Organic Field-Effect Transistors

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Half Title Page

Organic Field-Effect Transistors

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OPTICAL SCIENCE AND ENGINEERING

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Title Page

Zhenan Bao Jason Locklin

Boca Raton London New York

CRC Press is an imprint of the Taylor & Francis Group, an informa business

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CRC Press Taylor & Francis Group 6000 Broken Sound Parkway NW, Suite 300 Boca Raton, FL 33487-2742 © 2007 by Taylor & Francis Group, LLC CRC Press is an imprint of Taylor & Francis Group, an Informa business No claim to original U.S. Government works Printed in the United States of America on acid-free paper 10 9 8 7 6 5 4 3 2 1 International Standard Book Number-10: 0-8493-8080-4 (Hardcover) International Standard Book Number-13: 978-0-8493-8080-8 (Hardcover) This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. A wide variety of references are listed. Reasonable efforts have been made to publish reliable data and information, but the author and the publisher cannot assume responsibility for the validity of all materials or for the consequences of their use. No part of this book may be reprinted, reproduced, transmitted, or utilized in any form by any electronic, mechanical, or other means, now known or hereafter invented, including photocopying, microfilming, and recording, or in any information storage or retrieval system, without written permission from the publishers. For permission to photocopy or use material electronically from this work, please access www. copyright.com (http://www.copyright.com/) or contact the Copyright Clearance Center, Inc. (CCC) 222 Rosewood Drive, Danvers, MA 01923, 978-750-8400. CCC is a not-for-profit organization that provides licenses and registration for a variety of users. For organizations that have been granted a photocopy license by the CCC, a separate system of payment has been arranged. Trademark Notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation without intent to infringe. Library of Congress Cataloging-in-Publication Data Organic field-effect transistors / edited by Zhenan Bao and Jason Locklin p. cm. -- (Optical science and engineering series) Includes bibliographical references and index. ISBN-13: 978-0-8493-8080-8 (alk. paper) ISBN-10: 0-8493-8080-4 (alk. paper) 1. Organic field-effect transistors. I. Bao, Zhenan. II. Locklin, Jason. III. Series. TK7871.95.O734 2007 537.6’22--dc22 Visit the Taylor & Francis Web site at http://www.taylorandfrancis.com and the CRC Press Web site at http://www.crcpress.com

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Contents Section 1.1 Theoretical Aspects of Charge Transport in Organic Semiconductors: A Molecular Perspective ..........................................1 Demetrio A. da Silva Filho, Yoann Olivier, Veaceslav Coropceanu, Jean-Luc Brédas, and Jérôme Cornil Section 2.1 Charge Carrier Transport in Single-Crystal Organic Field-Effect Transistors ..........................................................................................27 Vitaly Podzorov Section 2.2 Charge Transport in Oligomers..........................................................73 Gilles Horowitz Section 2.3 Charge Transport Physics of Solution-Processed Organic Field-Effect Transistors....................................................................103 Henning Sirringhaus Section 2.4 Contact Effects in Organic Field-Effect Transistors .......................139 Matthew J. Panzer and C. Daniel Frisbie Section 3.1 Design, Synthesis, and Transistor Performance of Organic Semiconductors ................................................................................159 Abhijit Basu Mallik, Jason Locklin, Stefan C. B. Mannsfeld, Colin Reese, Mark E. Roberts, Michelle L. Senatore, Hong Zi, and Zhenan Bao Section 3.2 Dielectric Materials: Selection and Design .....................................229 Ashok Maliakal Section 4.1 Grazing Incidence X-Ray Diffraction (GIXD) ...............................253 Tae Joo Shin and Hoichang Yang

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Section 4.2 Near-Edge X-Ray Absorption Fine Structure (NEXAFS) Spectroscopy ....................................................................................277 Dean M. DeLongchamp, Eric K. Lin, and Daniel A. Fischer Section 4.3 Scanning Probe Techniques .............................................................301 Hoichang Yang Section 5.1 Vacuum Evaporated Thin Films ......................................................341 Alex C. Mayer, Jack M. Blakely, and George G. Malliaras Section 5.2 Solution Deposition of Polymers.....................................................371 Hoichang Yang Section 5.3 Solution Deposition of Oligomers ...................................................403 Howard E. Katz and Chad Landis Section 5.4 Inkjet Printed Organic Thin Film Transistors .................................419 Ana Claudia Arias Section 5.5 Soft Lithography for Fabricating Organic Thin-Film Transistors ........................................................................................433 Kimberly C. Dickey, Kwang Seok Lee, and Yueh-Lin Loo Section 6.1 Radio Frequency Identification Tags ...............................................489 Vivek Subramanian Section 6.2 Organic Transistor Chemical Sensors..............................................507 Luisa Torsi, M. C. Tanese, Brian Crone, Liang Wang, and Ananth Dodabalapur Section 6.3 Flexible, Large-Area e-Skins ...........................................................529 Takao Someya, Takayasu Sakurai, and Tsuyoshi Sekitani Section 6.4 Organic Thin-Film Transistors for Flat-Panel Displays..................551 Michael G. Kane Index ......................................................................................................................595

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Preface Since the early 1990s, remarkable progress has been made in the development of organic thin film transistors (OTFTs). The performance of the best organic materials now rivals that of the amorphous silicon TFTs commonly used as the pixel-switching elements in active matrix flat-panel displays. OTFTs also are of great interest from a technological standpoint, since a major advantage of organic materials is that they can be deposited onto substrates at low temperatures, thus providing for compatibility with plastic substrates. The purpose of this book is to provide a comprehensive survey of present theory, synthetic methodology, materials characterization, and current applications of organic field-effect transistors. The book is divided into six sections dealing with different aspects of organic transistors. Section 1 provides a theoretical description of charge transport in organic semiconductors at the molecular level. Understanding the influence of molecular parameters on charge transport is extremely important in designing new organic semiconductors. The conduction mechanism for organic materials is different from that found in traditional inorganic semiconductors. The molecules are held together by weak intermolecular forces and, except for a few single-crystal devices, the carriers usually move through the material by hopping instead of band transport. Section 2 is divided into four subsections describing the current understanding of charge transport in single-crystal devices, small molecules and oligomers, conjugated polymer devices, and charge injection issues in organic transistors. Section 3 begins with a detailed description of the synthetic methodologies used for organic semiconductors. An in-depth look at molecular design rules and rationales is given for each material system to tailor device characteristics of p- and n-type organic semiconductors. A survey of all reported molecules and correlations between their structure and transistor performance are presented. The second part of the section deals with the design of materials used as dielectric gate insulators in organic thin film transistor devices. Several requirements of dielectrics in OTFTs, including processability, high capacitance, high dielectric strength, high on/off ratio, and low hysteresis, are discussed. Different processing methods for dielectrics are surveyed. The performance of the best OTFTs, such as those fabricated with pentacene, depends not only on the molecular electronic properties of the organic semiconductor but also on their microstructure. Section 4 provides an overview of various characterization techniques used to probe interfacial ordering, microstructure, molecular packing, and orientation crucial to device performance. The section focuses on characterizing the orientation of organic semiconductors at interfaces, primarily through the use of grazing incidence x-ray diffraction (GIXD) and near-edge x-ray absorption fine structure (NEXAFS). These techniques provide valuable electronic and structural information about atoms, molecules, and local chemical functional-

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ities. The section also includes a discussion of scanning probe techniques used to probe local morphology, microstructure, and electrical properties of thin film devices. Unless they are carefully designed, conjugated oligomers are typically insoluble and thin films can only be deposited by vacuum evaporation. Imparting solubility by the introduction of chemical functionality disrupts the molecular π-system’s natural tendency to pack, and most semiconductor molecules that rely heavily on π-orbital overlap to achieve high carrier mobility suffer as a result of functionalization. Organic polymers are typically more soluble, since their chemical structure can be modified in a regular manner with organic side chains and has better filmforming ability than that of small molecules. Section 5 is divided into five parts describing the different processing techniques and associated organic semiconductor thin film growth mechanisms for molecules deposited by vacuum and solution. The last part includes a section on different printing techniques and nonconventional patterning methods by soft lithography. Section 6 provides some current technological examples that utilize OTFTs in their operation. This section includes specific examples of radio-frequency ID tags, chemical and pressure sensors, flexible scanners, display technology, and circuit design. Zhenan Bao Jason Locklin

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Acknowledgments We would like to thank Stefan C. B. Mannsfeld for his advice and critical reading of the manuscript. Also, Jordi Mata Fink, Ajay Virkar, Anna Reichardt, and Fei Qui are acknowledged for their help in formatting and editing.

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The Editors Zhenan Bao received her Ph.D. in chemistry from the University of Chicago in 1995. She is an associate professor in the Department of Chemical Engineering at Stanford University. Jason Locklin received his B.S. from Millsaps College in 1999 and Ph.D. from the University of Houston in 2004. He is currently an assistant professor in the Department of Chemistry and a member of the faculty of engineering at the University of Georgia.

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Contributors Ana Claudia Arias Palo Alto Research Center Palo Alto, California Zhenan Bao Stanford University Stanford, California Jack M. Blakely Cornell University Ithaca, New York Jean-Luc Brédas Georgia Institute of Technology Atlanta, Georgia University of Mons-Hainaut Mons, Belgium Jérôme Cornil Georgia Institute of Technology Atlanta, Georgia University of Mons-Hainaut Mons, Belgium Veaceslav Coropceanu Georgia Institute of Technology Atlanta, Georgia

Dean M. DeLongchamp National Institute of Standards and Technology Gaithersburg, Maryland Kimberly C. Dickey University of Texas at Austin Austin, Texas Ananth Dodabalapur University of Texas at Austin Austin, Texas Daniel A. Fischer National Institute of Standards and Technology Gaithersburg, Maryland C. Daniel Frisbie University of Minnesota-Twin Cities Minneapolis, Minnesota Gilles Horowitz ITODYS University Denis-Diderot Paris, France Michael G. Kane Sarnoff Corporation Princeton, New Jersey

Brian Crone Los Alamos National Laboratory Los Alamos, New Mexico

Howard E. Katz Johns Hopkins University Baltimore, Maryland

Demetrio A. da Silva Filho Georgia Institute of Technology Atlanta, Georgia

Chad Landis Johns Hopkins University Baltimore, Maryland

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Kwang Seok Lee University of Texas at Austin Austin, Texas

Vitaly Podzorov Rutgers University Piscataway, New Jersey

Eric K. Lin National Institute of Standards and Technology Gaithersburg, Maryland

Colin Reese Stanford University Stanford, California

Jason Locklin University of Georgia Athens, Georgia Yueh-Lin Loo University of Texas at Austin Austin, Texas Ashok Maliakal Bell Laboratories Murray Hill, New Jersey George G. Malliaras Cornell University Ithaca, New York Abhijit Basu Mallik Stanford University Stanford, California Stefan C. B. Mannsfeld Stanford University Stanford, California

Mark E. Roberts Stanford University Stanford, California Takayasu Sakurai University of Tokyo Tokyo, Japan Tsuyoshi Sekitani University of Tokyo Tokyo, Japan Michelle L. Senatore Stanford University Stanford, California Tae Joo Shin Pohang Accelerator Laboratory Pohang, Korea Henning Sirringhaus Cavendish Laboratory University of Cambridge Plastic Logic Ltd. Cambridge, United Kingdom

Alex C. Mayer Cornell University Ithaca, New York

Takao Someya University of Tokyo Tokyo, Japan

Yoann Olivier University of Mons-Hainaut Mons, Belgium

Vivek Subramanian University of California, Berkeley Berkeley, California

Matthew J. Panzer University of Minnesota-Twin Cities Minneapolis, Minnesota

M. C. Tanese Università degli Studi di Bari Bari, Italy

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Luisa Torsi Università degli Studi di Bari Bari, Italy

Hoichang Yang Rensselaer Polytechnic Institute Troy, New York

Liang Wang University of Texas at Austin Austin, Texas

Hong Zi Stanford University Stanford, California

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1.1

Theoretical Aspects of Charge Transport in Organic Semiconductors: A Molecular Perspective

Demetrio A. da Silva Filho, Yoann Olivier, Veaceslav Coropceanu, Jean-Luc Brédas, and Jérôme Cornil CONTENTS 1.1.1 Introduction....................................................................................................1 1.1.2 A Primer on Electron-Transfer Theory .........................................................3 1.1.3 Electron-Vibration Coupling and Reorganization Energy ............................5 1.1.3.1 Intramolecular Reorganization Energy ...........................................5 1.1.3.2 Intramolecular Reorganization Energy of Oligoacenes..................8 1.1.4 Electronic Coupling .....................................................................................10 1.1.4.1 Influence of Intermolecular Separation ........................................12 1.1.4.2 Influence of Long- or Short-Axis Displacements ........................13 1.1.5 From Molecular Parameters to Carrier Mobilities ......................................16 1.1.5.1 Influence of the Electric Field ......................................................17 1.1.5.2 Influence of the Reorganization Energy .......................................18 1.1.5.3 Influence of Intermolecular Distance............................................18 1.1.5.4 Influence of Molecular Translations .............................................20 1.1.5.5 Introduction of a Gaussian Disorder.............................................21 1.1.6 Concluding Remarks....................................................................................22 References................................................................................................................22

1.1.1 INTRODUCTION The development of the field of organic electronics has benefited from the unique set of characteristics offered by π-conjugated oligomers and polymers. These

1


2


materials combine the electrical properties of semiconductors with the properties typical of plastics: low cost, versatility of chemical synthesis, ease of processing, and flexibility. In organic field-effect transistors, the key steps of operation involve charge injection and formation of a conducting channel within the organic semiconductor due to application of a gate voltage; upon application of a drain voltage, the charges migrate across the organic layer and are collected at the drain electrode. Charge injection and collection processes and, in most instances, charge transport actually correspond to redox (electron-transfer) reactions. Much success in gaining a better understanding of charge-transport phenomena in organic materials has come recently from extending the theory of electron-transfer reactions, originally formulated by Marcus for the description of redox reactions in solution, to organic semiconductors [1–5]. The charge-transport properties in conjugated materials critically depend on the packing of the chains and degree of order in the solid state [6] as well as on the density of impurities and structural defects [7]. As a result, the measured mobility values can largely vary as a function of sample quality [8]. Overall, the transport mechanism results from a balance between the energy gained by electron delocalization in an electronic band and the energy gained by geometric relaxation and polarization around a charge on an oligomer or polymer segment to form a polaron [9]. In highly purified molecular single crystals, transport at low temperature can be described within a band picture, as shown by Karl and coworkers [10]. As a general rule of thumb, (effective) bandwidths of at least 0.1 eV are needed to stabilize a band regime [9]. In that case, the positive or negative charge carriers are fully delocalized and their mobilities are a function of the width of the valence or conduction band, respectively (i.e., of the extent of electronic coupling between oligomer or polymer chains). When temperature increases, the mobilities progressively decrease as a result of scattering processes due mainly to lattice phonons, as is the case in metallic conductors. Transport can then be described on the basis of effective bandwidths that are smaller than the bandwidths obtained for a rigid lattice. At elevated temperatures, localization steps in and transport operates via a thermally assisted polaron hopping regime where charge carriers jump between adjacent molecules or chains, as described, for instance, by Conwell and coworkers [11]. The hopping regime generally applies in the presence of significant static disorder, dynamic fluctuations [12], and/or impurities; this transport mechanism is thus expected to be operative in most organic field-effect transistors. At the microscopic level, polaron hopping can be viewed as a self-exchange electron-transfer reaction where a charge hops from an ionized site to an adjacent neutral site. In that context, the carrier mobilities are a direct function of the self-exchange electron-transfer reaction rates. In this chapter, we focus on the hopping regime and start with a primer on electron-transfer theory in Section 1.1.2. This section will underline the three major parameters that enter the expression of the electron-transfer rate: reorganization energy, electronic coupling, and driving force. We then discuss some examples of the impact of chemical structure and packing mode on these parameters. Section


Theoretical Aspects of Charge Transport in Organic Semiconductors

3

1.1.3 deals with reorganization energy, and Section 1.1.4 is devoted to electronic coupling. In Section 1.1.5, the role of the driving force (due to the application of an external electric field) is incorporated. This section provides an illustration of how the information on electron-transfer rates gathered at the intramolecular and intermolecular levels can translate into charge carrier mobilities at the macroscopic level.

1.1.2 A PRIMER ON ELECTRON-TRANSFER THEORY Electron-transfer processes, as well as energy-transfer processes, can be viewed as special cases of the nonradiative decay of an electronic state. In the framework of perturbation theory [1,2], the probability for a transition from a discrete initial state ψi (corresponding to the reactants) to a discrete final state ψf (corresponding to the products of the reaction) writes under application of a perturbation V to first order:

Pif =

1 < ψi V ψ f > 2

2

⎡ sin(ω fi t / 2 ) ⎤ ⎥ ⎢ ⎢⎣ ω fi / 2 ⎥⎦

2

(1.1.1)

where t denotes time, ωfi the transition energy between the electronic states i and f, and is the corresponding electronic coupling matrix element. To account for a continuous distribution of final (vibrationally coupled) electronic states, Equation 1.1.1 can be recast by introducing the density of final states ρ(Ef) and summing over all probability densities. Assuming that the function ||2 ρ(Ef) varies slowly with energy, the transition probability per unit time (or transition rate) adopts, in the long-time limit, the simple and widely exploited Fermi’s golden rule form:

kif =

2 2π < ψ i V ψ f > ρ( E f )

(1.1.2)

The transition mechanism involves vibrational motions driving the reaction coordinates from reactants to products. The expression for the rate obtained within the Frank–Condon approximation factorizes into an electronic and a vibrational contribution as:

kif =

2π Vif

2

( FCWD )

(1.1.3)

Here, Vif = is the electronic coupling matrix element and (FCWD) denotes the Franck–Condon weighted density of states. In the high-temperature regime — that is, when assuming that all vibrational modes are classical ( ω i kB T ], the rate kif becomes in the context of the Bixon and Jortner model (for details, see the review in Jortner et al. [2]):

kif =

2π Vif

2

1 4πλ 0 kB T

∞

∑ exp(− S ) Sn! exp[− (ΔΔG n i

i

n =0

0

+ λ 0 + n ω i )2 ] (1.1.6) 4 λ 0 kB T

Here, a single effective quantum mode, ωi, is assumed to contribute to λi. The Huang–Rhys factor, Si = λ i ω h , is a measure of the electron-vibrational coupling interaction. Equations 1.1.5 and 1.1.6 can be further generalized to account for the influence of external electric field and disorder. Thus, in the presence of an electric field (and when neglecting the entropy contributions), the free energy associated with the hop of an electron from site i to site j is given by: ΔGij0 = E j − Ei − e( F ⋅ dij )

(1.1.7)

Here, Ei and Ej are the site energies of molecules i and j; F is the vector associated with the electric field; and dij is the vector connecting the i to site j. Energy disorder in the form of diagonal disorder [13] can be simply introduced by defining a distribution (usually taken as Gaussian) for the site energies. Off-diagonal disorder, which arises from fluctuations in the electronic coupling due to the positional and orientational disorder of the molecules or chains, can also be incorporated. For instance, in the case of through-space interaction, the electronic coupling decays exponentially (vide infra) with intermolecular separation dij: Vij = V(0 )ij ⋅ exp(− γ ij dij )

(1.1.8)



5

Positional disorder can then be modeled by assuming a distribution of dij (or, alternatively, γij) values. We note that Equations 1.1.5 and 1.1.6 were obtained in the framework of perturbation theory and are therefore strictly applicable only in the limiting case of weak electronic coupling (nonadiabatic electron-transfer regime) [14,15]. In a more general case (any Vij < 0.5λ), the thermal activation barrier also depends on electronic coupling while the prefactor is a function of the attempting nuclear frequency νn (frequency for nuclear motion along the reaction coordinate) and the electronic frequency ν e that is equal to the prefactor in Equations 1.1.5 and 1.1.6 (i.e., νe = 2 π Vij

2

4π 2 λkB T ). The limiting case when the electronic coupling is

dominant (i.e., the condition νe >> νn is satisfied) is referred to as the adiabatic electron-transfer regime. In this case, the hopping rate for a self-exchange reaction (ΔG° = 0) is given by [14,15]: kif = ν n exp[−(λ − 2 Vif )2 / 4λkB T ]

(1.1.9)

This equation can also be further modified to accommodate other effects, such as site energy fluctuations, external fields, or disorder. In this chapter, however, we will limit ourselves to the nonadiabatic electron transfer regime. In the next two sections, we review some recent work that addresses, at the molecular level, the nature of the main parameters that govern electron-transfer processes in π-conjugated oligomers and polymers. This molecular approach contrasts with many models developed earlier for organic materials (where these processes were described on a phenomenological basis) and from a macroscopic perspective, thereby masking the actual chemical structures of the systems behind effective parameters.

