated spherulites present intermediate

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case maybe due to the quite high penetration depth of the laser light in rubrene. Another characteristic of amorphous phase is the molecular symmetry breakdown proved by the broad band at 1606 cm–1. The mode is only infrared active according to the molecular symmetry group of isolated rubrene (C2h). The seeded in the amorphous matrix spherulites preserve most of the bands of rubrene powder (Figure 3). The Raman spectra, recorded both for the center and at the boundary of a spherulite, reveal strong polarization dependences, for example, for the bands depicted by the phenyl groups (1303 cm–1), and tetracene backbone (1522 cm–1). The symmetry of the Ag mode for the phenyl groups and Bg mode for the tetracene core turns from the coupling between the electric field and the dipole transition in the rubrene structure. AM˙ Phenyl groups 1303 cm-1 (Ag)1373 cm-1 λexc=785 nm

z (xy) z˙

Tetracene core AM˙ 1522 cm-1(Bg)1606 cm-1(Ag) 102 cts. s-1

X 8.0

z (yy) z˙

1.

z (xy) z˙

1. 2.

z (yy) z˙

Amorphous

z (xu) z˙

2. 3.

z (xy) z˙

102 cts. s-1

X 1.0

z (yy) z˙

1. 1.

z (xy) z˙ z (yy) z˙

2.

z (xy) z˙

2.

z (yy) z˙

2. 2.

1200

1300

1400

1500

Raman shift / cm-1

1600

Rubrene powder Mica substrate

Fig. 3. Raman spectra of rubrene films: spherulites in amorphous matrix (top), and from closed rubrene layer (bottom); 1,2 – the center and close to border of a spherulite; 3 – amorphous matrix. Rubrene powder and mica substrate spectra are also shown as reference. AM* - labels the signatures of the amorphous matrix.

The coalesced spherulites forming closed rubrene layer are also sensitive to the polarized light. Depolarization ratio between different polarized spectra (below 0.6) indicates the crystalline nature of the rubrene in this stage. A common feature of many spectra was the sloped background, despite of the fact that low excitation energy (1.57 eV) was used for measurements. It comes from the high photoluminescence yield of the rubrene [10]. In particular, the

Crystalline Stages of Rubrene Films Probed by Raman Spectroscopy 41

background of the spherulites is sensitive to the light polarization, which indicates selection rules for the luminescence (Fig. 3).

4.

Conclusions

Raman investigations proved that rubrene layers grown by hot wall epitaxy have different structural properties depending on the growth stage. Rubrene films independent on their amorphous or crystalline nature were distinguished by the breathing mode at 1003 cm–1 (Ag symmetry), i.e. this mode represents the rubrene fingerprint feature. Analysis of the vibrational bands revealed that at earliest growth stages the film is amorphous. In particular, a broad band at 1373 cm–1 proves the amorphous nature of the film. On the other hand, the mode at 1606 cm–1, usually an infrared active band, proves a symmetry breakdown of the molecule at this growth stage. Additionally, the amorphous phase lacks of vibrational activity at the phenyl groups, and tetracene backbone. Therefore, it is likely that the geometry of the rubrene molecule is dramatically distorted. Further growth leads to seeding of spherulites in the amorphous matrix and further to their coalescence. This growth phase associated with spherulite-like shapes (spherulites in amorphous matrix and coalesced spherulites) was found to be highly sensitive to the light polarization, which is shown by the phenyl group band (1303 cm–1) and the tetracene core band (1522 cm–1). That proves clearly the crystalline nature of the rubrene in the spherulites.

References 1

D. Beljonne, et al., in “Handbook of conducting polymers”, Conjugated polymers: theory, synthesis properties and characterization, chap. 1, 3rd ed. Edited by T. A. Skotheim and J. R. Reynolds (CRC Press, 2007). 2 B. A. Paez, et al., Proc. SPIE Int. Soc. Opt. Eng. 5217, 63 (2003). 3 M. L. Shand, W. Richter, E. Burstein, and J. G. Gay, J. Nonmmetals 1, 53–62 (1972). 4 B. A. Paez, et al., Appl. Surf. Sci. 234, 168 (2004). 5 P. Colomban, Spectroscopy Europe 15 (6), 8 (2003). 6 J. R. W.-Wolf, L. E. McNeil, S. Liu and C. Kloc, J. Phys.: Condens. Matter 19 (2007) 276204 (15 pp). 7 S.-W. Park, et al., Appl. Phys. Lett. 90, 153512 (2007). 8 H. Sitter, chap. 5, in Organic Nanostructures for Next Generation Devices. Edited by K. Al-Shamery, H.-G. Rubahn, and H. Sitter (Springer-Verlag Berlin Heidelberg 2008). 9 Sh. Abd al-Baqi, et al., Proceedings Symposium O, E-MRS Spring Meeting 2008, Springer “Proceedings in Physics”, in print. 10 S. 8, Z. Peng, X. Zhang, S. Wu, Journal of Luminescence 121, 568–572 (2006).

Rubrene Thin Film Characteristics on Mica Sh. M. Abd Al-Baqi1, G. Henandez-Sosa1, H. Sitter1, B. Th. Singh 2, Ph. Stadler 2 and N. S. Sariciftci 2 1

Institute of Semiconductor and Solid State Physics, Johannes Kepler University, Linz, Austria 2 Institute of Physical Chemistry and Linz Institute For Organic Solar Cells (LIOS), Johannes Kepler University, Linz, Austria E-mail: [email protected]

Abstract. Rubrene thin films were deposited by Hot Wall Epitaxy on mica substrates.

To optimise the growth conditions, the growth rate and the substrate temperature were changed systematically. The surface morphology of the grown rubrene layers was investigated by polarized optical microscopy (POM), electron microscopy (SEM) and atomic force microscopy (AFM). After an initial nucleation and coalescence stage a continuous amorphous layer is formed. In a later stage of growth, spherulites embedded in the amorphous matrix are found, which furthermore cover the whole surface. It could be proven that the spherulite consist of polycrystalline material, which could be used for the fabrication of organic field effect transistors.

1.

Introduction

Many attempts were made to fabricate organic field effect transistors from rubrene thin films [1–4]. The largest mobility obtained so far in rubrene OFETs is 2.5 cm2. V–1. s–1, which is still much less than in monocrystlline bulk material [5]. The main difference between bulk and thin film material is the crystalline property. Consequently, the main effort goes in the direction of improving the crystallographic order in the rubrene layers. Since there is no lattice matched substrate available, the only chance to approach the goal of highly ordered structures is to use an optimised growth regime. Due to the weak Van der Wales type bonds of the molecules, a deposition process as close as possible to thermodynamic equilibrium seams to be appropriate. The local dynamic equilibrium at the growing surface allows a manifold impingement and re-evaporation of the molecules. Supported by an enhanced surface mobility the rubrene molecules can find the optimum position on the substrate to form a highly ordered structure, which can lead to a layer of enhanced crystallinity. The method of choice to provide such growth conditions is the Hot Wall Epitaxy (HWE).

2.

Experimental Procedure

The rubrene source material was purchased from Aldrich (purity > 98%) and purified by threefold sublimation under dynamic vacuum conditions. The rubrene layers were evaporated in a standard HWE reactor on freshly cleaved 2M1 muscovite mica. Two different growth rates were used by employing different

H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_7, © Springer-Verlag Berlin Heidelberg 2009

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source temperatures (Ts) for the evaporation of rubrene (Ts = 180°C and 235°C). The surface mobility of the rubrene molecules was influenced by the substrate temperatures (Tsub = 80°C, 90°C or 120°C). The surface structure of the rubrene layers was investigated routinely by polarized optical microscopy. Atomic Force Microscopy (AFM), (VEECO Di3100) in tapping mode was employed. In the case of high surface roughness SEM was used to investigate the surface morphology. The crystalline property of the rubrene layers was probed by polarization depend Raman spectroscopy and x-ray diffraction experiments reported in detail elsewhere [6,7].

3.

Results and Discussion

Since mica substrates can be freshly cleaved before evaporation and therefore provide a very clean surface, we selected this substrates to investigate the growth process of rubrene layers. The cleaving process was done ambient air, which means that some dust particles can still contaminate the substrate surface. In the early stage of growth single rubrene islands are formed which coalesce and form an amorphous layer [8]. The later stage of growth is dominated by the formation of spherulites. Figure 1 summarizes the optical micrographs to demonstrate the surface structure after different growth times. Very similar rubrene surface structures were observed previously on SiO2 substrates [9]. The development of spontaneously nucleated spherulites is clearly seen, which are embedded in an amorphous matrix. By increasing the deposition time, the spherulites grow in diameter and consist of an inner disc surrounded by an outer ring. Finally the spherulites coalesce and cover with the outer ring structures the whole substrate. Depending on their growth mechanisms spherulites can be divided into two categories nucleated either from single nucleation site or bunch of needles and both categories have been found in these thin films from TEM measurement [10]. The radii of the central disc and the outer rings are plotted as a function of growth time in Fig. 2, showing a linear increase with the onset of a saturation due to coalescence of the spherulites. The nucleation of the spherulites can be caused by static impurities or dynamic heterogeneities [10]. If the pin holes in the amorphous matrix would act as nucleation centers, as assumed in ref. [9], the density of spherulites should be the same as the density of the pinholes, which is by far not the case. On the other hand, if the spherulites are originated by defects in the amorphous larger, the spherulites would grow on top of the amorphous matrix which is also not the case. As shown by the AFM picture in Fig. 3, the spherulites are embedded in the amorphous matrix. The cross section across the border of the spherulite (see Fig. 3) shows the same height for the spherulites as for the amorphous matrix and a clear trench at the borderline. So we assume that starting from an impurity the amorphous matrix undergoes a phase change by forming poly crystalline structures. In that way the material of the amorphous matrix is consumed until the whole substrate is covered with spherulites.

Rubrene Thin Film Characteristics on Mica

45

Rubrene/Mica

Ts = 235°C Twall = 235°C Tsub = 80°C Growth time = 30 min Growth time = 1 hour Growth time = 3 hours

2 0 0 µm

5 0 0 µm

2 0 0 µm

Fig. 1. Optical Micrographs of the surface structure with increasing growth time. 1000 900

Radius (µm)

800 700 600 500 400 300 200

Inner Disc Outer Disc

100 0 0

100

200

300

t (minutes)

Fig. 2. Diameter of the inner disc and outer rings as a function of growth time, Ts = 235°C, Twall = 235°C, Tsub = 90°C.

trench

spherulite

amorphous

10µm

Fig. 3. AFM image of rubrene on mica together with a cross section across the border.

Optical microscopy using polarized light gives a first hint on the crystallinity of the obtained structures. A typical result of optical micrographs obtained without polarization, parallel and perpendicular orientation of the polarizer and analyzer are shown in the Fig. 4. If the relative position of the polarizer and analyser is changed; the different regions of the spherulites change their colour, while the

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surrounding amorphous matrix stays unchanged. This can be interpreted, that the spherulites consist of small crystallites oriented in a radial direction. A detailed investigation of the crystallographic order inside the spherulite was performed by polarization dependent micro Raman spectroscopy and x-ray diffraction [6,7]. Due to the higher surface roughness inside the spherulite the morphological details were studied by SEM. Figure 5a shows the typical structure of the inner disc of a spherulite. Zooming into the central part two different features can be observed. Out of a layer consisting of similar elongated mosaic blocks, (Fig. 5b) whisker like facetted structures grow in the third dimension (Fig. 5c). During the growth of the spherulites in lateral direction by recrystallization of the amorphous material, the flux of impinging rubrene molecules continuous. The additional molecules hitting the amorphous matrix contribute to the continuous growth of this part of the layer. The other molecules impinging on the polycrystalline spherulites find crystalline mosaic blocks as seeds for the formation of whiskers, pointing out of the surface in the central part of the sphurlites. Figure 4c clearly shows the faceted structure of such whiskers. Rubrene/Mica Tsource = 235°C Twall = 235°C Tsub = 80°C growth time = 1 hour Without polariztion 180° 90°

200µm

200µm

200µm

Fig. 4. Rubrene thin film on mica substrate under polarized optical microscope.

200 nm

1 μm

200 nm

Fig. 5. SEM pictures for rubrene on mica substrate Ts = 230°C, Twall = 240°C, Tsub = 120°C, Tpre-heat = 90°C.

4.

Conclusions

Rubrene thin films were evaporated by HWE on mica. After the formation of an amorphous matrix, spherulite structures start to grow, which finally cover the whole surface. Due to their polycrystalline property the spherulites are resistant against oxidization while the amorphous matrix becomes transparent upon exposure to air, which is a clear sign for oxidation. The rubrene spherulites provide therefore a promising material for the fabrication of OFETs.

Rubrene Thin Film Characteristics on Mica

47

Acknowledgements. The work was supported by the Austrian Science Foundation (FWF) within the National Research Network (NFN) “Interface controlled and Functionalized Organic Films”.

References 1 2 3 4 5 6 7 8 9 10

Se-W. Park, et al., Appl. Phys. Lett. 90, 153512 (2007). Se-W. Park, et al., Appl. Phys. Lett. 91, 033506 (2007). M. Nothaft et al., Phys. stat. sol. (b) 245, 788–792, (2008). C.H. Hsu, et al., Appl. Phys. Lett. 91, 193505 (2007). V. Podzorov, et al., Phys. Rev. Lett. 93, 086602 , (2004). T. Djuric, et al., E-MRS proceedings 2008, (in print). B.A. Paez, et al., E-MRS proceedings 2008, (in print). Gregor Hlawacek, et al., E-MRS proceedings 2008, (in print). Y. Luo, et al., phys. Stat. Sol. (a) 204, No. 6, 1851–1855 (2007). L. Gránásy, et al., Physical Review E 72, 011605 (2005).

Structural Properties of Rubrene Thin Films Grown on Mica Surfaces T. Djuric1, H.-G. Flesch1, M. Koini1, Sh.M. Abd Al-Baqi2, H. Sitter2, and R. Resel1 1 2

Institute of Solid States Physics, Graz University of Technology, Austria Institute of Semiconductor and Solid States Physics, University Linz, Austria E-mail: [email protected]

Abstract. Structural properties of rubrene thin films on cleaved mica (001) surfaces were investigated by optical microscopy and x-ray diffraction. Optical microscopy shows, that the crystallization of rubrene results in formation of spherulites. X-ray specular diffraction reveals polycrystalline and polymorphic nature of rubrene. The pole figure measurements of films prepared at low deposition rates reveal orthorhombic structure and indicate fiber textures with crystallographic planes (121), (131) and (141) preferentially oriented parallel to the substrate surface. High deposition rate thin films in addition show polymorphism, corroborating the existence of the orthorhombic and the triclinic phase.

1.

Introduction

Rubrene is an aromatic organic molecule with many desirable properties, which make it to a material of choice for the fabrication of single crystal organic field effect transistors. It has the highest reported charge carrier mobility in organic single crystal transistors (20 cm2/Vs) [1]. It is proposed that pristine crystalline rubrene is rather insensitive to oxidation due to its packing [2]. In spite of its very promising electronic behavior as a single crystal, the fabrication of rubrene thin films turned out to be a difficult task. Many efforts to grow well structured rubrene thin films on different substrates resulted in amorphous films with only small, usually spherulitic shaped, crystalline areas [3–5]. This growth of rubrene thin films can be divided into two subsequent stages. In the initial growth stage amorphous islands are formed [4,5]. When a critical coverage is obtained, amorphous islands merge together, coalescence starts and amorphous porous thin films are obtained [5]. The porous film serves as a template for the second growth stage and enables the nucleation of polycrystalline spherulites [5]. Measured charge carrier mobility of crystalline spherulitic areas revealed very small values, ranging between 10–6 and 10–3 cm2/Vs. Rubrene thin films reported here evinced the same morphological properties. Rubrene crystallizes in three different polymorphic phases. The first reported crystal structure of the rubrene is monoclinic (a = 15.500Å, b = 10.100Å, c = 8.800Å, β = 90.55°) [6]. Also a triclinic structure is described (a = 9.150Å, b = 11.600Å, c = 7.160Å, α = 103.53°, β = 112.97°, γ = 90.98°) [7]. Recent investigations report a orthorhombic structure (a = 7.184Å, b = 14.433Å and c = 26.897Å), which is the only one where a full structure solution is available [8–10]. H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_8, © Springer-Verlag Berlin Heidelberg 2009

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2.

Experimental Methods

Rubrene was purchased from Aldrich (elemental purity > 98%) and additionally purified by gradient sublimation. Freshly cleaved mica (001) was used as substrate. Rubrene thin films were deposited by hot wall epitaxy in a vacuum chamber with a base pressure below 10–4 Pa at different deposition rates and substrate temperatures (Ts). Pole figures were measured with a Philips X’pert x-ray diffractometer using CrKα radiation and a secondary side graphite monochromator. Specular scans were performed on a Bruker D8-Discover diffractometer using CuKα radiation. POWDER CELL and STEREOPOLE were used for the evaluation of the specular scans and simulation of pole figures.

3.