1.1.3 ELECTRON-VIBRATION COUPLING AND REORGANIZATION ENERGY 1.1.3.1 INTRAMOLECULAR REORGANIZATION ENERGY As emphasized in Section 1.1.2, the reorganization energy is one of the key quantities that control the rates for electron (or energy) transfer. From the rate expression given in Equation 1.1.6, it is clear that the lower the reorganization energy, the higher the rate. The reorganization energy is usually expressed as the sum of inner and outer contributions. The inner (intramolecular) reorganization energy arises from the change in equilibrium geometry of the donor (D) and acceptor (A) sites consecutive to the gain or loss of electronic charge upon electron transfer (ET). The outer reorganization energy is due to the electronic and nuclear polarization/relaxation of the surrounding medium. It is important to bear in mind that, due to the weakness of the van der Waals interactions among organic molecules, the separation of the reorganization energy


6


D

2

A

2

ii

(D2) λi

(A2) λi

Energy

ΔQA 1

i

(D1) λi

i

1

ii

QD

(A1) λi QA

FIGURE 1.1.1 Sketch of the potential energy surfaces (in the monomer coordinate representation) related to electron transfer, showing the vertical transitions, the normal mode displacement ΔQ, and the relaxation energies λi(1) and λi(2).

into inter- and intramolecular contributions remains largely valid even in the case of molecular crystals. We note that in most instances the outer contribution to the reorganization energy is expected to be of the same order of magnitude as the inner part (see Section 1.1.5); it is also expected to be less sensitive to the chemical structure of the constituents than the inner contribution is. The formalisms used to estimate the outer reorganization energy have been mainly developed to describe electron-transfer processes in solution and apply to isotropic media [16]; thus, it is desirable to extend these standard models to account for the anisotropy in the solid state. In this section, we focus on the intramolecular reorganization energy and its description in terms of vibrational modes. In order to illustrate the physical meaning of the intramolecular reorganization energy, we represent in Figure 1.1.1 the potential energy surfaces (PES) of the donor and acceptor involved in an intermolecular ET reaction of the type D + A+ → D+ + A. In the figure, the electronic states D1 or A1 and D2 or A2 correspond to the neutral and cation states of the donor or acceptor, respectively. The ET process can be formally divided into two steps: Step 1 is the simultaneous oxidation of D and reduction of A+ at frozen reactant geometries. (In Figure 1.1.1, this step corresponds to a vertical transition from the minimum of the D1 surface to D2 and a similar A2 to A1 transition.) Step 2 corresponds to the relaxation of the product nuclear geometries. As seen from Figure 1.1.1, the overall intramolecular reorganization energy upon electron transfer consists of two terms [1–5,17–21]: λ i = λ i( A1) + λ i( D 2 )

(1.1.10)

λ (i A1) = E ( A1) ( A + ) − E ( A1) ( A)

(1.1.11)



λ (i D 2 ) = E ( D 2 ) ( D ) − E ( D 2 ) ( D + )

7

(1.1.12)

Here, E(A1)(A+) and E(A1)(A) are the energies of the neutral acceptor A at the cation geometry and optimal ground-state geometry, respectively; E(D2)(D) and E(D2)(D+) are accordingly the energies of the radical cation D+ at the neutral geometry and optimal cation geometry. The vertical transitions involved in Figure 1.1.1 and Equations 1.1.10–1.1.12 are consequences of the Franck–Condon principle that requires the nuclear configurations of the system immediately before and after electron transfer to coincide. However, it is important to note that in addition to the Franck–Condon principle, the principle of energy conservation should also be satisfied for electron transfer to occur [1–5]. In the case of optically driven electron transfer, the mismatch between the electronic vertical transitions (see the lines labeled “i” in Figure 1.1.1) is balanced by the absorption of light. In the case of thermal (dark) electron transfer (which is our main concern here), to satisfy both principles, thermal fluctuations from the equilibrium nuclear configurations of the reactants are needed prior to electron transfer [3–5]. The contribution of each vibrational mode to λi can be obtained by expanding the potential energies of the neutral and cation states in a power series of the normal coordinates (denoted here as Q1 and Q2). In the harmonic approximation, the relaxation energy λi writes [1–5,17–21]: λi =

λj =

kj 2

∑ λ = ∑ ω S j

ΔQ 2j ,

j

j

S j = λ j / ω j

(1.1.13)

(1.1.14)

where the summations run over the vibrational modes ΔQj represents the displacement along normal mode j between the equilibrium positions of the two electronic states of interest kj and ωj are the corresponding force constants and vibrational frequencies Sj denotes the Huang–Rhys factor. The numerical procedure to obtain the reorganization energy consists of the following steps. First, the normal-mode coordinates and force constants of the electron donor and acceptor are determined. The standard rectilinear normal modes Q1(2) are obtained as a linear combination of Cartesian displacements [22]: Q1(2 ) j =

∑L

1( 2 ) kj

k

(q1(2 ) k − q1((02)) k )

(1.1.15)


8


The matrix L1(2) connects the 3n-6 (n is the number of atoms in the [nonlinear] molecule) normal coordinates with the set of 3n mass-weighted Cartesian coordinates q1(2); the vectors q1(0) and q2(0) correspond to the stationary points on the adiabatic potential surfaces of states 1 and 2, respectively. Then, the normal mode displacements ΔQ1(2) are obtained by projecting the displacements Δq = q1(0) – q2(0) onto the normal-mode vectors. Finally, substituting the calculated quantities into Equations 1.1.10, 1.1.13, and 1.1.14 provides the total relaxation energy.

1.1.3.2 INTRAMOLECULAR REORGANIZATION ENERGY OF OLIGOACENES We now describe how the numerical approaches discussed previously were applied to the case of oligoacenes containing from two to five rings: naphthalene, anthracene, tetracene, and pentacene [15,23–27]. These oligomers are of high current interest because of their high charge mobilities in the crystalline state, in particular tetracene, pentacene, and their derivatives such as rubrene [28,29]. We note that all calculations including geometry optimizations and normal-mode analyses were performed at the DFT (density functional theory) level with the hybrid B3LYP functional using the standard 6-31G** basis set. The bond-length modifications upon positive ionization show a consistent trend along the series. Naphthalene displays the largest geometry relaxations, with changes in C–C bond lengths on the order of 0.03 Å. This value is reduced to ca. 0.02, 0.015, and 0.01 Å in anthracene, tetracene, and pentacene, respectively. The geometry distortions, as well as the changes in atomic charge densities (Mulliken populations), are found to spread over the entire molecule. The theoretical estimates of the relaxation energies and total reorganization energies obtained from the normal mode analysis are in excellent agreement with the values computed directly from the adiabatic potential energy surfaces [30]. The calculated reorganization energies evolve from 187 meV in naphthalene to 137 meV in anthracene, 113 meV in tetracene, and 97 meV in pentacene [30]. These values are also in good agreement with the results of previous calculations by Kato and Yamabe [31–35] and Klimkans and Larsson [36]. Our results indicate that the main contribution to the relaxation energy comes from high-energy vibrations. This high-energy contribution is in fact divided over several vibrational modes with wave numbers in the range of 1200–1600 cm–1. The contribution to λi from low-energy vibrations is negligible in anthracene and tetracene and is very small in naphthalene and pentacene. We have also carried out the Frank–Condon simulation of the shape of the first ionization peak of anthracene, tetracene, and pentacene using the DFT/B3LYP estimates of the frequencies and Huang–Rhys factors obtained from normal-mode calculations. The vibrational frequencies and Huang–Rhys factors calculated for pentacene along with the Frank–Condon simulation of the shape of the first ionization peak are shown in Figure 1.1.2. In general, the positions and shapes of the peaks are very well reproduced. Similar results were obtained for the oligoacenes and their derivatives [24,26,27,37,38]. These results underline the importance of multimode effects to obtain a detailed understanding of the UPS band shapes in oligoacenes.



9

Intensity (arb. units)

1 0.8 0.6 0.4 0.2 0

6.4

6.6

6.8

7

Energy (eV) 0.125 0.1

S

0.075 0.05 0.025 0

0

250

500

750 Frequency

1000

1250

1500

1750

(cm−1)

FIGURE 1.1.2 (Top) DFT/B3LYP simulation (dashed lines) of the vibrational structure of the UPS first ionization peak of pentacene (solid lines) and (bottom) computed Huang-Rhys factors as a function of the frequency. The normal modes of the cation species with the largest Huang-Rhys factors ((10 modes) have been used for the simulations. A scaling factor of 0.9613 has been applied to the computed frequencies. The transition intensities were convoluted with Lorentzian functions with full-width at half-maximum (FWHM) of 0.060 eV. The HOMO wavefunctions obtained at the DFT/B3LYP level is also illustrated.

Overall, the intramolecular reorganization energies in tetracene (0.11 eV) and pentacene (0.10 eV) rank among the smallest λi values that have been calculated or measured for molecules. A smaller value of λi (0.045 eV) has been found only in the case of phthalocyanine [39]. The tetracene and pentacene values are about three times as small as in TPD (0.29 eV), which is a hole-transport material widely used in organic molecular devices. Interestingly, side-chain derivatizations of pentacene in the form of ethynylsilyl substitutions have been reported by Anthony and coworkers [40]. We have found that such substitutions actually lead to a significant increase (by about 50%) in the intramolecular reorganization energy, due to the involvement of the side chains in the geometry relaxation process upon ionization. In contrast, Wudl and coworkers have synthesized a tetra-methyl derivative of pentacene with the goal of improving the processability of the material [41]. These authors calculated the reorganization energy in the same way as described earlier and found that it remains exactly the same as in pentacene because, in this instance, the substituents have a saturated nature and do not couple to the geometry relaxations of the conjugated backbone [42].


10


The origin of the small reorganization energy values in tetracene and pentacene can be traced to a combination of macrocyclic rigidity and full delocalization of the frontier molecular orbitals [24,26,27]. Accordingly, other molecules that have been found to present small intramolecular λi values are fullerenes, as described by Devos and Lannoo [43], phthalocyanines [39,44], or discotic macrocycles [45]. We note that the reorganization energy λi is directly related to such quantities as the polaron binding energy (Epol = λi/2) and the dimensionless electron–phonon parameter λe–ph (λe–ph = λi N(EF), where N(EF) is the density of states at the Fermi level [15,23–25,43]. The electron–phonon parameter is a key value in the conventional theory of superconductivity. Therefore, the results discussed before are especially relevant in the development of adequate polaron models to understand superconductivity and charge transport in organic molecular systems.

1.1.4 ELECTRONIC COUPLING A number of computational techniques, based on ab initio or semiempirical methodologies, have been developed to estimate the electronic coupling Vif; they have recently been reviewed in references 1, 2, and 46. A robust approach to compute Vif is to describe the diabatic states of the reactants and products by means of a Slater determinant and to compute their splitting at the transition state [47,48]. Li and coworkers have applied this approach to benzene and biphenyl dimers using concerted linear reaction coordinates to define the geometry of the transition state [47,48]. Another approach is to use Koopman’s theorem and to estimate (in the context of a one-electron picture) the transfer integrals t for holes or electrons as half the splitting of the highest occupied molecular orbital (HOMO) or lowest unoccupied molecular orbital (LUMO) levels, respectively, in a system made of two chains in the neutral state. In the case of benzene and biphenyl dimers, a good quantitative agreement is observed between the two approaches [49]. The applicability of Koopman’s theorem was confirmed in a study by Pati and Karna [50]; we note that the direct coupling Vif was not evaluated at the transition-state geometry, on the basis that the coupling amplitude hardly depends on the actual nuclear configuration, a feature often referred to as the Condon approximation. The transfer integrals can be estimated in a yet simpler approach from the spatial overlap between the two molecular orbitals in interaction [51,52]. All these considerations explain why many theoretical studies have made use of Koopman’s theorem to estimate electronic couplings [51–57]. We note that much care has to be taken when Koopman’s theorem is used to estimate the transfer integrals in asymmetric dimers, as has been extensively discussed elsewhere [58,59]. In such instances, part of the electronic splitting can simply arise from the different local environments experienced by the two interacting molecules, which create an offset between their HOMO and LUMO levels prior to their interaction due to polarization and/or electrostatic effects. In order to evaluate the effective couplings, this offset can be accounted for by performing calculations using molecular orbitals localized on the individual units as basis set [59] or by applying an electric field to promote the resonance between the electronic levels, as done by Jortner and coworkers [56]. This artifact can also be prevented when



11

E LUMO

Conduction band

HOMO

Valence band

FIGURE 1.1.3 Illustration of the bonding-antibonding interactions between the HOMO/ LUMO levels of two ethylene molecules in a cofacial configuration; we also illustrate the formation of the valence and conduction bands when a large number of stacked molecules interact.

estimating the electronic coupling from the overlap between the molecular orbitals of the isolated molecules. The electronic splittings reported next have been calculated within Koopman’s theorem using the semiempirical Hartree–Fock INDO (intermediate neglect of differential overlap) method; interestingly, the INDO method often provides transfer integrals of the same order of magnitude as those obtained with DFT-based approaches [45,60]. It is of interest to note that when building infinite one-dimensional stacks of chains, the widths of the corresponding valence and conduction bands are usually found to be nearly equal to four times their respective t integrals; this indicates that in most instances the tight-binding approximation is relevant. In the remainder of this section, our goal is simply to lay out a basic understanding of how molecular packing and molecular size affect the electronic coupling between neighboring oligomers [61]. As a first step, it is useful to consider the simple example of a dimer made of two perfectly superimposed ethylene molecules. In an isolated ethylene molecule, the HOMO corresponds to the bonding situation between the two π-atomic orbitals (zero-node case), while the LUMO corresponds to the (fully) antibonding situation. In a cofacial dimer, the interaction between the two molecules leads to a splitting of the HOMO level and a splitting of the LUMO level, as illustrated in Figure 1.1.3. The HOMO splitting is very large (0.54 eV for an intermolecular distance of 4 Å); the reason is that the interaction between the HOMO wave functions of the two molecules is either fully bonding (which leads to the much stabilized HOMO-1 level of the dimer) or fully antibonding (leading to the much destabilized dimer HOMO level). For the LUMO level, the splitting is found to be much smaller (0.15 eV) because direct bonding interactions are compensated by diagonal antibonding interactions in the dimer LUMO wave


12


function or direct antibonding interactions are compensated by diagonal bonding interactions in the LUMO + 1 level (see Figure 1.1.3). Thus, as a qualitative rule in the case of perfectly cofacial configurations, one can state that the lower the number of nodes in the wave function of a given frontier level of an isolated chain is, the larger the splitting of that level in a dimer (or larger clusters) will be. Since in isolated π-conjugated chains the LUMO wave function has usually one more node than the HOMO wave function, the LUMO splitting is expected to be smaller than the HOMO splitting. Qualitatively, for large clusters, this will translate into larger HOMO bandwidths. This feature is what gave rise earlier to the conventional wisdom that, in crystals or crystalline films of π-conjugated chains, the hole mobility is expected to be higher than the electron mobility. We will see later that this concept breaks down in numerous instances.

1.1.4.1 INFLUENCE

OF INTERMOLECULAR

SEPARATION

We now consider perfectly cofacial dimers made of two tetracene molecules and examine the evolution of the electronic splittings of the HOMO (Figure 1.1.4a) and LUMO (Figure 1.1.4b) levels as a function of the distance, d, between the molecular planes (see inset in Figure 1.1.5). Although fully cofacial configurations are rarely encountered in actual crystal structures, it is of interest to start by studying such geometries since they provide a highly symmetric reference point and lead to the largest electronic splittings (here, for instance, 280 and 200 meV for the HOMO and LUMO levels, respectively, for an interchain distance of 4 Å). The results are illustrated in Figure 1.1.5. Consistent with our previous discussion, the HOMO splitting is calculated to be larger than the LUMO splitting whatever the interchain separation may be. The amplitudes of the electronic splittings are observed to decay exponentially when the interchain distance is increased; this simply translates the exponential decay in intermolecular overlap of the π-atomic orbitals when the two oligomers are pulled apart. By fitting the HOMO/LUMO

(a)

(b)

FIGURE 1.1.4 B3LYP/6-31G(d,p) HOMO (a) and LUMO (b) wavefunctions for tetracene.



13

1400

1200 d

Splitting (meV)

1000

800

600

400 HOMO 200

LUMO

0 3.3

3.6

3.9

4.2

4.5

4.8

5.1

5.4

5.7

d spacing (Å)

FIGURE 1.1.5 Evolution of the INDO-calculated electronic splittings of the HOMO and LUMO levels in a cofacial dimer made of two tetracene molecules as a function of the intermolecular separation.

splitting (Figure 1.1.5) to Equation 1.1.8, we obtained for holes or electrons V0 = 0.717 × 106 meV or 0.622 × 106 meV, and γ = 2.137 Å–1 or 2.187 Å–1, respectively. An important result is that the electronic splittings vary by as much as a factor of three to four between 3.4 and 4.0 Å. This range corresponds to the typical intermolecular distances found in organic conjugated crystals and thin films.

1.1.4.2 INFLUENCE

OF

LONG-

OR

SHORT-AXIS DISPLACEMENTS

In many instances, cofacial packing involves the displacements of adjacent molecules along their long and/or short molecular axes. Figure 1.1.6 describes the evolution of the HOMO and LUMO splittings in dimers where the top tetracene molecule is translated along its long axis (while keeping the interchain distance fixed at 3.4 Å). As expected, the overall effect of this displacement is to reduce the overlap and thus the electronic coupling. The most interesting result is the appearance of oscillations in the values of the splitting. The important consequence of this difference in oscillation period is that small translations can lead to situations where the electronic splitting is larger for the LUMO than for the HOMO and hence where electrons can possibly be more mobile than holes. For instance, for a shift of about 4.5 Å, the reversal in the relative amplitude of the splittings is very significant: We calculate a LUMO splitting of 381 meV and a HOMO splitting of 17 meV.


14


Short axis displacement

Long axis displacement

HOMO LUMO

1000 800

1000 800

600

600

400

Splitting (meV)

Splitting (meV)

HOMO LUMO

200 0 −200

400 200 0

−400

−200

−600

−400

−800

−600 0.0

2.5

5.0

7.5

Long axis displacement (Å)

10.0

12.5

0 1 2 3 4 5 6 Short axis displacement (Å)

FIGURE 1.1.6 Evolution of the INDO-calculated electronic splittings of the HOMO and LUMO levels in a dimer formed by two tetracene molecules separated by 3.4 Å as a function of the degree of translation of one molecule along its long axis (left) and short axis (right).

The calculated evolutions can be rationalized in terms of the phase and the nodal properties of the HOMO and LUMO orbitals of a single tetracene molecule [61,62]. In the HOMO level, the distribution of the positive and negative LCAO (linear combination of atomic orbitals) coefficients shows a change in the sign of the wave function of every monomer unit (see Figure 1.1.4a). This pattern leads to extrema in the calculated electronic splittings for degrees of translation corresponding roughly to multiples of the monomer unit size. Large electronic splittings are dominated by full bonding or antibonding interactions between the π-atomic orbitals localized over the carbon–carbon bonds within the benzene rings. In contrast, small splittings are calculated for geometries where the global overlap (and hence the HOMO splitting) is considerably reduced by the compensation of bonding and antibonding interactions between the two bonds of one chain and the two adjacent bonds of the other chain. The evolution of the LUMO splitting also displays maxima and minima and can be explained once again by considering the shape of the LUMO orbital. Large or negligible small values are observed when the benzene rings of one molecule overlap the rings or half of the ring of the second molecule, respectively. The oscillation



15

period of the curve is roughly the same as that calculated for the HOMO splitting, although an increasing dephasing is observed; when the displacement is about 6.0 Å, the LUMO splitting has opposite phase with respect to the HOMO splitting. Note that an overall decrease in the HOMO and LUMO splittings occurs for increasing translational shifts. This simply results from the progressive reduction in the overall extent of spatial overlap between the two oligomers. In the case of the long-axis displacement, when the spatial overlap of the two molecules is zero (e.g., displacements bigger than the size of the molecule), the splitting is also close to zero. We now turn to the impact of translating the top oligomer along its short axis. The HOMO and LUMO splittings decay to zero when the molecule is shifted by more than 6.0 Å. This is more then twice the lateral width of the tetracene molecule (see Figure 1.1.6), illustrating clearly that the estimate of the coupling based purely on the “spatial overlap” of the two molecules can be misleading. In the case of the long-axis displacement, zero spatial overlap does mean negligible HOMO/LUMO splitting, while in the case of short-axis displacement, the orbital splitting is still reasonably large in the situation where the spatial overlap tends to zero. Again, the oscillations in splittings can be rationalized by looking at the wave functions. The antibonding character over the bonds along the short molecular axis found in the HOMO level leads to the appearance of a minimum in the course of translation. In contrast, the LUMO wave function does not change sign along the short axis; thus, the translation preserves dominant bonding or antibonding interactions in the LUMO + 1 or LUMO level of the dimer, respectively. These interactions attenuate as the translational shift increases, leading to a decrease in splitting without change in sign. The case of rubrene — a tetraphenyl derivative of tetracene lately gaining much attention due to reported field-induced hole mobility of up to 20 cm2/Vs at room temperature [63] — illustrates how the results presented earlier can prove useful to rationalize the charge transport properties of this system. The largest electronic coupling is found for dimers along one of the crystallographic directions (see Figure 1.1.7). A combination of zero short-axis displacement and a reasonably large longaxis displacement (ca. 6.13 Å) still results in a large electronic coupling between the molecular units [37]. This result can be rationalized in terms of the oscillatory behavior of the electronic coupling as a function of the long-axis displacement (see Figure 1.1.6). Although the overall coupling decreases with the increase in displace-

a 3.74Å 6.13Å

FIGURE 1.1.7 Rubrene dimer along the a direction. The π-stacking distance and long-axis displacement are indicated.


16


ment — around 6.13 Å — an extremum exists and thus the splitting observed at this configuration is still a considerable fraction of the maximum splitting computed for the cofacial configuration. To conclude this section, we underline that combining the results of calculations on transfer integrals (electronic couplings) and on intramolecular reorganization energies (electron- and hole-vibration constants) allows one to gain a basic understanding of the impact of molecular parameters on the intrinsic electron and hole mobilities in π-conjugated materials. In the following section, we adopt a pragmatic approach and, assuming reasonable values for the outer reorganization energy, we show how it is possible to make the connection between the molecular parameters described earlier and the carrier mobilities in model one-dimensional stacks of pentacene molecules.

1.1.5 FROM MOLECULAR PARAMETERS TO CARRIER MOBILITIES In this section, our goal is to make the connection between molecular scale and mesoscopic scale. We evaluate charge transport over large distances, in the presence of an electric field, by inserting into Monte Carlo simulations the electron-transfer rates computed at the quantum-chemical level within the Bixon and Jortner model (Equation 1.1.6) for pairs of interacting molecules. We illustrate this approach on model systems made of one-dimensional stacks of pentacene molecules. The restriction to one-dimensional stacks allows us to rationalize easily the variations calculated in charge mobilities when: (1) the relative positions of the molecules are changed; (2) structural disorder is introduced; and (3) the external electric field is modulated. This methodology can, however, be readily extended to two- and three-dimensional structures. A detailed description of the methodology is given elsewhere [64]. Briefly, the electron- or hole-transfer rates calculated at the INDO quantum-chemical level are injected into Monte Carlo (MC) simulations to evaluate the propagation of a single charge along the pentacene one-dimensional stacks. MC methods rely on the use of random variables; in our case, the random variable is the occurrence of a hop between two adjacent molecules separated by a given distance. This distance is counted as positive if the hole propagates in the direction of the electric field and as negative in the opposite direction; this counting scheme is reversed for electrons. In the MC algorithm, the first step is to choose randomly the direction along which charge hopping takes place in a given iterative cycle (i.e., a jump to the left or right nearest neighbor in the one-dimensional stack). The second step is to calculate the probability of transferring the charge in the chosen direction (to the right in Equation (1.1.16)) as: p→ =

k→ k→ + k←

(1.1.16)

where k→ and k← correspond to the transfer rates in the right and left directions, respectively; these are different in the presence of the electric field.



17

Finally, a random number is generated between 0 and 1 and is compared to p→. If the random number is lower than the transfer probability, the transfer in the right direction is accepted. Otherwise, the transfer is rejected, the charge carrier remains in its initial position, and a new iterative loop is performed. After a large number of iterative cycles (typically between 1010 and 1013), the mobility can be estimated directly as: μ = d/(τ F)

(1.1.17)

with F the magnitude of the electric field, d the total distance traveled by the charge (summed as positive and negative contributions depending on the direction of the hop with respect to the field), and τ the total time calculated as the sum of the inverse of the transfer rates. The electric field is varied here from 104 to 107 V/cm; these are reasonable values for devices a few hundred nanometers thick under the application of 10–100 V between the electrodes (note that we assume a linear voltage drop across the device). Thus, this approach offers a way of connecting molecular parameters such as λ, ΔG°, and t to macroscopic quantities such as charge mobility. There is no limitation in size since the system is built progressively around the charge when moving along the stacks. We recall that the inner reorganization energy λi of pentacene has been calculated to be 0.10 eV for holes and 0.13 eV for electrons [25]. Since the electric field is consistently applied in the direction perpendicular to the molecular plane of the pentacene oligomers, the magnitude of λi is assumed to be field independent. This is further validated by recent calculations showing that λi is not affected in the presence of an external electric field applied along the long molecular axis of pentacene molecules [65]. The outer reorganization energy λs has been estimated on the basis of a modified expression, taking into account the actual shape of the donor and acceptor molecules; values on the order of 0.3–0.4 eV are obtained when considering the dielectric characteristics of organic matrices [66].

1.1.5.1 INFLUENCE

OF THE

ELECTRIC FIELD

We report in Figure 1.1.8 the field dependence of the hole and electron mobilities calculated for a one-dimensional array of pentacene molecules in a cofacial configuration, with an intermolecular distance fixed at 4 Å. The simulations have been performed at 300 K for an external reorganization energy of 0.4 eV. The transfer integrals associated with such an arrangement amount to 0.13 and 0.10 eV for holes and electrons, respectively. The mobility increases by about a factor of four for both holes and electrons when the electric field is varied from 104 to 107 V/cm. Since the magnitude of the electric field appears in the denominator of the mobility expression, the results further indicate that the charge velocity (d/τ) increases faster than the magnitude of the field. The ratio between the hole and electron mobilities at any field perfectly matches the ratio of the corresponding transfer rates. It is interesting to note that the calculated mobility values (in the range 10–1–1 cm2/Vs) are very reasonable; the mobilities reported for pentacene thin films, and crystals are usually in the range of 0.1–5 cm2/Vs [67].


18


0.8 0.7 4Å

Mobility (cm2/ Vs)

0.6 0.5 0.4 0.3 0.2 0.1

Hole Electron

0 0

1

2

3

4 Field

5

6

(106

V/cm)

7

8

9

10

FIGURE 1.1.8 Evolution of the hole mobility as a function of the magnitude of the electric field in a one-dimensional array of pentacene molecules separated by 4 Å.

1.1.5.2 INFLUENCE

OF THE

REORGANIZATION ENERGY

Since the external reorganization energy may vary significantly in specific cases (for instance, when the molecules are located close to structural defects or in the vicinity of metallic electrodes), we have investigated the way the hole mobility evolves as a function of the magnitude of λs. Monte Carlo simulations have been performed on the same one-dimensional array of pentacene molecules as discussed earlier, with an electric field fixed at 5 × 106 V/cm and at 300 K. Figure 1.1.9 illustrates the mobility drop with λs. The mobility is reduced by about a factor of five for an increase of λs from 0.2 to 0.4 eV or from 0.4 to 0.6 eV; it is strongly hampered beyond 0.7 eV. In this case, the evolution simply reflects that for the transfer rate (see inset of Figure 1.1.9).

1.1.5.3 INFLUENCE

OF INTERMOLECULAR

DISTANCE

Changes in the distance between molecules can be modulated by substituents attached to conjugated backbones. Accordingly, we have examined the evolution of hole mobility in the one-dimensional array of pentacene molecules when varying the intermolecular separation between 3.4 and 5 Å; see Figure 1.1.10. In the remainder, unless stated otherwise, the simulations have been achieved with λs = 0.4 eV and for two different magnitudes of the electric field (106 and 5 × 106 V/cm). In all cases, the mobilities drop exponentially with an increase in the intermolecular distance and are shifted rigidly depending on the magnitude of the electric field; this evolution has to be traced back to that of the transfer integrals since the overlap between the HOMOs decreases exponentially with distance.



1.8

Transfer rate (s−1)

1.6

Mobility (cm2/ Vs)

1.4 1.2 1 0.8

19

4.50E + 14 4.00E + 14 3.50E + 14 3.00E + 14 2.50E + 14 2.00E + 14 1.50E + 14 1.00E + 14 5.00E + 14 0.00E + 00 0.2

0.3

0.4 0.5 λs (eV)

0.6

0.6

0.7

0.4 0.2 0 0.2

0.25

0.3

0.35

0.4

0.45

0.5

0.55

0.6

0.65

0.7

λs (eV)

FIGURE 1.1.9 Evolution of the hole mobility as a function of the magnitude of the external reorganization energy in a one-dimensional array of pentacene molecules separated by 4 Å.

10 F = 106 V/cm F = 5 × 106 V/cm 1 Mobility (cm2/ Vs)

3.4

3.6

3.8

4

4.2

4.4

4.6

4.8

5

0.1

d 0.01

0.001 Intermolecular distance (Å)

FIGURE 1.1.10 Evolution of the hole mobility as a function of intermolecular distance in a one-dimensional array of pentacene molecules.


20


Transfer integral (eV)

0.45 0.4

Mobility (cm2/ Vs)

0.35 0.3 0.25

0.14 0.12 0.1 0.08 0.06 0.04 0.02 0

d

4Å

0

0.2

0.5

1

1.5 2 Shift (Å)

2.5

3

3.5

F = 106 V/cm F = 5 × 106 V/cm

0.15 0.1 0.05 0 0

0.5

1

1.5

2

2.5

3

3.5

Shift (Å)

FIGURE 1.1.11 Evolution of the hole mobility as a function of the shift applied to every other molecule in a direction perpendicular to the stacking axis in a one-dimensional array of pentacene molecules separated by 4 Å. The inset shows the evolution of the corresponding transfer integral.

1.1.5.4 INFLUENCE

OF

MOLECULAR TRANSLATIONS

We have looked at the impact of translating molecules in a direction parallel or perpendicular to the stacking axis. Figure 1.1.11 shows the variation of the hole mobility when translating every other molecule by a distance d perpendicularly to the stacking axis (note that a single transfer integral value characterizes the full system). The conventional wisdom here would be that the mobility goes down with an increase in displacement due to the progressive reduction in spatial overlap between two adjacent molecules. However, as discussed previously, the calculated values globally decrease with distance but in an oscillating way [61]. Again, this evolution fully reflects that of the corresponding transfer integrals. Once more, we emphasize that the balance between the number of bonding versus antibonding interactions in the electronic overlap between the wave functions of the two molecules dictates the magnitude of the transfer integral and thus of the hole mobility. Maxima are observed when one kind of interaction dominates and minima when a compensation occurs between them. We have also shifted every other molecule along the stacking axis by a distance ranging from 0.1 to 0.5 Å with respect to the initial situation where all the molecules are separated by 4 Å. In such an arrangement, the charge has to hop alternatively over a distance larger or shorter than 4 Å. For various magnitudes of the applied electric field, we obtain parallel variations that point to a reduction in the mobility when the displacement is amplified (Figure 1.1.12); the impact is much more pro-



21

1

Mobility (cm2/ Vs)

F = 106 V/cm F = 5 × 106 V/cm

4Å − Δx 4Å + Δx

0.1

0.01 0

0.1

0.2

0.3

0.4

0.5

Shift (Å)

FIGURE 1.1.12 Evolution of the hole mobility as a function of the shift applied to every single molecule over two in a direction parallel to the stacking axis.

nounced for large shifts. This behavior can be understood by comparing the transfer rates associated with the short and long hops. When the distance is smaller or longer than 4 Å, the transfer rate increases or decreases, respectively, as discussed in Section 1.1.5.3. However, since the evolution with respect to the initial distance is not symmetric, the time required to make two consecutive hops in the same direction (and thus to travel in all cases a distance of 8 Å) increases with the degree of translation. Interestingly, fluctuations in the separation as large as 0.5 Å reduce the mobility only by a factor of five.

1.1.5.5 INTRODUCTION

OF A

GAUSSIAN DISORDER

We generalize the simulations carried out in the previous section by introducing a Gaussian distribution of the intermolecular distances d along the cofacial stack, randomly among the pairs of interacting molecules. The corresponding transfer integrals are estimated from an analytical expression of the results obtained in Section 1.1.5.3; we further assume here that λs is not affected by variations in the intermolecular distances. The Gaussian distribution g(d) is centered around an average value d0 set equal to 4 Å and is characterized by a standard deviation σ set equal to 0.05, 0.1, and 0.2 Å, respectively, in the simulations:

g(d ) =

⎡ ( d − d0 )2 ⎤ exp ⎢ − ⎥ 2σ2 ⎦ σ 2π ⎣ 1

(1.1.18)

With this expression, 99% of the generated distances lie in the interval [d0 – 3σ, d0 + 3σ]. Figure 1.1.13 shows that, as expected, the mobility is reduced when the standard deviation increases, whatever the magnitude of electric field might be. However, this reduction is very moderate and points to the weak impact of such


22


1

Mobility (cm2/ Vs)

F = 106 V/cm F = 5 × 106 V/cm

0.1

0

0.05

0.1

0.15

0.2

σ (Å)

FIGURE 1.1.13 Evolution of the hole mobility as a function of the width of the Gaussian distribution of the intermolecular distances in the one-dimensional array of pentacene molecules.

types of fluctuations on the charge-transport properties. The calculated evolution is due to the fact that the transfer rate increases faster below 4 Å than it decreases above 4 Å.

1.1.6 CONCLUDING REMARKS A major message resulting from our discussions is that the amplitudes of the transfer integrals depend on both the relative positions of the interacting molecules/oligomers and the shape (bonding–antibonding pattern) of their frontier molecular orbitals. Thus, the transfer integral amplitudes can hardly be predicted from a simple examination of molecular packing. Fortunately for theoretical chemists, this underlines the useful role that quantum chemistry can play by providing a molecular-scale understanding of the charge-transport parameters in conjugated systems. However, the role of the interactions with lattice phonons, which modulate the transfer integrals and define the external reorganization energy, and of the induced electronic polarization and electrostatic effects, which introduce an offset between the electronic levels and hence another contribution to ΔG°, need to be better incorporated, as shown by a number of recent studies [59,68–70]. Work along these lines is necessary to be able to estimate carrier mobilities accurately.

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42. Chen, H.Y. and Chao, I., Effect of perfluorination on the charge-transport properties of organic semiconductors: Density functional theory study of perfluorinated pentacene and sexithiophene, Chem. Phys. Lett., 401, 539, 2005. 43. Devos, A. and Lannoo, M., Electron-phonon coupling for aromatic molecular crystals: Possible consequences for their superconductivity, Phys. Rev. B, 58, 8236, 1998. 44. Sun, S.-S. and Sariciftci, N.S., Organic photovoltaics: Mechanism, materials, and devices, Taylor & Francis: Boca Raton, FL, 2005. 45. Lemaur, V., da Silva Filho, D.A., Coropceanu, V., Lehmann, M., Geerts, Y., Piris, J., Debije, M.G., Van de Craats, A.M., Senthilkumar, K., Siebbeles, L.D.A., Warman, J.M., Brédas, J.L., and Cornil, J., Charge transport properties in discotic liquid crystals: A quantum-chemical insight into structure–property relationships, J. Am. Chem. Soc., 126, 3271, 2004. 46. Newton, M.D., Quantum chemical probes of electron-transfer kinetics: The nature of donor–acceptor interactions, Chem. Rev., 91, 767, 1991. 47. Li, X.-Y., Tang, X.-S., and He, F.-C., Electron transfer in poly(p-phenylene) oligomers: Effect of external electric field and application of Koopmans theorem, Chem. Phys., 248, 137, 1999. 48. Li, X.Y. and He, F.C., Electron transfer between biphenyl and biphenyl anion radicals: Reorganization energies and electron transfer matrix elements, J. Comp. Chem., 20, 597, 1999. 49. Li, Z.C., Xu, L., Sun, H., Xiao, Y.M., and Zhang, J., Investigation on performances of non-loss storage for cryogenic liquefied gas, Cryogenics, 44, 357, 2004. 50. Pati, R. and Karna, S.P., Ab initio Hartree–Fock study of electron transfer in organic molecules, J. Chem. Phys., 115, 1703, 2001. 51. Wolfsberg, M. and Helmholz, L., The spectra and electronic structure of the tetrahedral ions MnO–4, CrO–4, and ClO–4, J. Chem. Phys., 20, 837, 1952. 52. Pietro, W.J., Marks, T.J., and Ratner, M.A., Resistivity mechanisms in phthalocyanine-based linear-chain and polymeric conductors: Variation of bandwidth with geometry, J. Am. Chem. Soc., 107, 5387, 1985. 53. Paulson, B.P., Curtiss, L.A., Bal, B., Closs, G.L., and Miller, J.R., Investigation of through-bond coupling dependence on spacer structure, J. Am. Chem. Soc., 118, 378, 1996. 54. Liang, C.X. and Newton, M.D., Ab initio studies of electron transfer: Pathway analysis of effective transfer integrals, J. Phys. Chem., 96, 2855, 1992. 55. Palenberg, M.A., Silbey, R.J., Malagoli, M., and Brédas, J.L., Almost temperatureindependent charge carrier mobilities in liquid crystals, J. Chem. Phys., 112, 1541, 2000. 56. Voityuk, A.A., Rosch, N., Bixon, M., and Jortner, J., Electronic coupling for charge transfer and transport in DNA, J. Phys. Chem. B, 104, 9740, 2000. 57. Grozema, F.C., van Duijnen, P.T., Berlin, Y.A., Ratner, M.A., and Siebbeles, L.D.A., Intramolecular charge transport along isolated chains of conjugated polymers: Effect of torsional disorder and polymerization defects, J. Phys. Chem. B, 106, 7791, 2002. 58. Valeev, E.F., Coropceanu, V., da Silva Filho, D.A., Salman, S., and Brédas, J.L., Effect of electronic polarization on charge-transport parameters in molecular organic semiconductors, J. Am. Chem. Soc., 128, 9882, 2006. 59. Senthilkumar, K., Grozema, F.C., Bickelhaupt, F.M., and Siebbeles, L.D.A., Charge transport in columnar stacked triphenylenes: Effects of conformational fluctuations on charge transfer integrals and site energies, J. Chem. Phys., 119, 9809, 2003.


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Organic Field-Effect Transistors 60. Mattheus, C.C., Polymorphism and electronic properties of pentacene, Ph.D. dissertation, University of Groningen, Netherlands, 2002. 61. Brédas, J.L., Calbert, J.P., da Silva, D.A., and Cornil, J., Organic semiconductors: A theoretical characterization of the basic parameters governing charge transport, Proc. Natl. Acad. Sci. USA, 99, 5804, 2002. 62. Kazmaier, P.M. and Hoffmann, R., A theoretical study of crystallochromy. Quantum interference effects in the spectra of perilene pigments, J. Am. Chem. Soc., 116, 9684, 1994. 63. Podzorov, V., Menard, E., Borissov, A., Kiryukhin, V., Rogers, J.A., and Gershenson, M.E., Intrinsic charge transport on the surface of organic semiconductors, Phys. Rev. Lett., 93, 086602, 2004. 64. Olivier, Y., Lemaur, V., Brédas, J. L., and Cornil, J., Charge hopping in organic semiconductors: Influence of molecular parameters on macroscopic mobilites in model one-dimensional stacks, J. Phys. Chem. A, 110, 6356, 2006. 65. Sancho-Garcia, J.C., Horowitz, G., Brédas, J.L., and Cornil, J., Effect of an external electric field on the charge transport parameters in organic molecular semiconductors, J. Chem. Phys., 119, 12563, 2003. 66. Lemaur, V., Steel, M., Beljonne, D., Brédas, J.L., and Cornil, J., Photoinduced charge generation and recombination dynamics in model donor/acceptor pairs for organic solar cell applications: A full quantum-chemical treatment, J. Am. Chem. Soc., 127, 6077, 2005. 67. Facchetti, A., Yoon, M.H., and Marks, T.J., Gate dielectrics for organic field-effect transistors: New opportunities for organic electronics, Adv. Mater., 17, 1705, 2005. 68. Johansson, A. and Stafstrom, S., Interchain charge transport in disordered π-conjugated chain systems, Phys. Rev. B, 66, 085208, 2002. 69. Troisi, A. and Orlandi, G., Charge-transport regime of crystalline organic semiconductors: Diffusion limited by thermal off-diagonal electronic disorder, Phys. Rev. Lett., 96, 086601, 2006. 70. Leontyev, I.V. and Tachiya, M., The reorganization energy of electron transfer in nonpolar solvents: Molecular level treatment of the solvent, J. Chem. Phys., 123, 224502, 2005.

8080_C002.fm Page 27 Friday, April 6, 2007 10:21 AM

2.1

Charge Carrier Transport in Single-Crystal Organic Field-Effect Transistors

Vitaly Podzorov CONTENTS 2.1.1 Introduction: The Field Effect in Small-Molecule Organic Semiconductors............................................................................................28 2.1.2 Fabrication of Single-Crystal OFETs..........................................................30 2.1.3 Charge Transport on the Surface of Organic Single Crystals ....................38 2.1.3.1 Basic FET Operation.....................................................................38 2.1.3.2 The Multiple Trap-and-Release Model.........................................46 2.1.3.3 Anisotropy of the Mobility ...........................................................48 2.1.3.4 Longitudinal and Hall Conductivity in Rubrene OFETs..............50 2.1.3.5 Comparison with the Holstein–Peierls Model and Transport Measurements in the Bulk of Organic Crystals............................54 2.1.3.6 Tuning the Intermolecular Distance..............................................55 2.1.3.7 Surface versus Bulk Transport ......................................................56 2.1.3.8 Photoinduced Processes in Single-Crystal OFETs.......................58 2.1.4 Defects at the Surface of Organic Crystals .................................................59 2.1.4.1 Bulk and Surface Electronic Defects in Organic Crystals ...........61 2.1.4.2 Density of Defects in Single-Crystal OFETs ...............................63 2.1.4.3 Single-Crystal OFETs as Tools to Study Surface Defects...........64 2.1.5 Conclusion ...................................................................................................65 Acknowledgments....................................................................................................67 References................................................................................................................67

27


28


2.1.1 INTRODUCTION: THE FIELD EFFECT IN SMALL-MOLECULE ORGANIC SEMICONDUCTORS Organic semiconductors represent a large class of solids consisting of organic oligomers or polymers. This chapter focuses on crystals of small organic molecules (mostly, polyacenes containing typically 2–10 benzene rings) held together in a solid by van der Waals forces. These small-molecule organic semiconductors, together with polymers, represent the material basis for the rapidly developing field of organic electronics [1–5]. Due to the weak van der Waals bonding, many electronic properties of these materials (e.g., the energy gap between the highest occupied and lowest unoccupied molecular orbitals — HOMO and LUMO, respectively) are determined by the structure of an isolated molecule [6–8]. The weak intermolecular overlap of electronic orbitals results in narrow electronic bands (a typical bandwidth, W ~ 0.1 eV, is two orders of magnitude smaller than that in silicon), a low mobility of carriers (μ ~ 1–10 cm2/Vs at room temperature) and strong electron-lattice coupling. The anisotropy of the transfer integrals between adjacent molecules reflects the low symmetry of molecular packing in organic molecular crystals (OMCs). It is believed that the most adequate description of charge transport in these semiconductors is based on the concept of polarons — the electronic states resulting from interaction of charge with lattice polarization at a length scale comparable to or greater than the lattice constant [6,7,9,10]. After several decades of intensive research, our basic understanding of charge transport in small-molecule organic semiconductors remains limited. The complexity of transport phenomena in these systems is due to the polaronic nature of charge carriers and a strong interaction of small polarons with defects [6]. An especially challenging task is to develop an adequate model of high-temperature polaronic transport. At room temperature, which is typically comparable to or even higher than the characteristic phonon energies, the lattice vibrations might become sufficiently strong to destroy the translational symmetry of the lattice. In this regime, the fluctuation amplitude of the transfer integral becomes of the same order of magnitude as its average value [11], the band description breaks down, and a crossover from the band-like transport in delocalized states to the incoherent hopping between localized states is predicted with increasing temperature. At low enough temperatures (T), when the band description is still valid, the polaronic bandwidth, W, “shrinks” as T increases, leading to a decrease of the carrier mobility μ with T [12–18]. The benchmark for the study of charge transport in organic semiconductors was established by time-of-flight (TOF) experiments with ultrapure polyacene crystals, such as naphthalene and anthracene [19]. These experiments have demonstrated that the intrinsic (not limited by static disorder) charge transport can be realized in the bulk of these crystals. This transport regime is characterized with a rapid growth of the carrier mobility with decreasing temperature and a pronounced anisotropy of the mobility, which reflects the anisotropy of the intermolecular transfer integrals [12,13]. Numerous applications, however, are dependent on the charge transport on



29

the surface of organic semiconductors. The most important example is the organic field-effect transistor, in which field-induced charges move along the interface between an organic semiconductor and a gate dielectric. In these devices, conduction truly occurs at the surface because the thickness of the conducting channel does not exceed a few molecular layers [20–22]. The transport of field-induced carriers on organic surfaces may differ from the bulk transport in many respects. For instance, the density of carriers in the fieldeffect experiments can exceed that in the bulk TOF measurements by many orders of magnitude, approaching the regime when the intercharge distance becomes comparable with the size of small polarons [23,24]. Interactions between the polaronic carriers may become important in this regime. The motion of charge carriers in the field-induced conduction channel might be affected by the polarization of the gate dielectric [25]. Finally, molecular packing at organic surfaces could be different from that in the bulk. The investigation of polaronic transport on organic surfaces is crucial for a better understanding of the fundamental processes that determine operation and ultimate performance of organic electronic devices. This is an important issue. On the one hand, the first all-organic devices (e.g., the active matrix displays based on organic light-emitted diodes and organic transistors) are expected to be commercialized within a few years. On the other hand, our knowledge of the transport properties of organic semiconductors is much more limited than it is for their inorganic counterparts. This paradoxical situation contrasts sharply to the situation in inorganic electronics in the mid-1960s, when the first Si metal-oxide semiconductor field-effect transistors (MOSFETs) were developed [26]. Difficulties in fundamental research have been caused by the lack of a proper tool for exploring the polaronic transport on surfaces of organic semiconductors. The most common organic electronic device, whose operation relies on surface transport, is the organic thin-film transistor (TFT). Over the past two decades a large effort in the development of TFTs has resulted in an impressive improvement of the characteristics of these devices [27] so that, currently, the best organic TFTs outperform the widely used amorphous silicon (α-Si:H) transistors. However, even in the best organic TFTs, charge transport is still dominated by the presence of structural defects and chemical impurities. As a result, it has been concluded that TFTs cannot be reliably used for the studies of basic transport mechanisms in organic materials [28]. The recently developed single-crystal organic transistors with significantly reduced disorder [29–37] provide unique opportunities to explore fundamental processes that determine the operation and reliability of organic electronic devices. For the first time, these single-crystal organic field-effect transistors (OFETs) have enabled the observation of intrinsic (not limited by static disorder) transport of fieldinduced charges at organic surfaces [35,38,39]. The carrier mobility in these devices is an order of magnitude greater than that in the best organic TFTs [32]. Equally important, the single-crystal OFETs are characterized by a very good reproducibility: Devices fabricated in different laboratories exhibit similar characteristics. This reproducibility, which is crucial for the investigation of electronic properties of organic semiconductors, has never been achieved with thin-film devices, whose electrical


30


characteristics are strongly dependent on the details of fabrication processes and handling environment. In this chapter, we present a brief overview of the experimental results obtained with single-crystal OFETs over the last four years. In Section 2.1.2, we briefly describe the crystal growth and OFET fabrication techniques that preserve the high quality of pristine surfaces of as-grown crystals. Section 2.1.3 focuses on the intrinsic transport characteristics of surfaces and interfaces of organic crystalline devices. Electronic defects at organic surfaces and mechanisms of their formation are discussed in Section 2.1.4. Section 2.1.5 outlines several basic issues that can be experimentally addressed in the near future owing to the availability of singlecrystal OFETs.