Results and Discussion

Figure 1. (a)–(d) shows optical micrographs of the rubrene thin films, which were deposited under different growth conditions. Samples (a)–(c) were grown for 24h at low deposition rate (LDR) (Tsource = 180°C, Twall = 180°C) but at different Ts: 80°C, 90°C and 120°C. Sample (d) was grown for 3h at a high deposition rate (HDR) (Tsource = 235°C, Twall =235°C) while the Ts was held at 120°C. All samples exhibit spherulitic morphologies. Spherulites are aggregates of microcrystals, which arrange in radially growing fiber-like structures. At low Ts and low deposition

Fig. 1. Optical microscopy images of rubrene films deposited on mica substrates. Samples (a)–(c) were grown at low deposition rate at substrate temperatures: (a) 80°C, (b) 90°C and (c) 120°C. Sample (d) is grown at a high deposition rate. Micrographs below present growth stages of rubrene spherulites. Nucleating as single needle bunches (α), splaying through small-angle branching (β) to a spherulite with a spherical envelope (γ).

Structural Properties of Rubrene Thin Films Grown on Mica Surfaces

51

rubrene orth (020)

rubrene rubrene

rubrene

rubrene

rubrene

mica(006)

rubrene orth(002)

(1) LDR T sub= 80° (2) LDR T sub= 90° (3) LDR T sub= 120° (4) HDR T sub= 120°

mica(004)

log(intensity)

rubrene tric (001)

mica(002)

rates (LDR) thin films consist of separate spherulites. With increasing substrate temperature or with higher deposition rate spherulites are growing until impingement. In Fig. 1(c) one can see rather straight boundaries of impinging spherulites This indicates that neighboring spherulites start to grow at the same time with the same growth rate [11]. Otherwise hyperbolic boundaries are obtained, which partially appear in Fig. 1(d) (marked with arrows). Please note that the straight boundaries between the spherulites could appear due to cleavage steps of mica. Depending on their growth behavior spherulites can be divided in two categories [12]. Category I shows radial growth starting from a single nucleation site. Category II spherulites nucleate as single bunches of needles, through small-angle branching the needles spread until a spherical envelope is formed. Figure 2 (α)-(γ) shows the formation of spherulites observed on films grown with a HDR for 15min which indicates spherulites of category II. Figure 2 shows specular scans of rubrene thin films, measured with a specular offset of 3°. The position of the Bragg peaks enables an identification of the crystalline structure together with a preferred orientation of the crystallites. Since the thin film volume is mainly composed of spherulites, it can be supposed that the measured diffraction intensities reflect the crystallographic properties of the spherulites. For the phase identification primarily peaks at low values of the scattering vector qz were used. In comparison with the calculated values measured peak positions at qz = 0.47Å–1 and qz = 0.87Å–1 can be assigned unambiguously to net planes (002) and (020) of the orthorhombic phase, while the measured position at qz = 0.55Å–1 is in a good agreement with the calculated net plane (010) of the triclinic phase. Six other rubrene peaks were detected, which cannot be assigned unambiguously; even the monoclinic phase cannot be excluded definitely. Measured diffraction patterns of the LDR samples don’t show significant diffraction peaks, although crystalline rubrene has to be present.

rubrene

(4) (3) (2) (1)

0.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8

2.0

-1

q z [Å ]

Fig. 2. Specular scans of rubrene films grown on mica. Weak diffraction intensities are observed for the low deposition rate (LDR). The orthorhombic phase with (002), (020) and the triclinc phase with (010) parallel to the substrate can be assigned for the high deposition rate (HDR) film. Curves are shifted for clarity.

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In Figure 3 characteristic pole figures of LDR and HDR films are shown. In case of the LDR film (Fig. 3a,b) a so-called fiber texture is observed. The rather blurred features in Ψ-direction can be explained by multiple crystal orientations which are slightly tilted against each other. Due to defocusing effects the position of pole densities are shifted to smaller values. The crystallographic net planes (121), (131) and (141) of the orthorhombic structure are found to be parallel to the mica surface. For these orientations the aromatic backbone of the molecule is nearly parallel to the substrate surface. These planes were not observed in the specular scan, because their positions overlap with Bragg peaks of mica. In Fig. 3 c,d pole figures of the HDR film, which also reveal a fiber texture, are shown. By the reason of the huge variety of crystal orientations parallel to the mica surface on the one side and the existence of polymorphic phases on the other side, multiple orientations are necessary to explain the pole figures. One additional difficulty is given by the fact, that a full structure solution for the triclinic phase is not known; hence no prediction about high Bragg intensities can be given. Simulated positions of enhanced pole densities with net planes (001) and (010) of the orthorhombic and (010) of the triclinic phase parallel to the substrate surface are in good agreement with the measured pole figures. These orientations are corroborated by the specular scan (Fig. 2).

Fig. 3. Pole figures of rubrene films prepared at low deposition rate are measured at qz = 0.86Å–1 (a) and qz = 1.19Å–1 (b). The calculated pole densities of individual crystal orientations with (141), (131) and (121) are denoted by , and , respectively. High deposition rate films are measured at qz = 1.45Å–1 (c) and qz = 1.72Å–1 (d). The calculated pole densities due to (001), (010) of the orthorhombic phase and (010) of the triclinic phase are denoted by , and , respectively. The single high intensity spots are due to single crystalline mica substrate.

Structural Properties of Rubrene Thin Films Grown on Mica Surfaces

4.

53

Conclusions

Spherulitic growth of rubrene thin films grown on mica (001) surfaces is observed. The growth behavior of rubrene spherulites is assigned to the category II, which nucleates from single bunches of needles. By means of x-ray specular scans and pole figure measurements it is found that spherulites are polycrystalline and polymorphic crystal entities. Pole figures of the LDR film indicate a widely smeared fiber texture with net planes (121), (131) and (141) of the orthorhombic structure oriented parallel to the substrate surface. For these orientations the aromatic backbone of the molecule is nearly parallel to the substrate surface and consequently an identical orientation of the molecules within the first monolayer of spherulites is proposed. Pole figures of the HDR thin film also reveal a fiber texture but with other preferred orientations. Here orthorhombic net planes (001), (010) and the triclinic net plane (010) are parallel to the mica surface. Orthorhombic net planes (001) and (010) correspond to cleavage planes in rubrene. Acknowledgements. The work was supported by the Austrian Science Foundation (FWF) within the National Research Network (NFN) “Interface Controlled and Functionalized Organic Films”.

References 1

V. Podzorov, E. Menard, A. Borissov, V. Kiryukhin, J.A. Rogers and M.E. Gershenson, in Phys. Rev. Lett., Vol. 93, 8, 2004. 2 O. Mitrofanov, C. Kloc, T. Siegrist, D.V. Lang and W.Y. So, A.P. Ramirez, in Appl. Phys. Lett., Vol. 91, 212106, 2007. 3 F. Balzer, M. Schiek, A. Lützeu and K. Al-Shamery in Proc. SPIE, 64706, 2007. 4 Y. Luo, M. Brun, P. Rannou and B. Grevin in Phys. stat. sol., Vol. 204, 1851, 2007. 5 S. Seo, B. Park and Paul G. Evans, in Appl. Phys. Lett., Vol. 88, 232114, 2006. 6 W.H. Taylor in Z. Kristall., Vol. 93, 151, 1931. 7 S.A. Akopyan, R.L. Avoyan and Yu.T. Struchkov in Zh. Strukt. Khim, Vol. 3, 602, 1962. 8 D.E. Henn, W.G. Williams and D.J. Gibbons in J. Appl. Cryst., Vol. 4, 1971. 9 I. Bulgarovskaya, V. Vozzhennikov, S. Aleksandrov and V. Belsky in Latv. PSR Zinat. Akad. Vestis Fiz. Teh. Zinat. Ser., Vol. 4, 53, 1983. 10 I.D. Jurchescu, A. Meetsma and T.T.M. Palstra in Acta Cryst. B, Vol. 62, 330, 2006. 11 H. Müller, in phys. stat. sol. (a), Vol. 66, 199, 1981. 12 L. Granasy, T. Pusztai, G. Tegze, J.A. Warren and J.F.Douglas in Phys. Rev. E, Vol. 72, 011650, 2005.

Rubrene on Mica: From the Early Growth Stage to Late Crystallization Gregor Hlawacek1, Shaima Abd-al Baqi2, Xiao Ming He1, Helmut Sitter2 and Christian Teichert1 1

Institute of Physics, University of Leoben, Leoben, Austria E-mail: [email protected], [email protected] 2 Institute of Semiconductor and Solid State Physics, University of Linz, Linz, Austria Abstract. The fabrication of Rubrene thin films is of interest because of the high mobility observed for Rubrene single crystals. Here, we report on an atomic force microscopy (AFM) investigation of the growth of Rubrene thin films by Hot Wall Epitaxy on mica(001). During the initial formation of amorphous islands, a non-constant growth rate is observed due to temperature dependent changes in the sticking coefficient. Furthermore, the contact angle of these islands – also measured by AFM – depends on temperature. With continuous deposition, island coalescence starts resulting in ramified surface aggregates. The final growth stage is characterized by the formation of crystalline spherulites which also analyzed by AFM.

1.

Introduction

Rubrene is known for its large carrier mobility in single crystal form. Values of up to 20 cm²/Vs have been reported [1,2]. Although this would make it an interesting candidate for various devices, thin films grown from Rubrene have shown hole mobilities lower by 7 orders of magnitude [3]. Unfortunately, very often these thin films are amorphous. Recent studies have shown that in crystalline or polycrystalline films grown with a large overpressure of Rubrene or by utilizing heterostructures, mobilities of up to 0.2 cm²/Vs can be reached [4–6]. Recent work using the weakly interacting substrate SiO2 and an OTS layer to tune the surface energy of the substrate demonstrated mobilities between 0.1 cm²/Vs and 2.5 cm²/Vs [7]. Here, thin films of Rubrene are grown on mica(001) by means of hot wall epitaxy (HWE). The initially formed amorphous islands as well as the morphologies in crystalline spherulites, observed in thick films, are characterized by atomic force microscopy (AFM). The behavior of the growth rate, the contact angle of the Rubrene islands, and their fractal dimension are analyzed in dependence on growth temperature and film thickness.

2.

Experimental Methods

Mica(001). The (001) surface used is a cleavage plane of 2M1 muscovite. It allows easy ex situ cleavage offering large (>100 µm) atomically flat terraces. Furthermore it has shown the capability to align organic molecules by the a strong surface dipole [8].

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Standard quality mica was cleaved before insertion into the HV chamber, thus a fresh cleaved surface with only a few cleavage steps is used. Hot Wall Epitaxy (HWE). HWE is a high vacuum variant of physical vapor deposition with a base pressure of 10–6 mbar [9]. In contrast to many other growth techniques it utilizes the near field of the molecular beam by moving the sample close or even into the hot wall tube that holds the film material. The walls of the tube can be heated separately and are held on a higher temperature than the sample and the source. This prevents deposition on the tube wall and helps to create a uniform flux of molecules. The main advantages of HWE are that the films are grown close to the thermodynamic equilibrium. The main drawback is that the position of the sample close to the evaporator makes an in situ characterization of the film growth impossible. Rubrene. This organic semiconductor consists of a tetracene core with four additional phenyl groups connected by single bonds. In contrast to the desired crystalline phase the amorphous one is not stable against oxidation [10]. The two states can be distinguished easily, since the amorphous phase lacks the typical red color found for crystalline films. Atomic Force Microscopy (AFM). We used an Digital Instruments MultiMode IIIa AFM in intermittant mode to avoid damage to the organic thin film. Conventional Si probes with opening angles of 20° and tip radii of less then 10 nm have been employed. The typical resonance frequency of the used cantilevers is 300 kHz and the force constant is about 40 N/m.

3.

Results and Discussion

Figure 1 shows two series of AFM images obtained from amorphous Rubrene (or more likely oxidized Rubrene [10,11]) thin films grown with a substrate 363 K (a–d) and 393 K (e–h). The deposition times for both image sequences ranged between 2 min and 24 h. First, small circular islands of rather uniform size are formed which then grow (a,b,e,f) and start to coalesce (c,g,h). In Fig. 2 the distribution of island height and island base area for the case of 363 K is presented. The island height changes from 50 nm after 2 min to 120 nm after 60 min. While the width of the height distribution remains constant (20 nm) with ongoing deposition, the width of the island area distribution, however, broadens dramatically as soon as coalescence starts. The same holds for the films grown at higher temperature (lower row of Fig. 1). However, in this case the island density is much lower which can be related to the increased mobility of the molecules on the surface at higher temperatures. Figure 3(a) shows the nominal film thickness ( f ) vs. deposition time. These thickness values are obtained by calculating the total volume of Rubrene divided by the image area of at least three independent images. For the sample grown at higher temperature, f ranges significantly below the film thickness obtained under identical conditions at lower temperature. We can explain this by a change in sticking

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coefficient for Rubrene on mica(001). From the power law fits indicated we obtain exponents of nearly unity (linear dependence as expected) for the higher temperature growth and 1/2 for the low temperature growth. We can understand this behavior if we assume different sticking coefficients, namely a high one for Rubrene on mica(001) and a low one for Rubrene on Rubrene. The observed change in lateral shape of the islands can be evaluated by the fractal dimension D. Figure 3(b) shows the change in D vs. deposition time, where D is calculated by applying a linear fit to the power spectrum [12] of the corresponding 10 µm ×10 µm AFM images. Again, the growth of the amorphous Rubrene islands can be divided in three stages. First, the compact islands are formed and the highest fractal dimension is obtained. Then coalescence starts, and the fractal dimension is reduced as more one-dimensional aggregates are formed. The last stage shown in Fig. 1(d) leads to a further reduction of the dimensionality as the islands become more ramified.

Fig. 1. AFM images of Rubrene thin films grown by HWE on mica(001). Top row samples grown at 363 K, bottom row sample temperature 393 K. Deposition times are (a,e) 2min, (b,f) 15 min, (c,g) 60 min, and (d,h) 24 h.

Fig. 2. Island height (left) and island size histograms for Rubrene films grown at 363 K.

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From the islands three dimensional shape we can make qualitative estimations on the surface free energy. As the Rubrene islands in the initial stage have the shape of a droplet it is assumed that they are amorphous which is confirmed by several techniques [3,13]. Thus, AFM allows measuring the contact angle Ɵ, like for liquids, by analyzing cross sections through the island center [3]. Figure 4 shows two cross sections through islands grown at different temperatures. Both sections are taken from samples with deposition times shorter than the deposition time required for coalescence to allow the largest possible drop volume but before changes in island shape might influence the contact angle. Besides the obvious size difference, the contact angle for the film grown at elevated temperature is larger than the one grown at lower temperature. This is contrary to what would be expected at first glance, since surface energy in general decreases with increasing temperature. The observed behavior can be explained when considering different temperature dependencies of surface energies for Rubrene and mica(001). It seems that the surface energy of Rubrene decreases slower with increasing temperature as the one of mica. As a result the surface becomes more rubrenophobic with increasing temperature.

Fig. 3. (a) Nominal film thickness f from AFM images vs. deposition time. (b) Fractal dimension of the Rubrene islands vs. deposition time. Dashed lines separate the regimes.

Fig. 4. AFM cross section through Rubrene islands. The contact angle is 27° for 393 K and 22° for 363 K.

The later growth stage is characterized by the formation of large spherulites [14]. Figure 5 shows AFM images obtained from different parts of a spherulite typically found after 1 hour of Rubrene deposition with a source temperature of 508 K and a sample temperature of 363 K. Three areas can be distinguished optically: a dark red center region, a lighter iris region, and a transparent matrix. Already from the red color impression we conclude that the spherulites are crystalline. The matrix, however, has turned transparent when the sample was removed from the high vacuum system because the amorphous Rubrene film is not stable against oxidation. The center (Fig. 5(a)) is formed by a rough (rms-roughness: 33 nm), highly crystalline and faceted area that exhibits a slight radial orientation. The surrounding

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iris area (Fig. 5(b,c) rms-roughness 10 nm) shows branched, strongly radial orientated structures, typical for spherulitic growth. The edge of the spherulite is separated from the amorphous matrix by a deep trench (Fig. 5(c)). Furthermore, there is a significant change in height between the spherulite and the matrix region. Both observations are indications of a large mass transport towards the spherulite since it is fed from the amorphous matrix. Figure 5(d) finally represents the amorphous matrix (rms-roughness 0.7 nm) which is covered by many small holes of a few nanometer depth. These holes could be either due to the mentioned massive mass transport but are more likely a result of the oxidation process, since they exit all over the surface.

Fig. 5. Details of a Rubrene spherulite. (a) Center, (b) Iris, (c) spherulite rim, and (d) surrounding amorphous matrix. Z-scale: (a) 300 nm, (b) 50 nm, (c) 30 nm and (d) 5 nm. The insets show the cantilever position relative to the spherulite center.

4.

Conclusions

Quantitative morphological AFM analysis of HWE growth of Rubrene on mica(001) allowed to draw the following conclusions with respect to sticking coefficient and surface energies: Material and temperature depended sticking coefficients lead to non-linear growth rates for amorphous Rubrene on mica. The sensitivity of Rubrene growth with respect to the growth temperature is also reflected in the observed increase of the contact angle for Rubrene on mica(001) with increasing deposition temperature. Acknowledgments. Funding by Austrian Science Fund Projects S9707, S9706.