2.1.2 FABRICATION OF SINGLE-CRYSTAL OFETS The first step in the fabrication of single-crystal OFETs is the growth of ultrapure organic crystals. The best results to date have been obtained with physical vapor transport (PVT) growth in a stream of ultrahigh purity argon, helium, or hydrogen gases, similar to the method suggested by Laudise et al. [40]. A PVT furnace consists of a quartz tube with a stabilized temperature profile created along the tube by external heaters (Figure 2.1.1). The temperature gradient can be achieved by resistively heated wire unevenly wound on the quartz tube or, more conveniently, by coaxially enclosing the quartz tube in a metal tube (copper, brass, or stainless steel) with two regions of stabilized temperature: high T on the left and low T on the right (Figure 2.1.1). In such a design, the good thermal conductivity of the metal creates a linear temperature distribution between the heated and cooled regions, with the temperature (Tset) (Tgrowth) Water coil

Heater H2

Copper tube

Temperature profile

T (°C)

200 150 100 50 0

10

20 30 x (cm)

40

50

FIGURE 2.1.1 A sketch of the physical vapor transport (PVT) growth furnace (top) and an example of the temperature profile along the axis of the quartz tube (bottom).



31

gradient inversely proportional to the length of the metal tube. This helps to create a smooth temperature profile along the growth reactor. In addition, the outer metal tube reduces the access of ambient light to the organic material, which might be important in the case of photosensitive organic compounds. Starting material is loaded into the high-temperature zone and maintained at temperature Tset, where sublimation takes place, and molecules are carried by the gas stream into the region of lower temperature. For a given concentration of evaporated molecules, defined by the temperature Tset, there is a point located downstream at a lower temperature, Tgrowth, where crystallization occurs. At this point, the crystallization rate (proportional to the density of molecular vapor) becomes slightly greater than the rate of sublimation from crystal facets kept at temperature Tgrowth. Although both crystallization and sublimation occur at a facet simultaneously, the growth prevails and free-standing crystals grow. In the region to the left of the growth zone (upstream), sublimation prevails and no growth occurs; in the region to the right (downstream), the density of molecular vapor decreases and crystallization also does not occur. In this region, only smaller molecular weight impurities condense. If the temperature Tset is maintained very low (near the sublimation threshold of the material), heavier molecular impurities do not sublime and are retained in the load zone. This creates a 2- to 3-cm wide crystallization region that is typically well separated from the original material and from the impurities. Therefore, PVT process results in the crystal growth and material purification at the same time. For better separation of the crystals from the impurities, the temperature gradient along the tube should be sufficiently small (e.g., ~5–10°C/cm). Several factors affect the morphology and the quality of the grown crystals. Important parameters are, for instance, the temperature of the sublimation zone, Tset, and the carrier gas. For each material and each furnace, the optimal set of parameters has to be determined empirically. At least one common tendency has been observed for the common compounds, such as 7,7′,8,8′-tetracyanoquinodimethane (TCNQ), tetracene, and rubrene: The slower the growth process is, the higher the field-effect mobilities obtained in the resultant OFETs are. For this reason, Tset should be adjusted close to the sublimation threshold of the material. Typical Tset resulting in very slow growth of bulky crystals with large and flat facets is: 200, 210, and 300°C for the growth of TCNQ, tetracene, and rubrene, respectively. At such conditions, typical growth duration is 40–70 hours for 100–300 mg loaded material and H2 gas flow rate ~ 100 cc/min. Large, high-purity organic crystals can be obtained by the PVT technique (Figure 2.1.2). Most of the organic crystals are shaped as thin platelets or needles. The crystal shape is controlled by the anisotropy of intermolecular interactions: For many materials, the largest crystal dimension corresponds to the direction of the strongest interactions and, presumably, the strongest overlap of π-orbitals of adjacent molecules. For this reason, the direction of the fastest growth of elongated rubrene crystals (b axis) coincides with the direction of the highest mobility of field-induced carriers (see Section 2.1.3). In platelet-like crystals, the largest natural facet typically corresponds to the a–b plane. In-plane dimensions range from a few square millimeters


32


Rubrene

Tetracene

FIGURE 2.1.2 Single crystals of rubrene and tetracene grown from the vapor phase.

to a square centimeter. The crystal thickness also varies over a wide range and, in most cases, can be controlled by stopping the growth process at an early stage. For example, the thickness of the tetracene crystals grown for 24 hours ranges between ~10 and ~200 μm [41]; however, it is possible to produce crystals of submicron thickness by interrupting the growth after ~10 min. According to the atomic force microscopy (AFM) studies (Figure 2.1.3) [42], the slow crystal growth proceeds by the flow of steps at a very low growth rate (≤10 μm/hour in the direction perpendicular to the a–b facet) and results in molecularly flat facets with a low density of molecular steps, separated by relatively wide (0.5–1 μm) terraces. Several ultrahigh-purity gases have been used as a carrier agent. In de Boer, Klapwijk, and Morpurgo [31], the highest mobility of tetracene-based devices, μ = 0.4 cm2/Vs, was realized with argon, whereas other groups reported slightly higher mobilities in tetracene grown in hydrogen (0.8–1 cm2/Vs) [36,43]. The best reported mobilities in rubrene have been measured in the crystals grown in ultrahigh-purity

500 nm

1 μm

FIGURE 2.1.3 AFM images of the surface of uncleaved vapor-grown rubrene (left) and TCNQ (right) crystals, showing the molecular growth steps. The height of the steps is consistent with the lattice parameter along the c axis in these crystals.



33

(UHP) H2. The carrier gas might also influence the size and morphology of the crystals. For instance, growth of tetracene in helium gas yields very thin and wide crystals, inappropriate for fabrication of free-standing OFETs, though useful for lamination on hard substrates. On the other hand, slow growth in H2 or Ar yields much thicker (bulkier) and robust tetracene crystals that can be used to create free-standing devices using parylene gate dielectric with mobilities up to ~1 cm2/Vs [43]. Similar tendencies in morphology with the carrier gas have been observed for rubrene. At present, it is still unclear how, exactly, the transport gas affects the crystal quality, and more systematic studies are required. Poorly controlled factors such as parts per million (ppm) levels of water, oxygen, and other impurities in UHP gases that might create charge traps in organic material could complicate such studies. Another poorly controlled parameter is the purity of the starting material. Empirically, different grades of material with the same nominal purity might result in crystals of quite different quality (in terms of the field-effect mobility). Normally, the density of impurities can be greatly reduced by performing several regrowth cycles, in which previously grown crystals are used as a load for the subsequent growths. However, after a gradual increase of μ with the number of purifications, the mobility saturates already after two or three cycles. This indicates that some of the impurities cannot be effectively removed from the material or new defects might be forming during the growth process. Clearly, the higher the purity of the starting material is, the fewer regrowth cycles are required. In Podzorov et al. [30], the rubrene OFETs with μ > 5 cm2/Vs have been fabricated from the “sublimed grade” rubrene (Sigma–Aldrich) after only one or two growth cycles. Besides the growth from a vapor phase, other techniques, such as Bridgman growth from a melt or crystallization from a solution, can be used to produce organic crystals. For instance, vapor-Bridgman growth from a saturated vapor in a sealed ampoule has been used to grow large tetracene crystals for TOF studies (Figure 2.1.4) [44]. Crystallization from a solution usually results in mobilities substantially lower than those obtained in vapor-grown crystals. A clear demonstration of this has been recently obtained with OFETs based on single crystals of halogenated tetracene derivatives that are soluble in common organic solvents and can also be sublimed without decomposition [45]. OFET mobilities in the vapor-grown crystals were as high as 1.6 cm2/Vs, while the solution growth resulted in devices with μ ~ 10–3 cm2/Vs. A rare example of a high-mobility solution-grown crystalline system is dithiophene-tetrathiafulvalene (DT-TTF), with field-effect mobilities of up to 1.4 cm2/Vs [46]. It is worth noting that the mobility in single-crystal devices might be substantially improved if a zone refining process is used for prepurification of the starting material. Indeed, in the time-of-flight studies of organic crystals, the highest mobilities have been obtained after multiple cycles of zone-refinement purification. This process enabled reduction of impurity concentration in the bulk down to the part-per-billion level. It has to be noted that zone refinement cannot be applied to all organic materials, since this technique requires the existence of a coherent liquid phase (i.e., the melting temperature of a substance has to be lower than the temperature of decomposition of its molecules).


34


FIGURE 2.1.4 A tetracene crystal grown by the vapor-Bridgman technique and used for the time-of-flight studies. (From Niemax, J. et al., Appl. Phys. Lett., 86, 122105, 2005.)

X-ray diffraction studies show that most of the PVT-grown crystals are of excellent structural quality; they are characterized by a very small mosaic spread, typically, less than 0.050 [47] (in rubrene this value has been found to be even smaller, ~0.0160) [48]. Rubrene crystallizes in an orthorhombic structure with four molecules per unit cell and the lattice parameters a = 14.44 Å, b = 7.18 Å, and c = 26.97 Å [49] (crystallographic data for several other polyacenes have been reported in Campbell et al. [50]). The crystals are usually elongated along the b axis; the largest flat facet of the crystal corresponds to the (a,b)-plane. The possibility of surface restructuring or existence of a “surface phase” on free facets of organic crystals has not been addressed yet and remains to be studied experimentally. Deviations of the surface structure from the bulk phase that might be important for the charge transport in OFETs might occur similarly to the thin-film phase in monolayer-thick pentacene films studied by grazing incidence x-ray diffraction [51,52]. Recently, it has been demonstrated that scanning tunneling microscopy (STM) could be used to study the molecular organization at the surfaces of bulk crystals in certain cases of high-mobility systems, such as rubrene [53]. Figure 2.1.5 shows the first molecular-resolution STM image of the surface of a bulk organic crystal (rubrene) at room temperature. The common problem of charging of an insulating surface with the tunneling electrons is avoided here because the high mobility of carriers in rubrene facilitates fast removal of the tunneled electrons through the crystal into the conducting substrate. The surface quality of these crystals is unprecedented; the density of surface defects is very low, resulting in a low-noise image and the opportunity to resolve individual molecules at the surface. The packing motif observed with an STM at the surface is consistent with the bulk packing obtained by crystallography. It is a “herringbone” type of structure with a stack forming along the b axis.



35

a b a

b 2 mm

FIGURE 2.1.5 Scanning tunneling microscope images of a–b facet of a thick, as-grown rubrene crystal. The herringbone molecular organization at the surface, consistent with the bulk structure (shown in the lower left corner), is evident. (From Menard, E. et al., Adv. Mater., 18, 1552, 2006.)

Fabrication of field-effect structures on the surface of organic crystals is a challenge because the conventional processes of thin-film technology (such as sputtering, photolithography, etc.) introduce a high density of defects on vulnerable organic surfaces. For this reason, the first single-crystal OFETs have been realized only recently, after development of the two innovative fabrication techniques briefly described next. The first technique is based on the use of an unconventional gate dielectric: thin polymeric film of parylene, which can be deposited from a vapor phase on the surface of organic crystals at room temperature, producing a defect-free semiconductor–dielectric interface [29]. Conformal parylene coating is a well developed technology used commercially in electronic packaging applications [54]. A homebuilt setup for parylene deposition is depicted in Figure 2.1.6 (commercially available parylene coaters are not recommended for this research purpose because of their large volume and high cost). The reactor consists of a 20-mm ID quartz tube blocked at one end and a twozone furnace for sublimation and pyrolysis of the commercially available parylene dimers. The quartz tube extends from the high-T (700°C) section of the furnace by about 40 cm to the right; the sample(s) with prefabricated contacts and leads are placed in this portion of the tube, which is then connected to the mechanical pump through a liquid N2 trap. After evacuating the reactor to approximately 10–2 torr, the temperatures in the sublimation and pyrolysis zones are set to 100 and 700°C, correspondingly. Parylene dimers, sublimed at 100°C, split into monomers at 700°C and polymerize as they enter the room temperature section of the tube, producing a clear pinhole-free insulating coating on the sample’s surface.


36


CH2

CH2

CH2

CH2

CH2

2CH2

CH2

CH2 n

Di-para-xylylene (dimer)

Sublimation at 100 °C

Para-xylylene (monomer)

Pyrolysis at 700 °C

Poly(para-xylylene) (polymer)

Polymerization To liquid–N2 at 25 °C tap & mech. pump

Dimer powder

Sample

FIGURE 2.1.6 Reactions involved in parylene deposition process (top): sublimation of dimers at ~100°C, splitting into monomers at ~700°C, and polymerization at room temperature. Inexpensive system for parylene deposition (bottom) consists of a two-zone tube furnace and a 20-mm ID quartz tube containing parylene dimer powder and connected to a one-stage mechanical pump through a liquid nitrogen trap. A sample with prefabricated contacts and attached wire leads is placed in the tube at about 30 cm from the furnace.

The advantages of this technique are the following: • • • • •

•

High-energy charged particles, inherent to plasma-based deposition techniques (e.g., sputtering) and detrimental for the organic surfaces, are avoided. The sample is maintained at room temperature throughout the entire process. A high vacuum is not required, which, in combination with cheap parylene precursors, makes this a very low cost technique. The deposition process is fast: Growth of a 1-μm thick film lasts ~20 min. The physical properties of parylene films are remarkable: Parylene is a very good insulator with an electrical breakdown strength of up to 10 MV/cm, superior chemical stability, and high optical clarity. The coating is truly conformal, which allows working with crystals that have sharp features on their surfaces (e.g., steps, edges, etc.), without having problems with shorts.

The conformal properties of parylene coating are especially important in devices with colloidal graphite contacts that have rough surfaces. The parylene coating is the only technique available to date for the fabrication of free-standing single crystal OFETs. In comparison with the laminated devices, this has several advantages, such as elimination of substrate-related strains and a possibility to perform studies of photoinduced effects in OFETs by illuminating the conduction channel through the transparent parylene dielectric and a (semi)transparent gate electrode [55]. The



37

OFETs with parylene dielectric are very stable. For example, the characteristics of rubrene/parylene transistors remained unchanged after storing the devices for more than two years in air and in the dark. For the free-standing devices, deposition of metal contacts directly on the surface of organic crystal is necessary. While painting colloidal graphite contacts (e.g., Aquadag Colloidal Graphite, Ted Pella, Inc.) on the crystal surface works well and results in low contact resistance, it is difficult to create complex contact geometries and well-defined features with this simple technique. An alternative high-vacuum thermal evaporation of metals through a shadow mask is challenging due to several factors related to generation of defects at the organic surfaces as a result of (1) infrared radiation from evaporation sources; (2) contamination of the channel area with metal atoms able to penetrate under the shadow mask; and (3) interaction of organic surfaces with free radicals produced by hot filaments and high-vacuum gauges (the gauge effect) [56]. Nevertheless, devices with evaporated contacts have been successfully fabricated in Podzorov et al. [30] by using an optimized deposition chamber — a technique that might be useful for more complex contact geometries in OFETs, such as four-probe or Hall geometry. In the second technique of single-crystal OFET fabrication, the transistor circuitry is prefabricated by conventional microfabrication (lithography) methods on a substrate (this structure can be called a “stamp”), and the organic single crystal is subsequently laminated to it. This technique eliminates the need for deposition of metal contacts and dielectrics directly onto organic crystals. Hard inorganic (e.g., Si) and flexible elastomeric (polydimethylsiloxane = PDMS) stamps have been used for this purpose. In the first case, a heavily doped Si wafer with a thermal SiO2 plays the role of an insulated gate electrode [31,33]. After the deposition of gold contacts, a thin organic crystal can be laminated to such a stamp owing to van der Waals attraction forces. Similarly, field-effect transistor (FET) structure can be fabricated using PDMS substrates and spin-coated PDMS films [35]. The elastomeric stamps compare favorably with the Si stamps in two respects. First, the slightly conformal properties of PDMS enable establishing a good contact even with crystals that are not perfectly flat. Conversely, the use of hard Si stamps is restricted to perfectly flat crystals or to very thin and “bendable” crystals that could conform to hard substrate. Second, for the robust and bulky crystals such as rubrene, the PDMS stamps provide a unique opportunity to re-establish the contact many times without breaking the crystal and without degradation of the crystal’s surface. However, the achievable density of field-induced charges is typically greater in the Si-based stamps, especially if these stamps utilize high-ε gate insulators [23,57]. This is important for the exploration of the regime of high carrier densities, in which novel electronic phases might emerge (see Section 2.1.3.7). Even though the lamination of crystals on prefabricated substrates enables a “low-impact” probing of charge transport on organic surfaces, this impact may still be too strong for chemically reactive organic materials (e.g., a strong electron acceptor TCNQ). To minimize these effects and to preserve the pristine surface of organic crystals, modification of the PDMS stamping technique has been recently introduced that avoids these complications simply by eliminating the direct contact between the crystal and the gate dielectric [39]. The idea of these so-called vacuum-


38


PDMS stamp

Pr

Peel back; flip over; coat with Ti/Au; laminate crystal

Si

Single Crystal

Ti/Au

D

S G

FIGURE 2.1.7 Casting and curing procedures for fabrication of the “air-gap” transistor stamps. The recessed gate electrode is separated from the conductive channel by a micronsize gap. (From Menard, E. et al., Adv. Mater., 16, 2097, 2004.)

gap stamps is illustrated in Figure 2.1.7. In these devices, the conventional dielectric is replaced by a micron-size gap between the gate electrode and the surface of the organic semiconductor. A thin layer of gas (e.g., air) or vacuum between the bottom surface of the crystal and the recessed gate electrode plays the role of gate dielectric. This approach eliminates surface defects introduced in the process of lamination and enables studies of the effect of different gases and other environmental agents on the conduction channel in OFETs [56].

2.1.3 CHARGE TRANSPORT ON THE SURFACE OF ORGANIC SINGLE CRYSTALS In this section, after a brief introduction of the OFET operation principles, we outline the main signatures of the intrinsic polaronic transport observed in the experiments with single-crystal OFETs. We compare them to the results of TOF and space-charge limited current (SCLC) experiments that probe charge transport in the bulk.

2.1.3.1 BASIC FET OPERATION Contemporary OFETs are based on undoped organic semiconductors, and mobile charges in these devices must be injected from the metallic contacts. These devices can potentially operate in the electron- and the hole-accumulation modes, depending on the polarity of the gate voltage (the so-called ambipolar operation). Often, however, the injection barrier at the contact or the field-effect threshold for either n- or



39

Vg EVAC

G

S ϕM

VSD

D ISD

EHOMO(Vg < 0) EF ϕB

Metal

EHOMO(Vg = 0)

Semiconductor

FIGURE 2.1.8 Schematic energy diagram of a metal–organic interface (the contact). Evac is the vacuum energy level, EF is the Fermi energy of metal, EHOMO is the energy of the band edge of the semiconductor. The inset shows a two-probe OFET circuitry: S and D are the source and drain contacts, G is the gate electrode; Vg and VSD are the gate and source-drain voltages, respectively, and ISD is the source-drain current.

p-type conductivity is so large that an FET operates in a unipolar mode. For this reason, we will mainly discuss p-type conductivity, which is more commonly observed in OFETs. We will start the discussion with charge injection from contacts. An energy diagram of a hole injecting metal-semiconductor contact and a generic field-effect transistor circuit are schematically shown in Figure 2.1.8. The hole injection occurs through the interfacial Schottky barrier of height ϕB; the formation of the barrier is a complex process that depends on the metal work function ϕM, the ionization energy of the semiconductor, and the interfacial dipole moment formed due to a charge transfer at the interface. For a comprehensive review of the energetics of metal– organic interfaces, see, for example, the paper by Cahen et al. [58]. While the maximum height of the barrier, ϕB, remains fixed due to the pinning of energy levels at the interface, its width can be modified by an external electric field associated with either VSD or Vg. Figure 2.1.8 shows that when a negative Vg is applied, the effective width of the Schottky barrier for hole injection decreases. This results in a decrease of the contact resistance (RC), which depends on the barrier height ϕB, its effective thickness, and temperature. The “triangular” shape of the Schottky barrier allows the carrier injection via thermally activated excitation above the barrier and via tunneling under the barrier (the latter process does not require thermal excitation, but it is limited by Vgdependent barrier width). Both processes are possible, and the resultant injection mechanism, called thermionic emission, typically causes an exponentially fast increase of the contact resistance RC with lowering T and a strong dependence of RC on the gate voltage. The contact resistance enters the equations of OFET operation because the source-drain circuit is represented by two resistors connected in series — the contact


40


resistance, RC, and the channel resistance, RCH — so that the total source-drain resistance is RSD ≡ VSD/ISD = RC + RCH. The Schottky contact resistance in OFETs is typically high; in many cases, RC ≥ RCH, especially in short-channel TFTs, and the operation of such devices is contact limited. This is also the case for shortchannel single-crystal OFETs. However, if the channel is long enough (0.2–5 mm), rubrene and tetracene devices with either graphite or laminated gold/PDMS contacts are not contact limited at room temperature. The devices with evaporated silver or gold contacts generally have higher contact resistances. Nevertheless, RC typically decreases with VSD, and at large enough VSD and Vg < 0, even the devices with evaporated contacts are not dominated by contacts. However, it is important to be able to directly measure RC and RCH independently in each individual OFET in the entire temperature range of interest, because of a strong T-dependence of the contact resistance. In the prior studies, contact resistance was estimated by measurements of twoprobe OFETs and fitting the data with the Schottky model at different T and Vg [59] and by performing the channel length scaling analysis [60,61]. While these methods provide useful information about the contacts of a particular system, they do not allow for the direct and model-independent measurement of the contact resistance in each individual device. Recently introduced OFETs with four-probe contact geometry (source, drain, and two voltage probes in the channel) can be used to address this problem [29,30]. Before describing the four-probe measurements, let us introduce the operation of a conventional two-probe OFET, assuming that the contact resistance is negligible compared to the channel resistance. (This is practically valid for some cases of single-crystal OFETs with long channels at room temperature.) With an increase of the gate voltage |Vg| towards the threshold value |Vgth|, the carriers injected from the metallic contacts fill localized in-gap states of the organic semiconductor, associated with impurities and defects in the channel, whose energy is separated from the edge of the HOMO band by more than a few kBT (the deep traps; see Figure 2.1.9) (this simplified model assumes the existence of the HOMO band; this assumption may be violated at high temperatures) [12,62]. As the result, the Fermi level at the organic surface, E F, initially positioned within the HOMO–LUMO gap, approaches the edge of the HOMO band, EHOMO, which corresponds to the zero energy in Figure 2.1.9. As soon as EF – EHOMO becomes smaller than ~kBT, the OFET’s conductance increases by several orders of magnitude due to the thermal excitation of the carriers from the localized states into the HOMO band. As a result, a conduction channel is formed at the interface between the semiconductor and the gate dielectric. Overall device operation depends, to a large extent, on the energetics of the semiconductor bands and metal contacts, and therefore studies of the electronic structure of molecular interfaces are important [58]. Figure 2.1.10 shows the transconductance characteristics (i.e., the dependence of the source-drain current on the gate voltage, ISD(Vg), measured at a constant sourcedrain voltage, VSD) and ISD(VSD) characteristics typical for the p-type rubrene singlecrystal OFETs [29,30,39]. The channel conductance per square,



41

ε(eV) LUMO

2.2

~0.1 Deep traps Shallow traps

Few kBT

0 HOMO ν(ε)

FIGURE 2.1.9 The schematic diagram of the energy distribution of localized electronic states in the energy gap between the HOMO and LUMO bands in the rubrene single-crystal OFETs. (From Podzorov, V. et al., Phys. Rev. Lett., 93, 086602, 2004.)

μb = 12.3 cm2/Vs

ISD(μA)

3

VSD= 5 V

2

1 μa = 5.2 cm2/Vs

0

−50

−25 VG(V)

2.0

0

VG= −30 V

−ISD(μA)

1.5 1.0

−25 V

0.5

−20 V −15 V

0.0

−10 V 0

10

20 −VSD (V)

30

FIGURE 2.1.10 The transconductance ISD(Vg) (the upper panel) and ISD(VSD) (the lower panel) characteristics of rubrene single-crystal OFET (see, for example, Podzorov et al. [30] and Menard et al. [39]).


42


σ≡

I SD L , VSD W

increases linearly with Vg at |Vg| > |Vgth|. (Here, L and W are the length and width of the conduction channel, respectively.) This indicates that the carrier mobility [63]

μ≡

σ ⎛ 1 ⎞ ⎛ dI SD ⎞ L = en ⎜⎝ C iVSD ⎟⎠ ⎜⎝ dVg ⎟⎠ W

(2.1.1)

does not depend on the density of carriers field-induced above the threshold

(

)

n = C i Vg − Vgth / e

(2.1.2)

Here, Ci is the capacitance per unit area between the gate electrode and the conduction channel. A density-independent μ has been observed in devices based on single crystals of rubrene [30,35,36], pentacene [33,34], tetracene [31], and TCNQ [39]. This important characteristic of single-crystal OFETs contrasts sharply with a strongly Vg-dependent mobility observed in organic TFTs [64] and α-Si:H FETs [65]. In the latter case, the density of localized states within the gap is so high that the Fermi level remains in the gap even at high |Vg| values. The observation of Vg-independent mobility in single-crystal OFETs suggests that the charge transport in these structures does not require thermal activation to the mobility edge and the mobile field-induced carriers occupy energy states within the HOMO band. This is consistent with an increase of the mobility with cooling observed in high-quality single-crystal OFETs (see Section 2.1.3.4). (For comparison, μ decreases exponentially with lowering temperature in organic and α-Si:H TFTs.) The pronounced difference in the Vg- and T-dependences of the mobility in these two types of devices clearly indicates that the theoretical models developed for the charge transport in α-Si:H and organic TFTs [65] are not applicable to singlecrystal OFETs. In the cases when contact resistance is not negligible, four-probe OFETs are used to measure the channel and the contact resistances independently. In the fourprobe OFET geometry (Figure 2.1.11), in addition to ISD, voltage between a couple of extra probes located in the middle of the channel, V4w, can be measured as a function of Vg, VSD, and T. Gate voltage dependences of ISD and V4w for a typical four-probe OFET are shown in the upper panel of Figure 2.1.12. In the four-probe geometry, conductivity of the section of the channel between the voltage probes per square is: σ≡

I SD D V4w W



43

D

Vg

W L S

V4w

G

D

VSD ISD

FIGURE 2.1.11 Four-probe field-effect transistor with measurement circuitry: V4w is the voltage measured between the two voltage probes located in the middle of the channel. The inset shows the channel geometry: L and W are the channel length and width; D is the distance between the voltage probes.

Note that the voltage V4w does not necessarily remain constant when Vg or T is varied because the total VSD voltage applied to the device is distributed between the channel resistance, RCH, and the contact resistance, RC, both of which vary with Vg and T. Using the relationship σ = enμ, and n from Equation (2.1.2), we obtain for the contact-resistance-corrected channel mobility, μ4w: ⎛ 1 ⎞ ⎛ d ( I SD / V4w ) ⎞ D μ 4w = ⎜ ⎟W ⎜ dVg ⎝ C i ⎟⎠ ⎝ ⎠

(2.1.3)

and for the contact resistance RC: RC =

VSD L V4w − I SD D I SD

(2.1.4)

Typical Vg- and VSD-dependences of the contact resistance, normalized to the channel width, RCW, are shown in the lower panel of Figure 2.1.12. In agreement with the Schottky model, RC for a p-type device decreases with a positive VSD applied to the hole-injecting source contact and with a negative Vg applied to the gate. Interestingly, the relatively large magnitude of the contact resistance in Figure 2.1.12 (≥100 kΩ⋅cm) can be greatly reduced down to 1–2 kΩ⋅cm by treating the gold contacts with trifluoromethylbenzenethiol before the crystal lamination or by using nickel instead of gold, which has been recently reported to result in a remarkably low contact resistance ~ 0.1–0.4 kΩ⋅cm [61]. For several important applications in plastic optoelectronics, including the possibility of electrically pumped organic lasers, it would be very important to achieve an ambipolar operation in OFETs, with high electron and hole mobilities. Gatecontrolled electroluminescence from organic small-molecule thin-film transistors


44


2 10

1

5

V4w

2

ISD 0

0

400 RCW (kΩcm)

RCW (kΩcm)

107

105

V4w (V)

ISD (μA)

7

300 200 100 10

20 30 VSD (V)

103

101

−75

−50

−25 Vg (V)

40

0

25

FIGURE 2.1.12 Four-probe OFET characteristics, ISD(Vg) and V4w(Vg) (top), and the corresponding contact resistance RCW(Vg) (bottom). The inset shows the dependence of the contact resistance on VSD. (From Sundar V. C. et al., Science, 303, 1644, 2004.)

and, more recently, from single-crystal OFETs based on thiophene/phenylene cooligomers has been observed (see, for example, Nakamura at al. [66] and references therein). However, because these devices have not been optimized yet, hole and electron currents were not balanced, and only unipolar (p-type) electrical characteristics have been observed. Interestingly, Vg-controlled electroluminescence and ambipolar characteristics have been recently observed in conjugated polymer OFETs [67,68], which indicates a balanced electron and hole injection. However, low hole and electron mobilities (~10–3 cm2/Vs), typical for polymer semiconductors, limit the channel current and therefore may present a serious problem for realization of electrically pumped polymer lasers. For this reason, ordered small-molecule organic semiconductors with higher mobilities are very promising for research in this direction.


4-probe channel conductivity, σ = IDS/V4W (μS)


0.05

45

With “vac-gap” stamp μa4w ~ 1.6 cm2/Vs Ci = 0.19 nF/cm2

0.04

L

0.03

0.02

D

0.01

0.00

−20

−10

0

10 20 Vgate (V)

30

40

50

60

FIGURE 2.1.13 Channel conductivity along the a-axis of TCNQ single crystal measured in the “vacuum-gap” OFET. The mobility of n-type carriers is 1.6 cm2/Vs. (From Menard, E. et al., Adv. Mater., 16, 2097, 2004.)

Most of the small-molecule organic FETs operate in the p-type mode, and examples of n-type operation with high mobility are rare [39,69]. This “asymmetry” between n- and p-type carriers is due to several factors: the HOMO bandwidth is typically larger than the LUMO bandwidth [17], a stronger trapping of n-type polarons [19], and a larger Schottky barrier for electron injection into organic semiconductors from the most commonly used high work function metals. Figure 2.1.13 illustrates the n-type operation in a single-crystal TCNQ transistor. The surface of TCNQ, a semiconductor with a very high electron affinity, can be easily damaged (e.g., a direct contact of the crystal with PDMS dielectric in the contact stamps, such as those used in Sundar et al. [35], results in a very poor transistor performance with electron mobilities ~ (2–3)·10–3 cm2/Vs). The “air-gap” PDMS stamps [39] help to solve the problem. The observed carrier mobility ~ 1.6 cm2/Vs in the linear regime is significantly higher than in most of the n-channel organic TFTs. This value, however, is still limited by trapping (see Section. 2.1.3.2); more work is required to approach the fundamental limit of performance of n-type OFETs. In practice, realization of high-mobility ambipolar operation is a challenge because two difficult problems must be solved simultaneously: (1) the density of both n- and p-type traps should be minimized at organic/dielectric interfaces; and (2) an effective injection of both n- and p-type carriers from the contacts into the organic semiconductor must be realized. Among inorganic FETs, only devices based on carbon nanotubes [70] and single crystals of transition metal dichalcogenides (e.g., WSe2 and MoSe2) [71] demonstrated high-mobility ambipolar operation. The number of organic materials in which the ambipolar operation has been demonstrated is limited as well [67,69,72,73]. The organic single crystals, with their intrinsically


46


low density of traps, offer a unique opportunity to realize the ambipolar operation with a relatively high mobility of both types of carriers. Ambipolar operation has been recently observed in the single-crystal OFETs based on metal phthalocyanines (MPc), namely, FePc and CuPc [74] (Figure 2.1.14) and rubrene [75]. Because of the reduced density of electron traps at the interface and a relatively small HOMO–LUMO gap in the case of MPc ( |Vgth|) contribute to the current flow at any given moment of time. Some of the mobile charges can be momentarily trapped by shallow traps; the number of these charges depends on the density of shallow traps and temperature. Within the MTR model, the effect of trapping can be described using two approaches. In the first approach, one can assume that all carriers field induced above


48


the threshold, n (see Equation 2.1.2), contribute to the current flow at any moment of time, but their effective mobility μeff is reduced in comparison with its intrinsic, trap-free value μ0:

μ eff = μ 0 (T )

τ(T ) τ(T ) + τtr (T )

(2.1.5)

Here, τtr(T) is the average trapping time on shallow traps and τ(T) is the average time that a polaron spends diffusively traveling between the consecutive trapping events. In the alternative second approach, one can assume that only a fraction of the carriers field induced above the threshold voltage are moving at any given moment of time:

neff = n

τ(T ) τ(T ) + τtr (T )

(2.1.6)

However, these charges are moving with the intrinsic trap-free mobility μ0. These two approaches are equivalent for describing the channel conductivity σ = enμ, which depends only on the product of n and μ. The distinction between these approaches becomes clear in the Hall effect measurements, in which the density and the intrinsic mobility of truly mobile carriers can be determined independently (see Section 2.1.3.4). According to Equation 2.1.5, the intrinsic regime of conduction is realized when τ >> τtr. In this case, the dependence σ(T) reflects the temperature dependence of intrinsic mobility μ0(T). In the opposite limit τ VG . The current in both regimes is given by Equations (2.2.1) and (2.2.2) [18]: I Dlin =

⎡ Z V2 ⎤ Ciμ ⎢(VG − VT ) VD − D ⎥ 2 ⎦ L ⎣

(2.2.1)

Z 2 Ciμ (VG − VT ) 2L

(2.2.2)

I Dsat =

Here, C i is the capacitance of the insulator, μ the mobility in the semiconductor, and VT the threshold voltage. The meaning of the latter parameter will be detailed later. In short, VT is the gate voltage beyond which the conducting channel forms.


Charge Transport in Oligomers

77

Apart from showing the performance of the device, I–V curves are used to extract its basic parameter, primarily the mobility and threshold voltage. A widely used method for parameter extraction consists of plotting the square root of the saturation current as a function of the gate voltage. As seen in Equation (2.2.2), this is supposed to give a straight line, the slope of which gives μ while its extrapolation to the VG axis corresponds to the threshold voltage. Equations (2.2.1) and (2.2.2) rests on the following assumptions: (1) The transverse electric field induced by the gate voltage is largely higher than the longitudinal field induced by the gate bias (so-called gradual channel approximation); (2) the mobility is constant all over the channel. Assumption (1) is justified by the geometry of the device; that is, the distance from source to drain is most often much larger than the thickness of the insulator. Assumption (2) is more or less fulfilled in a conventional inorganic semiconductor. However, this is far from true in organic solids, as will be shown in this chapter. For this reason, the use of Equation (2.2.2) to extract the mobility may lead to an incomplete, if not erroneous, description of charge transport in organic semiconductors. Alternative approaches to circumvent this difficulty will be presented in the following sections.

2.2.3 CHARGE TRANSPORT IN CONJUGATED OLIGOMERS In contrast to its parent elements of column IV of the periodic table (Si, Ge …), carbon presents the unique feature of being able to exist under three different hybridization configurations: namely, sp, sp2, and sp3. The latter one is found in the so-called saturated compounds that are the constituting element of plastics. In this configuration, each carbon atom is linked to its neighbors by four strong σ bounds that point to the four verges of a tetrahedron. Because σ bounds are so strong, the distance between the bonding and antibonding energy levels (also called highest occupied and lowest unoccupied molecular orbitals — HOMO and LUMO) is high, which has two consequences: Plastics are transparent to visible light, and they are electrically insulating. All the organic compounds designated as semiconductors are those made of sp 2 hybridized carbons, also called conjugated organic materials. Under such circumstances, each carbon is linked to its neighbors by three σ bonds resulting from the hybridization of 2s , 2 p x , and 2 py orbitals, while the remaining 2 pz orbital forms a π bond, which presents significantly less overlap than σ bonds. For this reason, the energy distance between the bonding and antibonding orbitals is somewhat reduced, which has two consequences: The materials absorb visible light (dyes are conjugated materials) and may behave as a semiconductor at nonzero temperature. This concept is illustrated in Figure 2.2.3 in the case of ethylene C2H2. In larger molecules, typically benzene, the π orbitals become delocalized and form a π system that extends all over the molecule. The HOMO–LUMO gap becomes smaller with increasing delocalization. In the case of a long chain of carbon atoms, the π bonds delocalize over the whole chain and form a one-dimensional electronic system. The resulting


78

Organic Field-Effect Transistors σ+ π+ pz

pz π−

sp2

sp2

σ−

FIGURE 2.2.3 Energy scheme of ethylene C2H2. . . n Conduction band (π∗) pz sp2

Valence band (π)

FIGURE 2.2.4 Molecular and energy schemes of poly-para-phenylene-vinylene (PPV).

one-dimensional band has substantial band width, and the chain can be viewed as a one-dimensional semiconductor with a filled valence band originating from the HOMO and an empty conduction band coming from the LUMO. Figure 2.2.4 illustrates this image with the molecule of poly-para-phenylene-vinylene (PPV). The image depicted in Figure 2.2.4 gives rationale for why charge carriers can be injected and reside in a conjugated molecule. However, the limiting step for charge transport in a solid is not within single molecules; rather, it involves electron transfer between molecules or molecular chains; because orbital overlap between molecules is low, the phenomenon of charge transport in conjugated solids requires further investigations.

2.2.3.1 BAND TRANSPORT Band transport refers to the mechanism occurring in crystalline inorganic solids like metals and semiconductors. Band theory can be found in numerous textbooks [1] and will not be detailed here. In short, energy bands in solids form because when a very large number of interacting atoms are brought together, their energy levels become so closely spaced that they become indistinct. Any solid has a large number of energy bands, but not all these bands are filled with electrons. The likelihood of any particular band to be filled is given by the Fermi–Dirac statistics, Equation



79

(2.2.3), so that at zero temperature, bands are filled up to the so-called Fermi energy E F . f (E ) =

1 1 + exp E −kTEF

(2.2.3)

On this basis, solids can be divided into insulators, in which the highest occupied band (the valence band) is completely filled, while the lowest unoccupied band (the conduction band) is completely empty and metals present a partly empty and partly filled band (the conduction band). Semiconductors are a particular case of insulators where the energy gap between the top of the valence band and the bottom of the conduction band is small enough that, at nonzero temperature, the smoothing out of the Fermi–Dirac distribution causes an appreciable number of states at top of the valence band to be empty and an equivalent number of states at bottom of the conduction band to be filled. Note that the conductivity in semiconductors is highly temperature dependent. The simplest model of charge transport in delocalized bands is the Drude model, which assumes the carriers are free to move under the influence of an applied electric field, but subject to collisional damping forces. Note that the scattering centers are not the nuclei of the background material, but rather phonons (lattice vibrations) or impurities. A statistical equation for estimating the mean drift velocity of the carriers in the direction of the electric field Fx may be written as d q 1 vx = ∗ Fx − vx dt m τ

(2.2.4)

where q is the elemental charge and m ∗ the effective mass. τ is the mean free time between two collisions (or relaxation time). Steady state corresponds to dtd v x = 0 . Under such circumstances, the solution of Equation (2.2.4) writes vx =

qτ Fx = μFx m∗

(2.2.5)

which defines the mobility μ . It is important to note at this stage that the model is only valid when the mean free path λ — that is, the mean distance between two collisions — is much larger than a characteristic distance that can be the de Broglie length of the charge carrier, or the distance between two atoms in the crystal. The mean free path is given by λ = vth τ

(2.2.6)

where vth = 3 kT / m ∗ is the electron thermal velocity (~105 m/s at room temperature). From Equations (2.2.5) and (2.2.6) the mobility can also be defined as


80


μ=

qλ m∗vth

(2.2.7)

The temperature dependence of the mobility depends on the nature of the scattering centers (acoustical or optical phonons, charged impurities …) However, in all cases, it is found that the dependence follows the general law given by Equation (2.2.8) μ(T ) ∝ T − n

(2.2.8)

In most practical cases, n is positive, so the mobility increases when temperature decreases. Evidence for “band transport” is often claimed to be brought when the temperature dependence in Equation (2.2.8) is observed. The most celebrated example for such a behavior is that by Karl and coworkers on highly pure crystals of acenes [19]. However, as pointed out by Silinsh and Cápek [20], the argument does not resist further analysis because, at least for temperatures above 100 K, the mean free path calculated from Equation (2.2.5) falls below the distance between molecules in the crystal, which is not physically consistent with diffusion limited transport; so the exact nature of charge transport in these crystals is still unresolved for the time being.