References 1 2 3 4 5 6 7

(a) A. L. Briseno, et al., Adv. Mater. Vol. 18, 2320, 2006; (b) V. Podzorov, et al., Phys. Rev. Lett. Vol. 93, 086602, 2004. M. E. Gershenson, et al., Rev. Mod. Phys. Vol. 78, 973, 2006. S. Seo, et al., Appl. Phys. Lett. Vol. 88, 232114, 2006. D. Kafer and G. Witte, Phys. Chem. Chem. Phys. Vol. 7, 2850, 2005. J. H. Seo, et al., Appl. Phys. Lett. Vol. 89, 163505, 2006. S.-W. Park, et al., Appl. Phys. Lett. Vol. 91, 033506, 2007. C.-H. Hsu, et al., Appl. Phys. Lett. Vol. 91, 193505, 2007.

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G. Hlawacek et al. (a) H. Plank, et al., Thin Solid Films, Vol. 443, 108–114, 2003, (b) F. Balzer, et al., Appl. Phys. Lett., Vol. 79, 3860–3862, 2001. M. A. Herman, H. Sitter, Molecular Beam Epitaxy, Springer, 1989. M. Kytka, A. Gerlach, J. Kováč, F. Schreiber; Appl. Phys. Lett. 90, 131911 (2007). A. Otomo, et al., Opt. Let., Vol. 27, 891–893, 2002. A. Mannelquist, et al., Appl. Phys. A Vol. 66, 891, 1998. T. Djuric et al., this proceedings volume. (a) Y. Luo, et al., phys. stat. sol. (a) Vol. 204, 1851, 2007; (b) S. Abd-al Baqi et al., this proceedings volume.

β-Sheeted Amyloid Fibril Based Structures for Hybrid Nanoobjects on Solid Surfaces V. Bukauskas1, V. Strazdienė1, A. Šetkus1, S. Bružytė2, V. Časaitė2, and R. Meškys2 1 2

Semiconductor Physics Institute, Gostauto 11, Vilnius, Lithuania Institute of Biochemistry, Mokslininku 12, Vilnius, Lithuania E-mail: [email protected]

Abstract. A self assemblage of three-dimensional β-sheet protein-based supramolecular structures on solid state surfaces are investigated. A set of hybrid proteins containing the abeta40 peptide domain is constructed. Dimeric glucose dehydrogenases and thioredoxin are fused to abeta40 peptide. The supramolecular structures are immobilized on solid surfaces and the properties of the surface nanobjects are studied. Based on analysis of morphology and mechanical properties of these objects it is proved that Aβ40 peptide containing proteins preferably self assemble into island and grain type structures diameter of which is about 20–120 nm. In contrast to this TrxAβ40 fusion proteins preferably form thick (about 7–14 nm) and short (about 1–2 μm) objects. It is experimentally demonstrated that arrangement of fibrils on solid surfaces can be varied by duration of immobilization period and material of solid substrate.

1.

Introduction

During the last decade, artificial smart systems include large number of highly integrated single elements and completed modules dimensions of which extremely decreases due to advantages of nanotechnology. A progress in nanotechnology is partly achieved due to studies of self-assembling of supramolecules and hybrid structures and adopting of so called “bottom-up” approaches. It is extremely increasing interest in integration of biomolecules with nanosized materials aiming to create hybrid materials that combine the evolutionary optimized recognition and catalytic properties of biomaterials with the unique electronic, optical and catalytic functions of nanomaterials. Notable success of these emerging technologies in practical applications stimulates frontier research and development activities aimed at creation of new nanomaterials and nanostructures that are based on self-assemblage of supramolecules [1,2]. It is known that many proteins and peptides aggregate into extended β sheet like structures [3]. Formation of amyloid fibrils is recognized important not only in understanding of human diseases but also is found acceptable for development of ordered nanostructures with modifiable properties and applicable in biotechnology, material science, molecular electronics and related fields. Present study deals with specific aspects of development of hybrid proteins capable to create β-sheeted protein fibrils and integration of the fibrils with solid surfaces. Our study aims to define the most essential conditions and parameters on which the arrangement of the complex object on the solids depends. Morphology,

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mechanical and electrical properties are investigated by scanning probe microscopy in self-assembled complex hybrid systems on solid surfaces.

2.

Samples and Experimental Methods

A set of hybrid proteins containing the Aβ40 peptide domain and dimeric glucose dehydrogenase (GDH) or thioredoxin (Trx) have been constructed in present work. Expression and purification of the proteins was accomplished using prokaryotic system based on Escherichia coli. Details of the construction and expression of hybrid proteins are subject of special report and will be presented elsewhere. Hen egg lysozyme was fibrilized as described in our recent work. Photoluminescence spectra were measured for the hybrid proteins in a solution in the interval of wavelengths between 460 and 560 nm. The wavelength of stimulating light was 446 nm. Binding of thioflavin T to the proteins resulted in distinctive increase of the intensity of the photoluminescence. The supramolecular structures were immobilized on solid surface by deposition of biomolecular objects from colloidal solution. Insulating SiO2 layer on Si, SnO and In2O3 thin films and mica sheets were used for the substrates of combined structures. The substrate was placed on the surface of the liquid and pulled off the liquid after some fixed period of time between 1 and 30 minutes. The samples were dried in air with relative humidity about 30% and surrounding temperature about 295 K. The drying time was about 3 hours. Structure of the sample surfaces and electrical properties were investigated by SPM D3100/Nanoscope IVa (Veeco, Digital Instruments). The surface properties were characterized by the maps of specific parameters obtained in Contact, Tapping and tunneling current (TUNA) modes.

3.

Results and Discussion

Formation of big combined biomolecular structures in a buffer solution was verified by photoluminescence experiments during which amyloid specific dye (thioflavin T) was mixed into the colloidal solution. Thioflavin T related increase in the intensity of fluorescence was detected for the solutions with both GDH– Aβ40 and Trx-Aβ40 hybrid proteins compared to the solutions with single biomolecular components. The intensity increase was obtained after incubation of protein solutions at 310 K in PBS buffer (pH = 7.4) for two days. It was supposed that this period was required for formation of fibrils because the increase in the intensity of fluorescence was comparable with the results obtained for the solution with lysozyme fibrils. Self assemblage of extended biomolecular structures on solid surfaces was verified by the SPM measurements. Supramolecular structures were detected on solid surfaces of practically all samples. Aiming to characterize transferability of the structures from the solution onto solid surface, we performed experiments with sufficiently stable lisozyme fibrils. In these experiments hen egg lysozyme fibrils were immobilized on several types of solid substrates listed in section 2 by deposition of the structures from buffer solution. Typical results of these tests are

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illustrated in Fig. 1 by the SPM images obtained for the structures on the surfaces of In2O3 and mica. The same solution was used for deposition of fibrils in the experiment. The surfaces of mica (see Fig. 2b) and SiO2/Si (not illustrated in this report) substrates were found the most favorable for immobilization of the fibrils. In contrast to this, the surfaces of thin In2O3 films seemed more favorable for self arrangement of ball like objects than fibrils as it is illustrated in Fig. 2a.

(a)

(b)

Fig. 1. SPM images of topography of 10 × 10 μm area for self assembled combined structure based on lysozyme fibrils on surface of (a) In2O3 thin film and (b) mica.

It also follows from these results that the density of immobilized fibrils significantly depends on the type of the solid surface. Clearly higher density of fibrils was obtained on mica than Si. Thin SnO and In2O3 films were covered with comparatively small number of fibrils and far between. It can be noted here that frequently it was difficult to detect the rare fibrils on SnO due to significant roughness of the surfaces of these polycrystalline films. Analysis of the SPM force curves on probed substrates revealed that binding forces are slightly greater on mica surfaces that on the rest substrate materials. On the other hand, the binding force on the SiO2/Si surfaces was lower than on SnO films. In spite of this, the density of fibrils was clearly higher on SiO2/Si surfaces than on SnO. Based on the TUNA tests it was supposed that insulating surface is more favorable for immobilization of fibrils than conductive one even if adhesive forces are characterized by the vise versa proportion on these surfaces. Additional experiments are required for more explicit interpretation of dependences of the arrangement of fibrils on the material of substrates. It was proved experimentally that immobilization period can be effectively used for variation of the density of the supramolecular structures on the surfaces. Influence of the immobilization time on the arrangement of the combined structures is illustrated in Fig. 2 by typical SPM images obtained for hen egg lysozyme fibrils deposited on mica surfaces. After deposition of the lysozyme fibrils under similar conditions, the fibrils were rare on the surfaces if immobilization was shorter

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than 3 minutes while highly dense fibril coating was obtained after period longer than 10 minutes. Moderate and well controlled density of the coverage was obtained between 3 and 10 minutes. Base on the experiments with lysozyme it seems reasonable to suppose that fibrils are transferred from liquid onto solid surface without significant modifications. Dimensions and shape of fibrils seems independent on the origin of the surface and immobilization period. It should be pointed here that some solid surfaces (e.g. In2O3 in our case) and longer immobilization period (>10 minutes) can be assumed being more favorable for agglomeration of fibrils. Assuming formation of fibrils in the solution, self-arrangement of tree-dimensional (3D) strucures on the solid surfaces can be associated with folding, twisting and sticking of fibrils. Immobilized fibrils are stable on the solid surfaces for at least 3–4 weeks. During this period, several analogous SPM images were obtained for each tested sample by separate scans. In general, structures of supramolecular objects of hybrid proteins on the surfaces were individual for each set of fused molecules and different from lysozyme fibrils. Ball-like objects were the most typical objects detected on the surfaces in the experiments with the hybrid proteins. The structure is illustrated in Fig. 3a by typical SPM image. The size of the balls in these structures seemed dependent on the type of hybrid protein but quantitative description of the dependence requires additional studies. Comparatively thick (about 7–14 nm) fibrils were detected in the samples for which the Trx-Aβ40 proteins were used. Length of fibrils was about 1–2 μm. These fibrils were very rare on the solid surfaces. Typical SPM image of such fibril is illustrated in Fig. 3b. Practically continuous granular layer was also visualized in these samples as it is seen in Fig. 3a.