2.2.3.2 POLARON TRANSPORT 2.2.3.2.1 Polarization in Molecular Crystals The main reason why the band model is unable to account for charge transport in organic semiconductors is that it fails to account for a crucial phenomenon in these materials: polarization. The occurrence of polarization in organic solids has been analyzed in detail by Silinsh and Cápek [20]. The principle is the following. A charge carrier residing on a molecular site tends to polarize its neighboring region. As the resulting formed polarization cloud moves with the charge, the traveling entity is no longer a naked charge, but a “dressed” charge, and the formed species is called a polaron. In conjugated solids, the main polarization effect is that on the clouds formed by the π-electrons. The principle is illustrated in Figure 2.2.5, where the conjugated molecules are symbolized by benzene rings; the hexagons represent the (fixed) core of the six carbon atoms, while the circles stand for the delocalized π-electrons. Under the effect of the positive charge on the central molecule, the π-electron rings tend to move towards the central molecule, thus creating an electric dipole, the magnitude of which is the greater as the molecule is closer to the center. In order to estimate the stability of the polaron, we can define two typical times: (1) the residence time τres corresponds to the average time a charge will reside on a molecule; (2) the electronic polarization time τel is the time it takes for the polarization cloud to form around the charge. An order of magnitude for both times can be estimated by using Heisenberg’s uncertainty principle



81

+

FIGURE 2.2.5 The figure is a schematic representation of the formation of a polaron when a positive charge is placed on a molecule in a conjugated organic solid. The hexagons symbolize the core of the nuclei, while the circles represent the delocalized π-electrons.

τ< ~

ΔE

(2.2.9)

where ΔE is a characteristic energy. For the residence time, the pertinent energy is the width W of the allowed band, typically 0.1 eV in an organic semiconductor and 10 eV in an inorganic semiconductor, which gives a residence time of 10 −14 s for the former and 10 −16 s for the latter. As for the electronic polarization time, the corresponding energy is that of an electron transition — that is, the energy gap (~1 eV) — so the time of the order of 10 −15 s in both cases. Similarly to electronic polarization, other polarization mechanisms can be invoked: molecular polarization, which concerns the displacement of the nuclei of the molecule where the charge resides, and lattice polarization, which involves movements of the entire lattice. The energies and times corresponding to these processes are estimated from the intramolecular and lattice vibration frequencies. The energy and time of the various polarization processes are summarized in Table 2.2.1. Table 2.2.1 reveals a striking difference between inorganic and organic semiconductors. In the former, the localization time is shorter than all the possible polarization times. In other words, electrons and holes move so fast that the polarization cloud does not have enough time to form. This is actually included in the band theory through the so-called rigid-band approximation, which states that the band structure remains uncharged when a charge is injected in the solid. The situation is drastically different in organic materials. Here, the electronic polaron has long enough time to form, so the energy levels of a charged molecule are significantly shifted with respect to that of a neutral molecule, as shown in Figure 2.2.6. As for the molecular polarization, it forms in a time comparable to the residence time, so the situation varies from one compound to the other. Finally, the formation time of the lattice polaron is too long, so its occurrence is unlikely under all circumstances. The pertinent parameters in the energy diagram in Figure 2.2.6 are the polarization energies for positive P + and negative charge carriers P −. Both are composed


82


TABLE 2.2.1 Residence Time and Various Polarization Times

Residence τres Polarization

Energy (eV)

Time (s)

10 0.1 1 0.1 < 0.01

10–16 10–14 10–15 10–14 > 10–13

Inorganic SC Molecular SC Electronic τel Molecular τv Lattice τl

Note: The reference energy is bandwidth for residence time, energy gap for electronic polarization, molecular vibration (~1,000 cm–1) for molecular polarization, and lattice vibration ( λ , while the latter dominates when 2t < λ . A connection with the molecular polaron model developed before can be derived from the fact that the reorganization energy is linked to the the molecular polarization time, and the transfer integral to the residence time, so the first inequality can also write τres < τ v and the second one τres > τ v , which is precisely what was established previously. Table 2.2.3 gives values of calculated reorganization energy and transfer integral for the polyacene series [27]. It can be seen that localized transport is expected for naphthalene and anthracene and delocalized transport for pentacene; tetracene is located in between these two extreme cases.

2.2.3.3 HOPPING TRANSPORT The problem with hopping transport is that dozens of different models have been proposed, based on different physical principles and approximations. Moreover, most


86


of these models only give a qualitative description of charge transport, thus preventing the possibility of computational treatment. However, hopping models have proved useful in rationalizing charge transport in disordered materials, such as polymers. Time of flight (TOF) measurements have revealed that the carrier mobility in these organic materials is thermally activated. Another ubiquitous feature relates to the field dependence of μ obeying a ln μ ∝ F law, where F is the magnitude of the electric field. The current practice is to interpret this behavior in terms of Gill’s [28] or Poole–Frenkel (PF)-like equation − Δ0 −β F ) / kTeff μ = μ0e (

(2.2.18)

where 1 / Teff = 1 / T − 1 / T ∗ and β is the PF factor. The problem with Equation (2.2.18) is that it presents several physical inconsistencies, among which are the lack of physical meaning for the effective temperature and the fact that the actual values of the PF factor are far from that predicted by the conventional PF theory. The disorder model developed by Bässler [29] rests on the following assumptions: (1) Because of the randomness of the intermolecular interactions, the electronic polarization energy of a charge carrier located on a molecule is subject to fluctuations; (2) transport is described in terms of hopping among localized states; in analogy to optical absorption profiles, the DOS is described by a Gaussian distribution of variance σ ; (3) charge transport is random walk described by a generalized master equation of the Miller–Abrahams form [30]: ν = ν0 e−2 γΔRij e−Δεij / kT

(2.2.19)

where ΔRij is the intersite distance and Δεij the energy difference of the sites; and (4) in addition to the energetic disorder, there exists a position disorder with a Gaussian distribution of variance Σ (the so-called off diagonal disorder). From a Monte Carlo simulation, Bässler arrives at a universal law relating the mobility to the degree of both diagonal and off diagonal disorder: ⎫ ⎧ ⎡ ⎤ 2 ⎡ ⎛ 2 σ ⎞2⎤ ⎪⎪ ⎪⎪ ⎢⎛ σ ⎞ 2 ⎥⎥ ⎢ μ = μ 0 exp ⎢−⎜ ⎟ ⎥ exp ⎨⎪C ⎢⎜ ⎟ − Σ ⎥ F ⎬⎪ ⎝ ⎠ ⎢⎣ ⎝ 3 kT ⎠ ⎥⎦ ⎥⎦ ⎪⎭ ⎪⎩ ⎢⎣ kT

(2.2.20)

where C is an empirical constant.

2.2.4 TRAP LIMITED TRANSPORT IN ORGANIC TRANSISTORS In real organic transistors, charge transport is most of the time limited by localized states induced by defects and unwanted impurities. Clear evidence for such a process is given by the fact that the performance of the devices is strongly sample dependent.



Transport band Localized levels

87

E DOS

FIGURE 2.2.8 Principle of charge transport limited by multiple trapping and thermal release.

Two useful models for accounting for such a trend are the multiple trapping and thermal release (MTR) and variable range hopping (VRH) models. While hopping transport is appropriate to describe charge transport in disordered materials, the MTR model [31] applies to well-ordered materials, prototypes of which are vapor deposited small molecules like pentacene or the oligothiophenes, where thermally activated mobility is often observed. The basic assumption of the model is a distribution of localized energy levels located in the vicinity of the transport band edge. During their transit in the delocalized band, the charge carriers interact with the localized levels through trapping and thermal release (Figure 2.2.8). The model rests on the following assumptions: (1) Carriers arriving at a trap are instantaneously captured with a probability close to one; and (2) the release of trapped carriers is controlled by a thermally activated process. The resulting effective mobility μ eff is related to the mobility μ 0 in the transport band by an equation of the form μ eff = μ 0αe−( Ec −Et )/ kT

(2.2.21)

where Ec is the energy of the transport band edge. In the case of a single trap level of energy Et and density of state (DOS) N t , the total charge-carrier concentration ntot splits into a concentration of free carriers n f = N c e − ( Ec − E F )/ kT , where N c is the effective density of states at transport band edge, and a concentration of trapped carriers nt = N t e − ( Et − E F )/ kT . The ratio of trapped to total densities is given by [32] Θ=

nt 1 N = t e−( Ec −Et )/ kT nt + n f 1 + NNct e( Ec −Et )/ kT N c

(2.2.22)

In that instance, the effective mobility is μ eff = Θμ 0 , so that Et in Equation (2.2.21) is the energy of the single trap level and α the ratio of the trap DOS to the effective density of states at transport band edge. If traps are energy distributed, distribution-dependent effective values of Et and α must be estimated, as will be exemplified in the following. In all circumstances, whichever the actual energy distribution of traps, the main feature predicted by the MTR model is thermally activated mobility. An important outcome of the MTR model is that in the case of an energy distributed DOS, mobility is gate voltage dependent. The mechanism at work is schematically pictured in Figure 2.2.9.


88


VS

EC EF

DOS VG − VT = 0

VG − VT ≠ 0

FIGURE 2.2.9 Gate voltage dependent mobility induced by an energy distributed density of traps.

When a bias is applied to the gate, a potential Vs develops at the insulator–semiconductor interface, which results in shifting by the same amount the Fermi level towards the transport band edge, thus partly filling the distribution of localized states. Accordingly, the energy distance between the filled traps and the transport band edge is reduced, so trapped-carrier release is made easier, and the effective mobility increases. Such a gate-voltage dependence of mobility has indeed been reported on several occasions [32,33]. The shape of the gate voltage dependence depends on that of the DOS. We have seen before that in hopping models, the preferred DOS distribution is a Gaussian distribution. By analogy with what is found in hydrogenated amorphous silicon [31], the trap distribution used in the MTR model is an exponential band tail. This is because the trap distribution no longer corresponds to the transport band itself (as in the case of the hopping model); instead, the DOS is a tail to a delocalized transport band. Note, however, that an exponential tail distribution can also be associated to a Gaussian transport DOS. One of the major interests of the exponential distribution is that it leads to an analytical form of the gate voltage dependence of the mobility. The general form of an exponential distribution of traps is given by

Nt (E ) =

N t 0 −( Ec −E )/ kT0 e kT0

(2.2.23)

where N t0 is the total density (per unit area) of traps and T0 a characteristic temperature that accounts for the slope of the distribution. The previously defined trapped charge is connected to the density of traps through

nt = q

∫

+∞ −∞

N t ( E ) f ( E ) dE

(2.2.24)

where f ( E ) is the Fermi distribution. If N t ( E ) is a slowly varying function, the Fermi distribution can be approximated to a step function; that is, it equals zero for E < E F and one for E > E F . The integration of Equation (2.2.24) leads to



89

nt qN t ( EF 0 + qVs ) = nt 0 eqVs / kT0

(2.2.25)

As stated earlier, we have made use of the fact that the Fermi level E F is shifted towards the band edge Ec from the value E F 0 at zero gate bias by an amount qVs (see Figure 2.2.9). nt 0 = N t 0 e − ( Ec − E F 0 )/ kT0 is the density of trapped charge at zero gate voltage. Making use of the defined effective mobility and assuming n f 100 nm. In principle, for a given thickness of dielectric, a high-k dielectric is preferable to a low-k dielectric for an FET application, which requires the FET to exhibit a high drive current at low drive voltage. Various solution-processable high-k dielectrics for low-voltage OFETs have been used in the literature, such as anodized Al2O3 [72] (ε = 8–10), TiO2 [73] (ε = 20–41), or polyvinylphenol loaded with TiO2 nanoparticles [74] (for a review see Veres et al. [19]). Many polar, high-k polymer dielectrics, such as polyvinylphenol (ε = 4.5) or cyanoethylpullulan (ε = 12), are hygroscopic and susceptible to drift of ionic impurities during device operation and thus cannot be used for ordinary TFT applications [75]. Veres et al. have shown that the field-effect mobilities of amorphous PTAA [18] and other polymers are higher in contact with low-k dielectrics with ε < 3 than dielectrics with higher k [19]. The latter usually contain polar functional groups randomly oriented near the active interface, which is believed to increase the energetic disorder at the interface beyond what naturally occurs due to the structural disorder in the organic semiconductor film resulting in a lowering of the field-effect


114


mobility (Figure 2.3.3a). Low-k dielectrics also have the advantage of being less susceptible to ionic impurities, which can drift under the influence of the gate field, causing device instabilities (see Section 2.3.4). A range of techniques have been developed that allow fabrication of OFET sourcedrain electrodes with submicrometer channel length (see, for example, Hamadani and Natelson [76] and Sele et al. [77]). To ensure correct scaling of the device characteristics in such short channel devices, the dielectric thickness needs to be significantly thinner than the channel length. Ideally, the gate dielectric thickness should be one order of magnitude smaller than the channel length. In this way, saturation of the FET current when the gate voltage (corrected for the threshold voltage) exceeds the sourcedrain voltage can be achieved even for submicrometer channel lengths. Very thin, sub-20 nm organic dielectrics have been demonstrated using several approaches, such as self-assembled monolayer dielectrics [78], self-assembled molecular multilayers [79], or ultrathin polymer dielectrics [20]. Cross-linking of polyvinylphenol and polystyrene using bis(trichlorosilyl)alkyl reagents has been shown to result in improved dielectric properties and enable very thin spin-coatable polymer dielectrics [79]. Cho et al. have used self-assembled monolayers of docosyltrichlorosilane as the gate dielectric of a bottom-gate, top-contact P3HT FET with inkjet-printed conducting polymer source-drain electrodes [80]. The chemical purity and composition of the gate dielectric can have dramatic effects on interfacial charge transport. The reason for the absence of n-type fieldeffect conduction in “normal” polymers such as PPVs or P3HT with electron affinities around 2.5–3.5 eV has puzzled the community for some time because, in LED devices, many of these polymers support electron conduction. Chua et al. [56] have demonstrated that by using appropriate gate dielectrics free of electron-trapping groups, such as hydroxyl, silanol, or carbonyl groups, n-channel FET conduction is in fact a generic property of most conjugated polymers. In contact with trappingfree dielectrics such as BCB or polyethylene, electron and hole mobilities were found to be of comparable magnitude in a broad range of polymers. Some polymers, such as P3HT and OC1C10–PPV, even exhibit ambipolar charge transport in suitable device configurations (Figure 2.3.5). This demonstrates clean inversion behavior in organic semiconductors with band gaps > 2 eV. n-Type behavior has previously been so elusive because most studies were performed on SiO2 gate dielectrics for which electrochemical trapping of electrons by silanol groups at the interface occurs [56]. Light-emitting organic field-effect transistors (LEOFETs) have recently attracted much attention because they combine the switching characteristics of transistors with the light emission of diodes. Although several groups had reported lightemission from an OFET [81–85], no report of spatially resolved light emission from within the channel of an organic light-emitting FET had been made until recently. As a corollary to the realization of clean ambipolar transport in organic semiconductors at trap-free gate dielectric interfaces, light-emitting polymer field effect transistors with a well-defined recombination zone within the channel have recently been demonstrated [86,87]. Figure 2.3.6(a) shows a schematic diagram of an ambipolar OFET with a semiconducting layer of OC1C10–PPV in contact with BCB gate dielectric and two dissimilar source and drain contacts (Au and Ca) formed by an angled evaporation technique.


Charge Transport Physics

115 Si H3C CH CH3 CH3 3 Si Si O

10−6

10−6

10−7

10−7 Current (A)

Current (A)

Si

10−8 10−9 O

10−10 10−11 −10

0

10−8 10−9 10−10

n

MeO 10

20

30

40

50

60

Gate Voltage (V) (a)

n

10−11 −10

C6H13 ∗

S

0

10

n∗ 20

30

40

50

60

Gate Voltage (V) (b)

FIGURE 2.3.5 Transfer characteristics of bottom-gate OC1C10 PPV and P3HT FET with trap-free BCB gate dielectric exhibiting clean ambipolar transport (Vsd = 60V). (Courtesy of Jana Zaumseil, University of Cambridge.)

When such an ambipolar FET is biased with the gate voltage in between the source and the drain voltage, an electron accumulation layer is formed near one electrode coexisting with a hole accumulation layer near the other electrode. Electrons and holes can be observed to recombine where the two accumulation layers meet, leading to light emission from a well-defined zone, the position of which can be moved to any position along the channel by varying the applied voltages (Figure 2.3.6b). Since the semiconducting layer is unpatterned, light-emission can also be observed from the periphery of the device at distances of more than 500 μm from the edge of the electrodes (Figure 2.3.6c). The observation of a spatially resolved recombination in the channel provides an unambiguous proof of the coexisting electron and hole channels and the truly ambipolar nature of charge transport at such trap-free dielectric–organic semiconductor interfaces [86].

2.3.5 CHARGE TRANSPORT PHYSICS The electronic structure of conjugated polymer semiconductors reflects the complex interplay between intrinsic π-electron delocalization along the polymer backbone and strong electron–phonon coupling, and the existence of energetic and positional disorder in solution-processed thin films. In a hypothetical, infinitely straight polymer chain, the highest occupied molecular orbital (HOMO) and lowest unoccupied


116


Light A Ca Au

SiOx Holes

Electrons

Au Ca OC1C10-PPV

BCB (50 nm) Thermal SiO2 (300 nm) Gate (doped Si)

B

C

Au

Au

Ca

Ca

FIGURE 2.3.6 (A) Schematic diagram of bottom-gate, ambipolar light-emitting FET with an active semiconducting layer of OC1C10–PPV and BCB gate dielectric. (B) Photograph of light emission from within the channel of the FET (IFET = 30 nA, Vg = –75V). (C) Photograph of light emission from periphery of the device illustrating the spreading of both electron and hole accumulation layers into the unpatterned semiconductor region around the source-drain electrodes. The channel length is 80 μm. (From Zaumseil, J. et al., Nat. Mater., 5, 69–74, 2006. Reprinted with permission. Copyright 2006, Nature Publishing Group.)

molecular orbital (LUMO) states of the neutral polymer are fully delocalized along the polymer chain and, in fact, exhibit significant dispersion with calculated bandwidths of several electron volts [88]. However, as a result of the strong electron–phonon coupling and the disorder-induced finite conjugation length, charges introduced onto the polymer interact strongly with certain molecular vibrations and are able to lower their energy with respect to the extended HOMO/LUMO states by forming localized polaron states surrounded by a region of molecular distortion [89]. There is clear, experimental evidence that the charge carriers carrying the current in a conjugated polymer FET are indeed of polaronic nature. Due to the surrounding molecular distortion and electronic relaxation, the charged molecule exhibits characteristic optical transitions below the absorption edge of the neutral molecule. These can be observed in operational FETs using charge modulation spectroscopy (CMS), which detects changes of the optical transmission of a semitransparent FET device upon gate voltage induced modulation of the carrier concentration in the accumulation layer [90]. In polymers such as poly(di-octyl-fluorene-co-bithiophene) (F8T2) in which close interchain interactions are weakened by the sp3-coordinated carbon atom on



117

au

(a)

au ag bg bu

LUMO

bg C2

π−π∗

C2 C3 au ag bg bu

au

C1

HOMO

bg

−1 −0.5

S

S

n C8H15 C8H15

0

π−π∗

C2 0.5

ΔT/T (104)

ΔT/T (104)

CT C1

−2 ∗

1

C3'

C6H13 ∗

−1 0

C3

C2

S

n∗

π−π∗

1 2

1

1.5

2 2.5 Energy (eV)

3

1

1.4

1.8

2.2

2.6

Energy (eV)

(c)

(b)

FIGURE 2.3.7 (a) Schematic energy diagram of neutral polymer (center), polaronic absorptions in the case of isolated chains (left) and interacting chains (right); charge modulation spectra of F8T2/PMMA (b) and P3HT/PMMA (c) top-gate FETs. (Courtesy of Shlomy Goffri, University of Cambridge.)

the fluorene unit, two characteristic sub-bandgap polaronic absorptions (Figure 2.3.7b) can be accounted for by the dipole-allowed C1 (≈0.4 eV) and visible C2 (1.6 eV) transitions of a simple isolated chain model (Figure 2.3.7a) [91]. In contrast, the charge-induced absorption spectrum of P3HT (Figure 2.3.7c) can only be explained by taking into account interchain interactions [92]. In addition to the C1 (0.3 eV) and C2 (1.3 eV) transitions, the CMS spectrum of high-mobility P3HT exhibits an additional C3 transition (1.6–1.8 eV), which is dipole forbidden in the isolated chain case, and low-energy charge transfer (CT) transitions at 60–120 meV [15,93]. Polarons in P3HT are not confined to a single chain, but are spread over several π-stacked chains. As a result of their two-dimensional nature, the polaron binding energy in P3HT is much reduced. From the position of the CT transition [89], the polaron binding energy Ep can be estimated to be on the order of Ep ≈

ECT ≈ 30–60 meV 2

At sufficiently high temperatures, charge transport of polaronic carriers in conjugated polymers should be governed by the physics of electron transfer processes, which was established by Marcus for chemical reactions and biological electron transfer processes [94]. In order for the localized polaron to hop between neighboring sites, the molecular configuration of the initial (occupied) site and the final (empty) site need to be distorted to a common configuration, where the molecular distortion


118


of both sites is equal (Figure 2.3.8a). This leads to thermally activated transport even in the absence of disorder. In the nonadiabatic limit, where the time scale for electron hopping is longer than that of the lattice vibrations, the mobility is given by: μ=

⎛ Ep ⎞ e ⋅ a2 ⋅ ν ⋅ exp ⎜− ⎟ ⎝ 2 ⋅ k ⋅T ⎠ k ⋅T

(2.3.1)

where ν is the attempt frequency ν=

π ⋅ J2 2 ⋅ E p ⋅ kT

and a is the typical hopping distance. However, in most experimental systems, the manifestations of the polaronic character of the charge carriers are masked by the effects of disorder. In any solutiondeposited thin film, disorder is present and causes the energy of a polaronic charge carrier on a particular site to vary across the polymer network. Variations of the local conformation of the polymer backbone, presence of chemical impurities or structural defects of the polymer backbone, or dipolar disorder due to random orientation of polar groups of the polymer semiconductor or the gate dielectric result in a significant broadening of the electronic density of states. The transport of charges injected into a molecular solid dominated by the effects of disorder is well understood from the work on molecularly doped polymers and other organic photoconductors used in xerography. Assuming a disorder-broadened Gaussian density of transport states with a characteristic width σ, Bässler [95] has shown on the basis of Monte Carlo simulations that an injected carrier hopping through such an otherwise empty density of states (DOS) relaxes to a dynamic equilibrium energy ε∞ = −

σ2 kT

below the center of the DOS, leading to a characteristic logμ ∝

1 T2

temperature dependence of the mobility (Figure 2.3.8b). The model has been improved by Novikov et al. [96], who showed that the dominant source of diagonal disorder is due to charge–dipole interactions and that spatial correlations of such interactions need to be taken into account in order to explain the commonly observed Poole–Frenkel dependence of the mobility on the electrical field. These researchers derived an expression for the electric field (E) and



119

3

EC

2

1

EF

Hopping at EF

Hopping in the band tail 2 EV

3 Extended state conduction (a)

5

ε/kT

ρ(ε/kT)

0

(ε∞)/kT

−5

−10

−1

0

1

3

2

4

5

lg (t/to) (b)

FIGURE 2.3.8 (a) Schematic energy diagram of DOS of a disordered semiconductor with a mobility edge. (b) Relaxation of energy distribution of an injected charge carrier hopping in a Gaussian DOS as a function of time. The DOS is shown as a dashed line on the right. (From Bässler, H., Phys. Status Solidi B, 175, 15, 1993. Reprinted with permission. Copyright 1993, Wiley.)