(a)

(b)

Fig. 2. SPM images of 10 × 10 μm area of lysozyme nanofibrils on surface of mica after immobilization period of 3 (a) and 8 (b) minutes. The height (h-scale) is 15 nm.

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(b)

Fig. 3. SPM images of structures based on GDH-Aβ40 (1.5 × 1.5 μm area, h-scale 6 nm) (a) and a fibril based on Trx-Aβ40 (1 × 1 μm area, h-scale 14 nm) (b) on mica surface.

4.

Conclusions

Hybrid proteins based on Aβ40 peptides can self-arrange in three-dimensional objects that are stable on solid surfaces. Supramolecular structures of hybrid proteins are initially obtained in special solutions by original technology. These structures can be transferred from the solution onto solid surfaces by the method used for immobilization of hen egg lysozyme fibrils. SPM images typically display linear objects of diameter about 2 nm and length about 1–10 μm on dry solid surface after immobilization of the lysozyme fibrils. Agglomeration of the fibrils was detected quite often but this effect was much more frequent for the hybrid proteins. Extended object based on hybrid proteins were larger than lysozyme fibrils. The Trx-Aβ40 hybrid protein formed thick (diameter about 7–10 nm) and short (about 1–2 μm) objects. We think that understanding of formation of the objects can be significantly improved after additional studies.

Acknowledgement. The study was supported by the Lithuanian State Science and Studies Foundation (project ProNanoFiHi) and, in SPI, partly by FP6 project Woundmonitor. References 1 2 3

I. Willner, B. Willner and E. Katz, in Bioelectrochemistry, Vol. 70, 2–11, 2007. T. Liedl, T.L. Sobey and F.C. Simmel, in Nanotoday, Vol. 2, 36–41, 2007. G. Colombo, P. Soto and E. Gazit, in Trends in Biotechnology, Vol. 25, 211–218, 2007.

Characteristics of Vacuum Deposited Sucrose Thin Films F. Ungureanu1, D. Predoi1, R.V. Ghita1, R.A.Vatasescu-Balcan2, and M. Costache 2 1

National Institute of Materials Physics, P.O. Box MG-7, Magurele, Bucharest, Romania, Fax:(040)-21-369 01 77, E-mail: [email protected] 2 University of Bucharest, Faculty of Biology, Molecular Biology Center, Bucharest, Romania Abstract. Thin films of sucrose (C12H22O11) were deposited on thin cut glass substrates by thermal evaporation technique (p ~ 10-5 torr). The surface morphology was putted into evidence by FT-IR and SEM analysis. The experimental results confirm a uniform deposition of an adherent sucrose layer. The biological tests (e.g., cell morphology and cell viability evaluated by measuring mitochondrial dehydrogenise activity with MTT assay) confirm the properties of sucrose thin films as bioactive material. The human fetal osteoblast system grown on thin sucrose film was used for the determination of cell proliferation, cell viability and cell morphology studies.

1.

Introduction

The obtaining of uniform and adherent thin films of sucrose with good biological properties (as maintaining cell morphology in an in vitro system) represents an important goal in the field of biological research. The difference between biomaterials and passive materials consists mainly in their specific biochemical function. Sucrose is a compatible osmolyte, belonging to a class of low molecular weight compounds (C12H22O11) present both in prokaryotic and eukaryotic cells in order to protect proteins against aggressive effects of harsh environmental conditions such as cold and heat stress [1,2]. In this paper we present a study regarding the properties of sucrose thin films deposited on cut glass in medium vacuum conditions, used as culture medium for human fetal osteoblast system. The biological tests confirm the interest regarding the culture medium to be studied in tooth decay, due to the fact that the increasing frequency of sugar application alters dental plaque by reducing its mineral protection capacity [3].

2.

Experimental Methods

2.1. Sample Preparation Powder of sucrose (Merck at 99, 99% purity), was deposited by thermal evaporation using a HOCH VAKUUM Dresden system. The thin films were deposited on cut

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glass substrates. For evaporation in vacuum (p ~ 8 × 10–6 torr) it was used a wolfram boat, and an intensity of the maximum current through boat of Imax ~ 40 A for t ~ 5 sec. The thickness of the sucrose thin films was ~190 nm (sample S1). The “as deposited” sucrose thin films were characterized by different techniques namely: FTIR, SEM together with biological tests. 2.2. Sample Characterizations IR spectroscopic studies were performed in the range 1800–400 cm–1 using a FTIR Spectrum BX apparatus (4000–350 cm–1) in transmission mode with the resolution 8 cm–1. The surface morphology for the deposited sucrose thin films was investigated by scanning electron microscopy (SEM) in a XL-30-ESEM TMP system. Cell morphology and viability were both investigated in this study to assess the biocompatibility of sucrose thin films in an in vitro environment. Cell seeding. Human fetal osteoblasts (hFOB 1.19, CRL-11372, American Type Culture Collection) were seeded at an initial density of 1 × 104 cells cm–2 in a medium containing a 1:1 mixture of Dulbecco’s Modified Eagle’s Medium (DMEM) without phenol red and Ham’s F12 medium and supplemented with 0.3 mg/ml G418, antibiotics (100 U/ml penicillin and 100 μg/ml streptomycin) and 10% foetal bovine serum. Cells were grown for 48 h at 37°C in an humidified atmosphere of 5% CO2 on the cell culture plastic supports (control), and sucrose thin films deposited on glass substrates (S1 substrate). Prior to cell culture, the samples were sterilized by UV light exposure. Cell morphology studies. Phase contrast microscopy was used every 24 h, until the culture reached the confluence, to examine the cell morphology of osteoblasts in contact with the test materials. Substrate dependent changes in cell morphology, density and orientation were evaluated by actin labelling with FITC conjugated phalloidin after 48 h of culture. With this purpose the cells were fixed in 4% paraformaldehyde solution, permeabilised using by 2% BSA/0.1% Triton X-100 and incubated with FITC conjugated phalloidin. Then, the samples were rinsed with PBS and analyzed by microscopy in fluorescence. The cell viability was evaluated by measuring mitochondrial dehydrogenize activity with MTT assay. This assay measures the cell activity, proliferation rate and cell viability. The yellow tetrazolium MTT (3-(4, 5-dimethylthiazolyl-2)-2, 5-diphenyltetrazolium bromide) is reduced to insoluble purple formazan granules by all living, metabolically active cells. The precipitated formazan was dissolved in isopropanol, and the absorbance was read at 595 nm. Absorbance values that are lower than the control cells indicate a reduction in the cell activity and viability. Conversely, a higher absorbance rate indicates an increase in cell viability/ proliferation.

3.

Results and Discussion

The deposited thin films of Sucrose were investigated by FTIR spectrometry with the aim to obtain first information about their molecular structure as compared to the powder material (sample S) used for the targets preparation (Fig. 1). The IR spectra for S and S1 samples show the vibration modes of sucrose. In these spectra

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the characteristic bands of sucrose are observed in the 1250–800 cm–1 range both for S and S1 and from this point of view the exposed spectra are similar [4]. As presented in literature [5] the most suitable region for the IR measurements of sucrose has been found to be the 1250–800 cm–1 region. Namely, the shoulder at 800 cm–1 of sample S can be related to CH2 group and the range 950–1300 cm–1 is related to vibration mode of C-O-C group. The vibration range related to hydrogen bonded water molecules adsorbed on the surface is present in the region 1600 cm –1 for the sample S (as can be observed in Fig. 1). The bands at 1750–1850 cm –1 are related to a C = O bond. The IR spectra of the films (sample S and S1) were recorded in ATR (attenuated total reflection) regime. The difference in intensity of different transmission peaks are related to the thickness of sucrose thin films as the measured exposed volume decreases.

Fig. 1. The FT-IR spectra of sucrose powder (sample S) and thin films (sample S1).

We present in Fig. 2 the SEM micrographs for Sucrose powder (sample S) and Sucrose thin film deposited on glass (sample S1). In order to record a SEM image the Sucrose powder was deposited on a double adhesive carbon band and afterwards was deposited a fine layer of gold (the sample S becomes conductive). We remark on sample S a disordered aspect with scratches and valleys on a uniform background specific for non-crystalline samples. For the Sucrose thin film (sample S1 with a thickness of 190 nm) we remark an ordered aspect of droplets in a uniform matrix.

Fig. 2. The SEM images of sucrose powder (sample S) and thin films (sample S1).

One of the purposes of this study was to investigate the effects of sucrose thin films deposited on glass substrates on cell morphology and viability, in an in vitro human fetal osteoblast system. Comparative morphological examination of the cell monolayer adhered to the plastic support (control) and sucrose thin films (S1), by phase contrast microscopy, showed no differences in cell shape and density.

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In both cases, the cells displayed a typical elongated phenotype (Fig. 3A, 3B). The localization of actin was examined and correlations between the cytoskeletal organization and morphology of the cell were evaluated. Phalloidin binds to actin filaments much more tightly than to actin monomers, leading to a decrease in the rate constant for the dissociation of actin subunits from filament ends, which essentially stabilizes actin filaments through the prevention of filament deploymerization [6]. Overall, phalloidin is found to react stoichiometrically with actin, strongly promotes actin polymerization, and stabilize actin polymers [7]. In our study, the cells showed an organized actin network dispersed throughout the cell on both analyzed surfaces (Fig. 4A, 4B). From these experiments we observed that the thin films of Sucrose are an appropriate environment for osteoblast cells proliferation. The difference between the osteoblast cells attached to Sucrose in sample S1 medium relative to control osteoblast culture is not obvious. This fact proves that Sucrose thin films are an appropriate medium to be used together with polysaccharides and iron oxides in biological active medium.

Fig. 3. Appearance in phase contrast microscopy of hFOB 1.19 at 24 h of culture: (A) on plastic dishes (control); (B) on the sucrose thin film deposited on glass support (S1). Original magnification 16x.

Fig. 4. Actin cytoskeletal organization (labeled with phalloidin-FITC) of hFOB 1.19 grown for 48h on: a) plastic support (A) – original magnification 16x; b) sucrose film deposited on glass support S1 (B) – original magnification 32x.

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Conclusions

The sucrose thin films vacuum deposited on glass presented a good adhesion and for in IR analysis it is observed that the band intensity decreases when the thickness of sucrose thin films decreases. The in vitro studies concerning cell morphology and viability displayed by the foetal osteoblast cell line hFOB 1.19 suggests a high biocompatibility of the studied sucrose thin films. They present the feature of preserving the function and structure of osteoblast cells. It is important to conclude that S1 sucrose thin films have micro cells configurations that allow them to be used for obtaining medical biocompatible supports. Acknowledgements. The authors thank to Romanian Scientific Program PNCD II (71-097 and 71-037)/2007 for financial support.

References 1 2 3 4 5 6 7

P. Cioni, E. Bramati, G.B. Strambini, Biophysical Journal, Vol. 88 (2005), 4213–4222 L.C. Provinata, Y. Tou, R.D. Ludescher, Biophysical Journal, Vol. 88 (2005), 3551–3561 E.I.F. Pearce, C.H. Sissons, M. Coleman, X. Wang, S.A. Anderson, L. Wong, Caries Res. Vol. 36 (2002), pp. 87–92 R. Jantas, B. Delczyk, Fibres & Textiles in Eastern Europe, Vol. 13, No. 1(49) (2005), 60–63 F. Cadet, B. Offmann, J. Agric. Food Chem, Vol. 45, No. 1 (1997), pp. 166–171 J.A. Cooper, Effects of Cytochalasin and Phalloidin on Actin. J. Cell Biol. 105 (4): 1473–1478, 1987 J. Wehland, M. Osborn, and K. Weber, Phalloidin-induced actin polymerization in the cytoplasm of cultured cells interferes with cell locomotion and growth. Proc. Natl. Acad. Sci. Vol. 74(12): 5613–5617, 1977

Electropolymerization of Polypyrrole Films in Aqueous Solution with Side-Coupler Agent to Hydrophobic Groups H.M. Alfaro-López1,2, J.R. Aguilar-Hernandez1,*, A. Garcia-Borquez1, M.A. Hernandez-Perez3, and G.S. Contreras-Puente1 1

E.S.F.M. – Instituto Politecnico Nacional Edificio No. 9 U.P.A.L.M. Lindavista C.P. 07738, Mexico D.F., México * E-mail: [email protected] 2 E.S.I.M.E.- I.E – Instituto Politecnico Nacional Edificio No. 2 U.P.A.L.M. Lindavista C.P. 07738, Mexico D.F., México. 3 E.S.I.Q.I.E. – Dpto. Ing. Metalurgica – Instituto Politecnico Nacional U.P.A.L.M. Lindavista C.P. 07738, Mexico D.F., México Abstract. A preliminary study of the electrochemical synthesis and optical characterizacion of the conducting polymer polypyrrole (PPy) was carried out, in order to understand the in-situ electropoliymerization of PPy. Electropolymerization was performed in a three electrode cell by using freshly prepared monomer solutions in presence of a side-coupler agent to hydrophobic groups: sodium dodecylsulfate (DDS), in order to improve the adherence polymer film to the surface of the working electrode. The adherence of the polymer film promotes either the ionic or electronic transport at the interface solution-working electrode. A polar chemical reactive, sodium tetrafluoroborate, TFB, was also used. In order to taylor and optimize the physical properties of the PPy films we varied some parameters during the electropolymerization: the monomer concentration, the electrolyte concentration, pH of the solution and the cell potential. This allowed us to control the oxidation level (impurification or doping) of the polymer. The obtained PPy films were characterized by using UV-Vis and IR spectroscopy. Moreover through scanning electronic microscopy we were able to observe well developed helical structures of polypyrrole.

1.

Introduction

Conducting polymers, like polypyrrole, are semiconducting materials which can be easily processed into thin films for application in organic electronic devices. The physics of these polymers is important, since the disposition of excited state energy governs the efficacy of the polymers as active elements in different devices [1]. Semiconducting polymers have non-degenerated states and in this case the change from double to single bond gives rise to electronic structures with different energy levels. The main difference between semiconducting conjugated polymers and inorganic semiconductors is due to the structural conformation. Conjugated polymers are flexible in nature due to the flexibility of the polymer chains, which also influence positively the exchange of electrical charges [2–3]. Starting from the monomer, pyrrole, the electropolimerization gives rise to the formation of the isomer linked at α positions (α−α). The respective reaction H.-G. Rubahn et al., Interface Controlled Organic Thin Films, DOI: 10.1007/978-3-540-95930-4_12, © Springer-Verlag Berlin Heidelberg 2009

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mechanism is depicted below showing each one of the reaction steps: a) oxidation, b) dimerization, c) deprotonation and oxidation, d) formation of oligomers and deprotonation, e) oxidation f) propagation or g) overoxidation.

In this kind of materials the excess of electrical charge will occupy localized energy stated in the band-gap. If an electrical charge is added to the polymer chain, it will slowly wander to a localized state. This also will originate a local deformation of the chain, giving rise to an increase of the elastic energy of the system. Usually the additional charge to the polymer chain is not a single electron but a charged radical (cation or anion). This local distortion together with the excess of electrical charge produces localized electronic states in the gap region. This state is usually called polaron, a term borrowed from condensed matter physics [4]. Electrical, optical as well as structural properties of conducting polymer, particularly PPy, depends at great extent upon the growth condition [2–4].

2.

Experimental Methods

Electropolymerization of PPy films was carried out in aqueous solutions in a three electrode cell, together with a Ag/AgCl reference electrode provided with a 1 M KCl junction. All the solutions were freshly prepared and purged with nitrogen during 10 minutes, just before electroplymerization. Pyrrole monomer was reagent grade obtained from Aldrich. As working electrode gold sputtered films onto fused quartz and microscope glass slides were used, which was placed in front of the platinum foil counter electrode. An EG&G Princenton Applied Research (PAR) Potentiostat/Galvanostat model 273 semi-automatically controlled was used for the electropolymerization by applying a constant voltage of 0.6 V against the Ag/ AgCl reference electrode. After deposition PPy films were rinsed with deionized water and dryed with nitrogen. Films thickness was measured with a DEKTAK DK2 profilometer. Measured thicknesses of the films were 10 μm in average. Two different chemicals were used in order to introduce the counter-anion into the polymer chains: sodium dodecylsulfate (DDS), pH = 8, which works also as a surfactant agent and improves the adherence of the PPy film to the substrate. The other salt employed was sodium tetrafluoroborate (TFB), pH = 6, both of them at 0.1 M. Optical characterization in the UV-Vis region was performed with a full automatized Lambda 35 Perkin-Elmer spectrophotometer between 275–1100 nm. IR measurements were recorded with a FTIR-System 2000 Perkin-Elmer spectrometer

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in specular reflectance mode, with a MCT (mercury-cadmium telluride) detector, between 400–2000 cm–1, i.e. in the so called “finger print” region.