120


temperature dependence of the mobility in a correlated DOS with diagonal as well as nondiagonal positional disorder: 2 ⎡ ⎛ ⎛⎛ σ ⎞3/2 ⎞ e⋅a⋅ E ⎤ 3σ ⎞ ⎥ μ = μ 0 ⋅ exp ⎢− ⎜ ⎟ + 0.78 ⋅ ⎜⎜⎜ ⎟ − 2 ⎟⎟ ⎢⎣ ⎝ 5 ⋅ k B ⋅ T ⎠ ⎥⎦ σ k T ⋅ ⎝ ⎠ B ⎝ ⎠

(2.3.2)

The model describes the transport of individual injected carriers at zero/small carrier concentrations (i.e., in principle, it should not be directly applicable to the relatively high carrier concentrations p = 1018–1019 cm–3 present in the accumulation layer of FETs). Vissenberg [97] has developed a percolation model for variable range hopping transport in the accumulation layer of an FET assuming an exponential DOS with width T0. An expression for the field-effect mobility as a function of carrier concentration p was derived: ⎡ ⎛ T ⎞4 ⎛ T ⎞ ⎤ 0 ⎢ ⎜ 0 ⎟ sin ⎜ π ⎟ ⎥ T0 σ 0 ⎢⎝ T ⎠ −1 ⎝ T0 ⎠ ⎥ = pT 3 e ⎢ (2 ⋅ α) ⋅ Bc ⎥ ⎢ ⎥ ⎣ ⎦ T /T

μ FE

(2.3.3)

where σ0 is the prefactor for the conductivity, α is the effective overlap parameter between localized states, and Bc ≅ 2.8 is the critical number for onset of percolation. Transport in this model can be effectively described as activation from a gate voltagedependent Fermi energy to a specific transport energy in the DOS. Tanase et al. [98] have shown that in a series of isotropic, amorphous PPV polymers the large difference between the low mobility values extracted from spacecharge limited current measurements in LEDs and the comparatively higher fieldeffect mobilities can be explained by the largely different charge carrier concentrations (Figure 2.3.9). It was possible to fit the temperature dependence of the zerofield LED mobility to Equation 2.3.2 and the carrier concentration dependence of the FET mobility to Equation 2.3.3 with a consistent value of σ = 93–125 meV. Building on this work, Pasveer et al. showed that at room temperature the currentvoltage characteristics are dominated by the carrier concentration dependence of the mobility, while at low temperatures and high fields the field dependence of the mobility also needs to be considered [99]. The gate voltage dependence of the FET mobility of MEH-PPV has also been analyzed by Shaked et al. [100]. In several higher mobility amorphous hole transporting materials such as PTAA [18] and TFB [20], as well as in nematic, glassy polyfluorene-co-bithiophene [16], a somewhat different behavior was observed. The field-effect mobility was found to be independent of gate voltage within the carrier concentration range of 1018–1019 cm–3. In PTAA the low-density time-of-flight and high-density field-effect mobilities are of similar magnitude, with the bulk TOF mobility even higher by a factor of two to three at room temperature than the field-effect mobility. The Gaussian disorder



121

10−3 P3HT 10−4

T0 = 425 K

OC1C10-PPV

10−5 0.50 0.45

10−6

Ea (eV)

μh, μFE (cm2/Vs)

T0 = 540 K

10−7

1014

LED

0.40 FET

0.35 0.30 0.25

10−8

OC1C10-PPV

1015

1016

1017 p (cm

−20

−15

1018

−10 −5 Vg (V) 1019

0 1020

−3)

FIGURE 2.3.9 Hole mobility as a function of charge carrier concentration in diode and fieldeffect transistors for P3HT and a PPV derivative. (From Tanase, C. et al., Phys. Rev. Lett., 91, 216601, 2003. Reprinted with permission. Copyright 2003, American Physical Society.)

model was used to extract significantly smaller values of σ = 57 meV and σ = 68–90 meV from the temperature dependence of the time-of-flight and field-effect mobility of amorphous PTAA, respectively (Figure 2.3.3b). The increased σ-value in the case of the FET mobility was attributed to the contribution to energetic disorder from polar disorder in the dielectric close to the charge-transporting accumulation layer. The reason for the different behavior observed in PPVs with room-temperature field-effect mobility < 10–3–10–4 cm2/Vs and the higher mobility PTAA and polyfluorene polymers (μFET = 10–3–10–2 cm2/Vs) might be related to the lower degree of energetic disorder in the latter. With narrow DOS (σ < 60–90 meV), the expected concentration dependence of the room-temperature mobility over a concentration range of 1014–1019 cm–3 spanned by LED/FET measurements is relatively weak (i.e., less than an order of magnitude) and might be masked by other effects such as differences in bulk and interface microstructure, effects of interface roughness, or disorder effects induced by polar or charged groups in the dielectric. An alternative theoretical framework for understanding the effects of disorder is the multiple trapping model, which is well established for describing transport in amorphous silicon and has been claimed to be more appropriate for describing the charge transport in microcrystalline polymers such as P3HT [22] and poly(bis(alkylthienyl-bithiophene) [101,102]. This model assumes that disorder broadening is sufficiently weak that, in a certain energy range, the DOS becomes high enough that electronic states above the so-called mobility edge are extended, while electronic states below the mobility edge remain localized (Figure 2.3.8a). The current is


122


assumed to be transported by carriers thermally activated into the delocalized states above the mobility edge, while carriers in localized states are effectively trapped and do not contribute to the current. Assuming a specific DOS and a mobility for carriers above the mobility edge, the FET current can be obtained by first determining the position of the quasi-Fermi level at the interface for a particular gate voltage and then calculating the number of free carriers thermally excited above the mobility edge using Fermi–Dirac statistics. Salleo et al. [101] found that the multiple trapping model explained the temperature dependence of the FET mobility of poly(bis(alkyl-thienyl-bithiophene) more consistently than the Vissenberg hopping model; the latter yielded an unphysical dependence of σ0 and T0 on the processing conditions. In spite of detailed investigations to model the charge transport in a mobility regime between 10–2 and 1 cm2/Vs, it can be difficult to distinguish between hopping and band transport models. Many of the qualitative features are common to both hopping and multiple trapping and release models, such as the mobility decreasing with decreasing temperature and the dependence of the mobility on the carrier concentration. Therefore, characterization of the charge transport by techniques that provide complementary information is needed. One of the techniques providing such information is CMS. The spectroscopic properties of polarons in P3HT have been characterized as a function of molecular weight and film deposition conditions by CMS [32]. CMS experiments on regioregular P3HT have revealed a pronounced low-energy charge transfer (CT) transition in the midinfrared spectral region [15]. This transition can be interpreted in the framework of Marcus–Hush electron transfer theory describing the transfer of electrons between neighboring molecules in the presence of strong electron–lattice interactions [103]. The process is governed by two main parameters: The relaxation energy λ (which is twice the polaron binding energy) measures the energy lowering that charged molecules can achieve by adopting a relaxed conformation as a result of the electron–lattice coupling. The transfer integral t is a measure of the strength of the interchain coupling of the electronic wave functions on neighboring molecules. In the weak coupling case (λ > 2t), the lower adiabatic potential surface has a number of minima, and the charge is localized on an individual molecule (Figure 2.3.10a). Under such conditions, a charge transfer optical transition is observed centered at an energy ωCT = λ. In contrast, in the strong coupling case (λ < 2t), the lower adiabatic potential surface has only one minimum and the charge is delocalized over a certain number of neighboring molecules. In this case, an optical charge transfer transition can also be observed; it is not centered around λ, but rather around ωCT = 2t (Figure 2.3.10b). In intermediate MW samples with mobilities > 0.05 cm2/Vs, we observe an intense CT transition centered around 0.1 eV (Figure 2.3.10c). In the highest mobility, highest MW samples, the transition is similarly intense and appears to peak at slightly lower energies below 0.08 eV, which is the low energy cutoff of our experimental setup. In contrast, in the low MW samples with mobilities less than 10–2 cm2/Vs, a much less intense CT transition is observed, and the transition peaks at significantly higher energies on the order of 0.3 eV.


Energy

123

Energy


λ

Δ = 2Hab

Q

Q

(a)

(b)

12

−ΔT/T (10−4)

10 8 76 kD

6 4

29 kD

2 15.4 kD 0

0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

Energy (eV) (c)

FIGURE 2.3.10 Potential energy diagram as a function of configuration coordinate illustrating electron transfer between two sites in the case of weak coupling (a) and strong coupling (b). (c) Charge modulation spectra in the midinfrared spectral range of TCB spin-cast P3HT films for different MW. The spectra were obtained by subtracting the infrared absorption spectra of the device structure taken at 10 and –30 V.

The position and intensity of the CT transition appears to be very directly correlated with the field-effect mobility. In high-mobility P3HT, the strong coupling situation applies [103]. The lower intensity of the CT transition indicates a lower degree of interchain polaron delocalization in the low MW samples. A plausible explanation for the reduced intensity and higher energy of the CT transition in the low MW samples is that, due to the enhanced disorder and shorter conjugation length in these samples, the weak coupling regime might apply. Such a crossover behavior between localized and delocalized polarons as a function of MW would provide an intriguing microscopic explanation for the observed rapid increase of mobility with MW below 15–20 kD (see Figure 2.3.4b) [32].


124


5.2

ΦELP/SUB (eV)

5.0 4.8

PFO P10AF TFB P3HT

4.6 4.4 4.2

{ , }

4.0 3.8 4.0

4.5

5.0

5.5

6.0

ΦSUB (eV)

FIGURE 2.3.11 Dependence of work function of polymer coated substrate, ΦELP/SUB, on the work function of bare substrate, ΦSUB, for four studied materials: P3HT, TFB, poly(9-1decylundecylidene fluorene (P10AF), and polydioctylfluorene (PFO). (From Tengstedt, C. et al., Appl. Phys. Lett., 88, 053502, 2006. Reprinted with permission. Copyright 2006, American Institute of Physics.)

2.3.6 CHARGE INJECTION PHYSICS Another important aspect of the device physics, particularly in the context of short channel OFETs with L < 5 μm, is the injection of charges from a metal source-drain contact into the organic semiconductor. In contrast to inorganic semiconductors, controlled doping of organic semiconductors is still difficult, since dopants incorporated in the form of small molecule counter ions can migrate and cause device instabilities. Since most organic semiconductors that have shown useful FET performance have band gaps > 2 eV, the formation of low-resistance ohmic contacts with common metals is often challenging. The energy barrier for hole injection at the metal–poymer interface is determined by the vacuum work function of the metal contact ΦW and the ionization potential IP of the polymer. For conjugated polymer films spin-coated onto hole injecting metal electrodes, it has been reported that as long as ΦW is smaller than a critical value characteristic of the polymer, no interface dipole is formed [104]. In this case, the barrier for hole injection can be estimated simply by aligning the vacuum levels of the metal and the polymer (Mott–Schottky limit); the measured work function of the metal with the polymer deposited on top increases linearly with ΦW with a slope of one (see Figure 2.3.11). However, when ΦW exceeds said critical value a significant interface dipole can be formed. Positive charges are transferred from the metal to the semiconductor and the position of the Fermi level at the interface becomes pinned at an energy level interpreted as the hole polaron/bipolaron energy level in the polymer semiconductor. This simple picture suggests that, at least in the case of solution-deposited polymers on common hole-injecting contacts, chemical interactions between the metal and



125

the polymer and other defect states in the band gap of the polymer do not influence strongly the contact formation. We emphasize that this is less likely to be the case for metals deposited on top of the polymer semiconductor as well as for reactive metals employed to achieve electron injection into common organic semiconductors. It has been shown that deposition of gold contacts on top of an organic semiconductor, such as pentacene, can result in formation of trap states in the organic semiconductor [105]. There are intriguing reports of efficient charge injection in systems for which Schottky barriers calculated using Mott–Schottky theory should exceed 1 eV, such as hole injection from Ca into P3HT [56] or electron injection from Au into fluorocarbon-substituted oligothiophenes [106]. It is likely that in such systems chemical interactions and interface states are important factors that determine contact formation. In order to understand the contact injection in the OFET not only the interface electronic structure, but also the device configuration and injection geometry need to be taken into account because they determine the potential profile in the vicinity of the contact and the transport of injected charges away from the contacts. In the bottom gate device configuration, the charge injection physics can be studied directly using scanning Kelvin probe microscopy (SKPM) [107–109]. SKPM uses an atomic force microscope tip with a conducting coating operated in noncontact mode to probe the electrostatic potential profile along the channel of the OFET with a spatial resolution on the order of 100 nm (Figure 2.3.12a). The voltage applied to the conducting tip is regulated by a feedback loop so that the electrostatic force between tip and sample is minimized. For polymer TFTs in accumulation the tip potential essentially follows the electrostatic potential in the accumulation layer. Figure 2.3.12(b) shows typical SKPM potential profiles obtained for bottomgate, bottom contact P3HT devices on SiO2 with Au contacts comprising a Cr adhesion layer (Cr–Au) (inset) and pure Cr contacts. It can be seen that the contact resistance at the source and drain contacts exhibits very different behavior in the two cases. In the case of Cr–Au contacts, generally, in the case of contacts for which the Schottky barrier Φb is less than 0.3 eV, the voltage drop across the source and drain contacts is small and the contact resistance at the source contact is very similar to that of the drain contact. This is somewhat unexpected since in normal FET operation the source contact is reverse biased while the drain contact is forward biased, implying that the source contact resistance should be significantly larger than the drain contact resistance. This implies that under conditions that might be typical for high-performance OFETs, the contact resistance is not determined by the Schottky barrier at the interface, but by bulk transport processes in the semiconductor in the vicinity of the contact. Consistent with this interpretation, the contact resistance was found to depend on temperature in the same way as the mobility [107] so that the potential profiles become independent of temperature. This result was explained by invoking the existence of a depletion layer in the vicinity of the contacts. Similar results have been reported using channel length scaling analysis [76]. In contrast in the case of Cr contacts or generally for systems with Schottky barriers exceeding 0.3 eV, the voltage drops across the contacts become very significant and the source resistance is found to be larger than the drain resistance, as


126


Tip-height control

PSD

+

Channel

Tip Organic layer Source Insulator: SiO2

Drain Id

Gate: n+-Si Vg

Vd

Accumulation layer (a)

Drain

Source 0 0 −2 −4 −6 −8

Local potential (V)

−1 −2 −3

Source

Drain

ΔVs

145 K

−4

T = 300 k Cr-Au

−1 −2 −3

Cr 6

4

2

0

−2

190 K

−5

−4 −5

−6

−6 300 K

−7

−7

ΔVd

−8 −2

0

−1

0

1

2

3

4

5

−8 6

7

Distance from source (μm)

(b)

FIGURE 2.3.12 (a) Schematic diagram of experimental setup for scanning Kelvin probe microscopy (SKPM). (b) Profiles of an L = 5.5 μm P3HT transistor with Cr electrodes taken at three different temperatures (Vg = –40 V, Vd = –8 V). The inset shows a profile obtained after switching the source and drain contacts on the same TFT with both Cr and Cr–Au contacts Vg = –40 V, Vd = –8 V). (From Burgi, L. et al., J. Appl. Phys., 94, 6129–6137, 2003. Reprinted with permission. Copyright 2003, American Institute of Physics.)



127

expected. This implies that in this regime the contact resistance is determined by the injection physics at the interface [107]. It is remarkable that, in spite of the significant expected Schottky barrier height, the contact resistance shows only a very weak increase with decreasing temperature, which is even weaker than that of the field-effect mobility (i.e., the voltage drop at the contacts decreases compared to that across the bulk of the polymer with decreasing temperature and charge transport becomes less contact limited at low temperatures). This behavior cannot be explained in the framework of the commonly used diffusion-limited thermionic emission model [110], which takes into account backscattering into the metal due to the small mean-free path in the organic semiconductor and predicts the activation energy of the contact resistance to be larger than that of the mobility and larger than Φb/kT. Explanation of the experimental data required taking into account disorderinduced broadening of the density of states of the organic semiconductor, which provides carriers with injection pathways through deep states in the DOS, leading to a reduced effective barrier at low temperatures. Similar conclusions have recently been drawn on the basis of channel length scaling experiments [111].

2.3.7 DEFECT STATES AND DEVICE DEGRADATION MECHANISMS Electronic defect states in the semiconductor at the interface between semiconductor and dielectric and inside the dielectric layer can cause instabilities of the threshold voltage of the TFT. For practical applications, the threshold voltage stability is a figure of merit as important as, if not more important than, the field-effect mobility because it is closely related to the operational and shelf lifetime of the device. Most TFT technologies, including those based on a-Si, suffer from threshold voltage shifts induced by bias-temperature stress (BTS). In a-Si TFTs, generation of defect states inside the semiconducting layer, such as dangling bond defects, as well as charge injection into the SiNx gate dielectric has been found to contribute to VT shifts upon BTS; charge injection into the dielectric is the dominant mechanism in high-quality material [112]. Several groups have recently reported systematic BTS investigations and studies of organic TFT characteristics upon exposure to atmospheric conditions and humidity. In most p-type organic semiconductors, a negative shift of the threshold voltage is observed upon prolonged operation of the device in accumulation, which is generally attributed to charge trapping in the organic semiconductor and/or at the active interface. Matters et al. reported negative VT shifts due to charge trapping for a PTV precursor polymer in contact with inorganic SiO2 dielectric; these were more pronounced in the presence of water than when the device was operated in vacuum or dry air [113]. Street et al. reported significant negative VT shifts in F8T2/SiO2 bottom-gate, bottom-contact TFTs [114], which were more pronounced than reported for top-gate F8T2 devices with a polymer dielectric [16]. Street et al. also found the VT stability of PQT/SiO2 devices to be significantly better than that of F8T2/SiO2


128


devices. It is clear from these experiments whether the device configuration, choice of contacts, and dielectric play a crucial role in determining the device stability. There is little known at present about the nature of the electronic states involved in defect formation and device degradation. Few experimental studies have been aimed at understanding at a microscopic level the nature of defect states in organic semiconductors [115]. Device modeling has been used to understand the subthreshold characteristics of OFETs [116]. Based on an analysis of the relationship between the current decay at early times after FET turn-on and the hole concentration in the channel, Street et al. have suggested that charge trapping occurs due to formation of low-mobility bipolarons by reaction of two polarons [114,117]. However, Deng et al. performed optical spectroscopy of field-induced charge on F8T2/PMMA TFTs exhibiting significant VT shifts, but were unable to detect the spectroscopic signature of bipolarons [91]. When discussing threshold voltage shifts in OFETs, it is important to distinguish between reversible and irreversible charge trapping effects [118]. Reversible charge trapping depends on the duty cycle during operation and can be recovered by not operating the device for several minutes or hours. Threshold voltage shifts due to irreversible charge trapping are independent of duty cycle and do not recover on timescales of hours if the device is not operated while being kept in the dark. However, the irreversible threshold shift can be erased by illuminating the sample with above-bandgap light [119,120] (Figure 2.3.13a). The spectral dependence of the light-induced recovery follows the absorption spectrum of the organic semiconductor (Figure 2.3.13b). Charge traps that can be emptied in this way must be located inside the organic semiconductor or directly at the interface, but cannot be located inside the gate dielectric. It has also been reported that a positive gate voltage stress leads to a shift of VT to more positive values [119]. This has recently been explained by injection and trapping of negative electrons at the interface [56]. Zilker et al. have reported that, in films of p-type solution-processed pentacene in contact with an organic photoresist dielectric, the threshold voltage shifts to more positive values for negative gate bias stress during operation in air [121]. The VT shift was the more pronounced the smaller the source-drain voltage was. This was interpreted as the result of mobile ions drifting in the gate dielectric in the presence of water. Negative ions drifting towards the active interface cause accumulation of positive countercharges in the semiconducting layer. Only during operation in vacuum or in dry air was a negative VT shift of –3 V after application of Vg = –20 V for 1000 s observed resulting from charge trapping at or near the interface. Rep et al. have investigated the role of ionic impurities originating from the substrate on the conductivity of P3HT films [122]. On Na2O containing glass substrates, Na+ ions were found to drift towards the negatively biased contact, leaving behind negative charge centers on the glass surface. Gomes et al. have claimed recently that the bias stress instability in organic FETs is caused by trapped water in the organic semiconductor [123]. The preceding results point to the crucial role of the gate dielectric in determining the operational and shelf stability of the device. Several groups have recently reported very encouraging BTS and shelf lifetime data for solution-processed



129

8 × 10−8

1st meas. stressed 150 s. 350 s. 900 s.