3. Results and Discussion The electronic characteristics of the PPy films was determined through the absorption spectrum. Figure 1 shows the absorption spectra of a couple of PPy films: one of them grown by using the tensoactive agent (DDS): PPy-DDS and the other one grown with TFB: PPy-TFB. In both cases a well defined broad band, with maximum around 3.8 eV, can be observed. On the low energy side a couple of structures around 1.45–1.50 and 1.95 eV can be observed. All of these absorption bands are shifted to higher energies due to the doping level, as compared to a fully doped PPy film [5].

Fig. 1. Visible absorption spectra for PPy films: DDS solid-line and TFB dotted-line.

A fully doped PPy film usually shows two main absorption bands at 1.40 and 2.6 eV, which are related to the transition of valence band to the bonding level of the bipolaron state [6]. The band at 2.6 eV can be considered as the envelope which includes transition from the valence band to the antibonding level of the bipolaron and polaron states [7,8]. Spectroelectrochemical studies of PPy films have demonstrated that the ratio of the absorbance bands at 1.4 and 2.6 eV (γ = I1.4/I2.6 ), mirrors the doping level of the polymer: the higher the ratio γ the lower the doping of the polymer [9,10]. Thus according to this fact, and because band B appears as a shoulder in both spectra of figure 1, the ratio γ should have a high value, then both PPy-DDS and PPy-TFB films are slightly doped due to the

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overoxidation, which also gives rise to a reduction of the conjugation length. Nevertheless both films are electrochemically active due to the presence of the broad band at 3.86 eV, which is related to the πÆπ* interband transition, i.e. a transition from valence band to conduction band, which means that the band gap of both films is about 3.86 eV. Overoxidation of the studied films was confirmed by infrared measurements. Figure 2 shows the IR spectra of each film, spectra are up shifted in order to be able to compare them. Besides the usual IR bands, overoxidation is confirmed in both samples due to the presence of the carbonyl group ( C = O ) around 1700 cm–1 for the PPy-DDS sample and at 1833 cm–1 [13] for the PPy-TFB sample [11]. The formation of the C = O radical can be explained by considering the fact that pyrrole is a heterocyclic compound, considered as aromatic because of the delocalisation of the π-electrons wich stabilize the ring. These delocalized π-electrons are very reactive and they can promote aromatic electrophilic substitution, wich produces to nitration reactions. Usually aromatic nitriles are obtained by means of nitrogen salts (N2+), wich comes from aromatic amines, pyrrole in our case. The final step of the overall reaction is the production of the radical carbonyl. However, the polymerization of conductive PPy it must be avoided. Plays de role of TFB as electrolyte acts as activator of the reaction. Ar – NH Æ Ar - N2+ Æ omatic amine nitrogen salt

Ar – C ≡ N Æ Ar – COH nitrile carbonyl

(2)

According to the IR measurements and the presence of the carbonyl group, the overoxidation degree must be higher for the PPy-DDS than for the PPy-TFB. All vibrational modes for both films are summarized in Table 1.

Fig. 2. FTIR- spectra of the PPy films: DDS solid-line, and TFB dotted-line.

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Table 1. IR-vibrational modes of the PPy-DDS and PPy-TFB samples

PPy-DDS Group C=C pyrrole ring C4H5N C-N carbonyl C = O aromatic ring PPy-TFB Band Group a double bonds b aromatic ring c C=N d pyrrole ring C4H5N e C-N f carbonyl C = O Band a b c d e

Wavenumber (cm–1 ) 800/924 1042 1218 1286/1700 1570 Wavenumber (cm–1 ) 670/1000 775 1020 1104 1411 1833

Overoxidation will inhibit to a certain extent further growth of the polymer film due to the lost of electrical conductivity. Conducting PPy polymer films usually show a three dimensional homogeneous couliflower-like growth, as recorded either by scanning electron microscopy or atomic force microscopy [12], as shown in Figure 3, right image for the PPy-TFB sample. However, for the PPy-DDS sample the homogeneous growth is stopped due to the high resistivity of the polymer chains in some parts of the film. In the case the polymer chains tend to roll in a kind of helical growth, figure 3, left image. Moreover on the surface of the substrate cilindrical-like tubes, with diameter of the order of 250 nm, can be distinguished. For the sample PPy-TFB, grown in acid medium, there is no total overoxidation of the polymer chain, the IR band corresponding to the carbonyl group is much less intense, due to the interaction of the carbonyl group with the NH- functional group.

Fig. 3. SEM images of the surface of the a) PPy-DDS films (helical-like growth), left, and b) PPy-TFB, (couliflower-like growth) right.

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The formation of the C = N radicals is probably due to two mechanisms: on the one hand, to the protonation of the NH- radical, because it has no paired electrons. On the other hand, due to the acid medium and the presence of another kind of free radicals. This facts is corroborated by the presence of the IR band, C = N, around 1411 cm–1 .

4.

Conclusions

Polypyrrole (PPy) films were electroplymerized by using sodium dodecylsulfate (DDS), pH = 8, and sodium tetrafluoroborate (TFB), pH = 6, in aqueous solution. Both kind of polymer films showed absorption bands due to polaron and bipolaron electronic states. The PPy-DDS films has a higher overoxidation grade than the PPy-TFB one, as observed from presence of the carbonyl band at 1700 cm–1 in the IR spectrum. The nucleophilic attack of the carbonyl group by the NH- compound gives rise to the formation of C = N radical, as shown in the respective IR spectrum. SEM images show an homogeneous growth for the PPy-TFB films, whereas helical structures are seen for the PPy-NaDDS, due mainly to the overoxidation of the polymer chains. Acknowledgements. One of us H.M.A.L thanks for the financial support of CONACyT-Mexico for a Doctoral scholarship. J.R.A.H, A.G.B and G.S.C.P COFAA, SNI, EDI fellows. Work partially supported by IPN-SIP-20070990 and CONACyT-52972

References 1 2 3 4 5 6 7 8 9 10 11 12 13

H. Nalva (Ed.), in Organic Conductive Molecules and Polymers, Vol. 1, John Wiley & Sons, Chichester, England 1997. H.S. Nalwa (Ed.), Handbook of Advanced Electronic and Photonic Materials and Devices, Academic Press, San Diego, 2001. T.A. Skotheim, R.L Elsenbaumer, J.R. Reynolds (Eds.), Handbook of Conjugated Polymers, Marcel Dekker, New York, 1996. W.R. Salaneck, I. Lunstrom, B. Ranby, Conjugated Polymers and Related Materials, Oxford Science Publications, Oxford 1993. O. Chauvet, S. Paschen, L. Forro, L. Zuppiroli, Synth. Met., 63, 115, 1994. Y. Li, R. Qian, Synth. Met., 26, 139, 1988. D. Kim, J. Lee, D. Moo, C. Kim, Synth. Met., 69–71, 471, 1995. E. Genies, J. Pernaut, J. Electroanal. Chem., 191, 1515, 1985. T. Lewis, G. Wallace, C. Kim, D. Kim, Synth. Met., 84–86, 403, 1997. K. Yakushi, L.J. Lauchlan, T.C. Clarke, G.B. Street, J. Chem Phys., 79, 4774, 1983. F. Beck, P. Braun, M. Oberst, Ber. Bunsenges. Phys. Chem., 91, 967, 1987. Q. Pei, R. Qian, Synth. Met., 45, 2123, 1991. http://www.cem.msu.edu/reusch/VirtualText/Spectrpy/InfraRed/infrered.htm

Surface Modification of Polymer Powders by a Far Cold Remote Nitrogen Plasma in Fluidized Bed Lynda. Aiche1, 2, Hugues. Vergnes2, Bernard. Despax1, Brigitte. Caussat2 and Hubert. Caquineau1 1

Laboratoire des Plasmas et Conversion des Energies, 118 Route de Narbonne, 31106 Toulouse E-mail: [email protected] 2 Laboratoire de Génie Chimique, UMR CNRS 5503, ENSIACET/INPT, 5 rue Paulin Talabot, BP 1301, 31106 Toulouse Cedex 1, France E-mail: [email protected] Abstract. In this work, nitrogen was grafted on the surface of polyethylene powders in a fluidized bed coupled to a nitrogen microwave post-discharge, under low pressure (10Torr) and low temperature (