Drain current (A)

7 × 10−8 6 × 10−8 5 × 10−8 4 × 10−8 3 × 10−8 2 × 10−8 1 × 10−8 0 −25

−20

−15

−10

−5

0

Gate voltage (V) (a) 0.08 hv

0.6

0.06

0.4

0.04

0.2

0.02

0.0

2

3

4

5

6

7

τ1/2−1 (s−1)

Absorption (arb. units)

0.8

0.00

E (eV) (b)

FIGURE 2.3.13 (a) Pulsed transfer characteristics of bottom-gate F8T2/SiO2 FET after applying a negative gate bias stress showing subsequent recovery of the threshold voltage shift after illuminating the device for different periods of time. (From Salleo, A. and Street, R.A., J. Appl. Phys., 94, 471–479, 2003. Reprinted with permission. Copyright 2003, American Physical Society.) (b) Comparison of the wavelength dependence of the time constant for the light-induced trap release in TFB/SiO2 with the absorption spectrum of the organic semiconductor. (From Burgi, L. et al., Syn. Met., 146, 297–309, 2004. Reprinted with permission. Copyright 2004, Elsevier.)

OFETs measured and stored in air without special encapsulation. PTAA combined with low-k dielectrics exhibits excellent shelf life with no discernible VT shift upon storage in air and light for periods of several months [18]. Similarly, TFTs based on TFB with BCB dielectric exhibit very good operational stability during accelerated lifetime testing at temperatures of 120°C [20]. In both cases, the good stability is believed to be related to the use of an apolar, low-k dielectric, which is less susceptible to ionic impurities, and the amorphous microstructure of the aryl-aminebased polymer semiconductor with good thermal and photostability and low degree of energetic disorder.


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The group at Plastic Logic has recently reported excellent operational stability results on unencapsulated polymer TFTs fabricated on PET substrates [10]. Although, of course, significant work to assess and improve the operational and shelf life of OFETs under realistic application conditions and to understand degradation mechanisms in much more detail remains, these early results strongly suggest that solution-processed OFETs can exhibit device stability and reliability similar to if not higher than their a-Si counterparts.

2.3.8 OUTLOOK Solution-processable organic FETs have become a promising emerging technology for low-cost, large-area electronics on flexible plastic substrates. FET performance is approaching that of a-Si TFTs, and solution/printing-based manufacturing processes have been developed. Device operational and environmental stability has improved significantly recently as a result of availability of organic semiconductors with higher inherent oxidative stability, better understanding of the requirements for gate dielectrics, and more controlled manufacturing processes. In this chapter, we have reviewed recent progress in understanding the device physics of solution-processable organic semiconductors. It should be apparent from the discussion that although much progress has been made in understanding the materials physics and requirements for high-performance FETs, understanding of the fundamental excitations and processes at a microscopic level involved in charge transport and injection as well as device degradation is still much more superficial than the corresponding level of fundamental understanding available in inorganic semiconductors. Particularly, many fundamental aspects of the correlation between the structure and physics of charge transport at solution-processed organic–organic heterointerfaces remain to be explored. However, the field of organic electronics is gaining momentum. Continued breakthroughs in materials and device performance, concrete industrial applications in active matrix flexible electronic paper displays, and simple, low-cost intelligent labels are emerging on the horizon to be commercialized within the next three to five years. It is likely that new scientific discoveries and technological advances will continue to cross-fertilize each other for the foreseeable future.

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2.4

Contact Effects in Organic Field-Effect Transistors

Matthew J. Panzer and C. Daniel Frisbie CONTENTS 2.4.1 Introduction................................................................................................139 2.4.2 Definition of an Ohmic Contact ................................................................140 2.4.3 Origins of Contact Resistance ...................................................................140 2.4.3.1 Electronic Structure and Potential Barriers at Metal–Organic Interfaces .....................................................................................140 2.4.3.2 Charge Transport across Metal–Organic Interfaces ...................142 2.4.3.3 Influence of Channel Dimensions...............................................145 2.4.3.4 Influence of Device Architecture ................................................146 2.4.4 Measuring Contact Resistance...................................................................148 2.4.4.1 Extrapolation of Device Resistance to Zero Channel Length....148 2.4.4.2 Gated Four-Probe Measurements................................................149 2.4.4.3 Kelvin Probe Force Microscopy .................................................150 2.4.4.4 Measured Contact Resistance Values..........................................151 2.4.5 Contact Engineering ..................................................................................154 2.4.5.1 Chemical Modifications ..............................................................154 2.4.5.2 Ambipolar and Light-Emitting OFETs.......................................155 2.4.5.3 Channel Dimensions: How Small? .............................................155 References..............................................................................................................155

2.4.1 INTRODUCTION In an ideal organic field-effect transistor (OFET), the source and drain contacts are ohmic, meaning the value of the contact resistance is negligibly small in comparison with the electrical resistance of the semiconductor (i.e., the channel resistance). While this situation can be achieved in real devices, there are several practical considerations for fabricating OFETs that are not contact limited [1]. In this chapter, we begin by defining an ohmic contact and continue with a discussion of the origins of contact resistance. Subsequent sections cover techniques used to quantify contact resistance in working OFETs, along with tables of contact resistance values for 139


140


L Source Rsource

Rchannel

Drain Rdrain

Insulator Gate

FIGURE 2.4.1 Schematic of an OFET showing equivalent resistances corresponding to the source contact resistance, the channel resistance, and the drain contact resistance. The total contact resistance is Rc = Rsource + Rdrain and the total device resistance is Rtot = Rc + Rchannel.

common OFET geometries based on typical organic semiconductors. We conclude with comments on contact engineering to improve OFET performance.

2.4.2 DEFINITION OF AN OHMIC CONTACT In traversing an OFET channel from source to drain, charge carriers are (1) injected from the source contact into the semiconductor channel; (2) transported across the length of the channel; and (3) extracted from the channel into the drain. These processes can be roughly thought of as three resistors in series (Figure 2.4.1). The resistances associated with carrier injection and collection steps can be grouped into the contact resistance (Rc), while the resistance associated with crossing the channel length in the semiconductor is termed the channel resistance (Rchannel). Keeping the contact resistance small compared to the channel resistance is crucial to the realization of “ohmic contacts” in OFETs (i.e., for an ohmic contact, Rc 0.3 eV) exist at many metal–organic semiconductor interfaces, it is possible to make ohmic source and



143

∆dip = 0.60 eV

Evac

Wf = 5.05 eV

Bulk

Interface

Evac

Φe = 1.17 eV EF

Eg = 2.68 eV

Au

Bulk

Interface

Φh = 0.47 eV

Pentacene

FIGURE 2.4.3 Band line-up diagram for the Au-pentacene interface. An interface dipole results in a vacuum level shift of 0.6 eV. The diagram indicates that the valence band (HOMO level) and conduction band (LUMO level) have finite widths. The hole injection barrier (Φh) is determined by taking the difference in energy between the Fermi level and the edge of the valence band. The slight shrinking of the band gap at the interface is due to polarization effects by the metal. (Reprinted from F. Amy, C. Chan, and A. Kahn, Polarization at the gold/pentacene interface, Organic Electronics, 2005, 6, 85–91, with permission from Elsevier.)

drain contacts in OFETs. A likely reason for this is that the charge injection mechanism is probably not simple thermionic emission in which carriers must surmount the full potential barrier, as originally indicated in Figure 2.4.2(b). Instead, at large interfacial electric field strength, field emission (tunneling) through the barrier can become possible; this is a process that effectively lowers the potential barrier. Another possible injection mechanism involves defect-assisted transport in which carriers bypass the barrier by hopping through midgap states. Figure 2.4.4 shows a simple comparison of these different charge injection mechanisms. There is mounting experimental evidence that the charge injection process at the source electrode in OFETs is not simple thermionic emission [10–13]. First, measurements of the source contact resistance as a function of temperature reveal that the injection process is indeed thermally activated (which is consistent with thermionic emission), but the activation energies are generally much smaller than the estimated potential barriers determined by photoemission spectroscopy. In some cases, the activation energy associated with the source contact resistance is very comparable to the activation energy associated with the carrier field effect mobility, suggesting that transport of charge in the semiconductor near the contact is the limiting bottleneck, not the actual metal-to-semiconductor emission process. Second, channel potential measurements by Kelvin probe force microscopy and the four-probe method, which are described later in this chapter, indicate that the contact resistances and temperature dependences associated with the individual source and drain electrodes are nearly identical. From a thermionic emission


144


Evac

Evac

Evac

EF +

EGAP EF-EV +

VB (HOMO)

V (a) Thermionic emission

CB (LUMO)

CB (LUMO)

CB (LUMO) EF

EGAP + +

+

V (b) Field emission (tunneling)

VB (HOMO)

EGAP

EF ++ +

VB (HOMO)

V (c) Defect-assisted injection

FIGURE 2.4.4 Comparison of different charge injection mechanisms at a biased metalsemiconductor contact: (a) thermionic emission, (b) field emission (tunneling), (c) defect assisted injection.

viewpoint, one would expect the resistance at the source electrode (the injecting contact) to be much larger than the resistance at the drain (the collecting contact). The fact that the source and drain contact resistance behaviors are very similar in most devices also indicates that the bottleneck at the contacts is related to charge transport in the (disordered) organic semiconductor near the contacts and not simply due to an injection barrier at the metal–organic interface. An additional point is that, in general, the source and drain contact resistances are gate voltage dependent; specifically, they decrease with increasing gate voltage. The variation of the contact resistance with gate voltage is essentially identical for both the source and drain and it is also similar to the variation of the channel resistance. The close tracking of the gate voltage dependence on the source, drain, and channel resistances also indicates that contact resistance depends on film transport properties near the contact. The simplest picture that can explain these collective observations is that the source contact resistance is the sum of (1) resistance arising from charge injection over or through the potential barrier; and (2) resistance due to transport in the disordered depletion region near the contact. The drain resistance would simply be due to the latter part. Perhaps because the presence of a strong gate field facilitates field emission through the barrier, resistance (1) is often not limiting for the source; instead, resistance (2) dominates, and thus the source and drain resistances are comparable. This description provides a useful physical picture, but quantitative modeling of the injection and collection processes is complicated by details of the device geometry and the fabrication process, such as the evaporation of hot metal onto the organic semiconductor, which might produce many defects. Work on detailed understanding of transport at OFET contacts is ongoing.



145

Insulator Semiconductor

W Source

Drain

L

FIGURE 2.4.5 Top view of an OFET showing the relevant channel dimensions.

2.4.3.3 INFLUENCE

OF

CHANNEL DIMENSIONS

As mentioned earlier, channel dimensions are also critical in determining the importance of contact effects in OFETs. Figure 2.4.5 shows the top view of a typical OFET with channel length (L) and channel width (W). It is useful for comparison purposes to report the normalized or specific contact resistance (Rc′) as a “raw” resistance value multiplied by the width (W) of the channel. Thus, the units of Rc′ are Ωcm. Physically, the reciprocal of this value is the contact conductance per unit width of contact. The total device resistance can be expressed as RTOT = Rchannel + Rc

(2.4.3)

sheet In terms of the channel sheet resistance ( Rchannel , units of Ω/square, and VG dependent) and the specific contact resistance (Rc′, units of Ω cm and also VG dependent), this equation becomes:

sheet RTOT = Rchannel

L Rc′ + W W

(2.4.4)

This equation facilitates understanding of how the channel dimensions (L and W) affect the relative magnitudes of the contact resistance and the channel resistance. Note that the channel resistance scales as L/W but the contact resistance scales as 1/W; it does not depend on L. Consider two different OFET devices on the same semiconductor/insulator/gate substrate: both have the same channel width (equal W), but the length of the channel of the second device is 10 times smaller than that of the first (L2 = L1/10), as depicted in Figure 2.4.6(a). Both devices have equal contact resistances Rc because W is the same. But because the channel resistance scales with L/W (the source-drain current scales with W/L), the channel resistance of the second device is 10 times smaller than that of the first device. This means that contact resistance is potentially much more important in the shorter channel device because it contributes a larger fraction of the total resistance.


146


Device 1

W1 = W2 L1 = 10 × L2

Device 2

Equal contact resistances; Device 2 has a lower channel resistance (by 10X) (a)

Device 3

W3 = 0.5 × W4 L3 = 0.5 × L4 (W/L)3 = (W/L)4

Device 4

Equal channel resistances; Device 4 has a lower contact resistance (by 2X) (b)

FIGURE 2.4.6 (a) Decreasing the channel length (L) at constant channel width (W) leads to lower channel resistance and increased relative contact resistance, Rc/RTOT. (b) Increasing the channel size at constant (W/L) decreases the relative contact resistance.

A second hypothetical pair of OFET devices is shown in Figure 2.4.6(b). In this case, both contact patterns share the same W/L ratio, but both W and L are twice as large for the second device of this duo. While the channel resistances are now equivalent, the contact resistance of the larger device is half that of the smaller one. The larger device in this case is less likely to be contact limited. In general, one must pay attention to the magnitude of the contact resistance when scaling OFETs to very small lengths (L) because the contact resistance of short-channel devices can quickly become the dominating resistance. As a final note, to avoid nonidealities in transistor output curves, W/L should always be ≥10 in order to avoid fringing field effects at the edges of the channel.

2.4.3.4 INFLUENCE

OF

DEVICE ARCHITECTURE

There are two main architectures to choose from in OFET fabrication: the top contact and the bottom contact configurations. The physical difference between the two is the order of fabrication steps. That is, the source/drain contacts are either deposited before or after the semiconductor layer is deposited to create a bottom contact or top contact device, respectively. One can also build the entire transistor on top of the semiconductor layer (the so-called top gate architectures), in which the insulator



Top contacts Source

147

Bottom contacts Source

Drain

Drain

Semiconductor Insulator Gate (a)

(b) Bottom contacts/top gate

Top contacts/top gate Gate Source

Drain

Insulator

Source

Drain

Semiconductor Substrate (c)

(d)

FIGURE 2.4.7 Four possible OFET architectures (in cross-section), including (a) top contacts, (b) bottom contacts, (c) top contacts with a top gate, and (d) bottom contacts with a top gate.

and gate contact are sequentially deposited on top of either of the two contact configurations. All four of these OFET architectures are shown schematically in Figure 2.4.7. Top contact OFETs (Figure 2.4.7a) generally exhibit the lowest contact resistances. This is likely because of the increased metal–semiconductor contact area in this configuration. A major contribution to contact resistance in the top contact configuration is access resistance (see Figure 2.4.8a). Access resistance results from the requirement that charge carriers must travel from the source contact on top of the film down to the accumulation layer (the channel) at the semiconductor–insulator interface and then back up to the drain contact to be extracted. In order to minimize access resistance, the thickness of the organic semiconductor layer should not be too large. However, some researchers have proposed that access resistance is less than might be expected for top contact OFETs because the contact metal penetrates the film down to the accumulation layer (perhaps due to large peak-to-valley roughness of the semiconductor film or the nature of the metal deposition process) [11]. This scenario is shown in Figure 2.4.8(b). With the bottom contact architecture (Figure 2.4.7b), access resistance is not an issue because the contacts are in the same plane as the OFET channel. In addition, very small channel dimensions (W, L < 10 μm) can be prepatterned on the insulator using conventional photolithography. A limitation to the bottom contact configuration, however, is that film morphology in the vicinity of the contacts is often nonideal. A number of researchers have demonstrated that the organic semiconductor grain sizes are very small near the contacts, presumably due to heterogeneous nucleation phenomena [14]. Pentacene molecules, for example, prefer to “stand up” with the long axis of the molecule perpendicular to the plane of the substrate when deposited on the commonly used insulator SiO2 [15]. When deposited on top of gold contacts,


148


Polycrystalline semiconductor

Access resistance Source

Drain

Source

Drain

Insulator Insulator Gate (a)

(b)

FIGURE 2.4.8 (a) Access resistance in a top contact OFET represented by arrows indicating the injection and extraction of charge through the bulk semiconductor film between the source/drain contacts and the conductive channel. The plus signs represent the hole accumulation layer of a p-channel device. (b) Detailed view of the contact/semiconductor interface near the accumulation layer of a top contact OFET with a very rough, polycrystalline semiconductor thin film. Ovals highlight the penetration of the top contacts deep into the film; in such devices, access resistance can often be negligible.

however, strong interactions between the pentacene π-clouds and the metal surface lead to tiny grains at the contact and, in some cases, voids are observed [14]. Semiconductor growth at the complex triple interface (contact-semiconductor-insulator) is not very well understood, although it is clear that the bottom contact configuration almost always creates greater contact resistance than in the case of top contacts. Of the two top-gate OFET architectures (Figure 2.4.7c and 2.4.7d), the top contact/top gate configuration (Figure 2.4.7c) is the more favorable of the two because bottom contact/top gate devices suffer from access resistance. However, it should be noted that both top gate architectures face the additional concerns of semiconductor top surface roughness (since this is where the channel will form) and forming a stable interface between the insulator and the top of the semiconductor film. Solution deposition of the top insulator material, for example, may damage the underlying semiconductor film. Finally, regarding the alignment of the top gate contact to the OFET channel in top gate devices, care must also be taken to ensure that the gate reaches completely across the entire length (L) of the device. If the length of the gate electrode is less than the channel length or if the gate is simply misaligned, additional contact resistance will be introduced as a result of ungated semiconductor regions at one or both of the contacts.

2.4.4 MEASURING CONTACT RESISTANCE 2.4.4.1 EXTRAPOLATION OF DEVICE RESISTANCE CHANNEL LENGTH

TO

ZERO

Since the total resistance a charge carrier experiences during its journey from source to drain (i.e., VD/ID) is the sum of the contact resistance and the channel resistance



149

Device resistance (kΩ⋅cm)

200

150

100

Contact resistance

50

0

0

20

40 60 80 Channel length, L (μm)

100

FIGURE 2.4.9 Example of a transmission line (RTOT · W vs. L) plot at a given VG value. Extrapolation of the data to a channel length of zero yields the specific contact resistance Rc′ as the y-intercept.

as discussed earlier, one of the simplest ways to quantitatively measure OFET contact resistance involves isolating these two resistances by extrapolation. By fabricating several pairs of source and drain contacts with different channel lengths (but constant W) on the same semiconductor film, one measures the total OFET resistance and makes a plot of resistance versus channel length. Linear extrapolation of the plot to L = 0 effectively eliminates the channel resistance and yields the contact resistance as the y-intercept. This method of measuring contact resistance is known as the transmission line method or R vs. L technique. Figure 2.4.9 shows an example of a resistance (RTOT·W) versus channel length plot with extrapolation used to determine the specific contact resistance, Rc′. Because both the channel and contact resistances are affected by the charge density in the channel, the resistance measurements are made at a single gate voltage (usually in the linear operation regime). The drain voltage could also be scaled to L in order to obtain the same source-to-drain lateral electric field for each device. While this technique is straightforward to perform and understand, it has two disadvantages. First, it can be tedious to fabricate and test several devices in order to obtain the resistance versus channel length plot. In addition, the contact resistance obtained by the transmission line technique lumps the individual source and drain resistances together.

2.4.4.2 GATED FOUR-PROBE MEASUREMENTS In order to separate the individual contributions of the source and drain contacts to the total contact resistance, a more sophisticated measurement technique is required. The gated four-probe technique utilizes two narrow, voltage-sensing electrodes situated between the source and drain electrodes and slightly protruding into the channel, as shown in Figure 2.4.10(a). During the course of normal OFET electrical characterization, these voltage-sensing probes are connected to high-input impedance electrometers that sense the channel potential at the two probe positions (V1,


150


15

Drain

Potential (V)

Source

10

x

V2

V1 5

ΔVS 0 0 L 2L L 3 3 (a)

ΔVD

VD = 15 V VG = 75 V

V2

V1

0

L/3 2L/3 Channel position, x

L

(b)

FIGURE 2.4.10 (a) Top view of an OFET contact pattern with two voltage sensing probes (V1, V2) penetrating slightly into the channel in order to perform gated 4-probe measurements. (b) Hypothetical extrapolated linear channel potential profile based on the voltage-sensing probe data, which reveals the voltage drops (ΔVS, ΔVD) at the source and drain contacts that are responsible for the contact resistance. The dashed line is the ideal (no contact resistance) linear potential profile for an applied drain voltage of 15 V and a gate voltage of 75 V.

V2) without passing any current. In the linear regime of OFET operation (VG >> VD), the channel should be uniform in charge carrier density with a linear drop in electrostatic potential along L from source to drain. Therefore, any drops in electrostatic potential that occur at the contacts (due to contact resistance) will be manifested upon extrapolation of the channel potential profile based on the voltage-sensing probe measurements. As depicted in Figure 2.4.10(b), contact resistance at the source and drain electrodes results in a smaller than expected slope of the potential versus channel position profile. The profile is estimated by linear extrapolation between V1 and V2. Individual source and drain contact resistances are calculated by dividing the voltage drops ΔVS and ΔVD by the source-drain current, respectively. By isolating the source and drain contact contributions to the total contact resistance, the gated four-probe technique provides more information than the transmission line technique, and it is possible to determine Rc in one device (vs. several). An important caveat for the gated four-probe technique is that the extrapolated channel potential profile will only be valid for strict linear regime OFET operation (VG >> VD), where the channel potential profile can be expected to be linear and uniform.

2.4.4.3 KELVIN PROBE FORCE MICROSCOPY While the previously described techniques both require extrapolation of measured data in order to calculate the contact resistance, Kelvin probe force microscopy (KFM, also known as scanning surface potential microscopy or scanning potentiometry) can be used to determine the source and drain contributions to the contact resistance directly. In KFM, a conductive atomic force microscope (AFM) tip is scanned over the operational OFET channel twice. On the first pass, the topography



Semiconductor Insulator Gate L 2L 3 3

x L

Potential (V)

Drain

Source

0

15

(1) Topography (2) Surface potential

AFM tip

151

ΔVD

VD = 15 V VG = 75 V

10

5

ΔVS 0

0

(a)

L/3

2L/3

L

Channel position, x (b)

FIGURE 2.4.11 (a) Application of KFM to characterize an OFET. (b) Hypothetical channel potential profile measured by the KFM technique; voltage drops at the source and drain contacts (ΔVS, ΔVD) are measured directly. The dashed line is the ideal (no contact resistance) linear potential profile for an applied drain voltage of 15 V and a gate voltage of 75 V.

of the device is recorded; then, the tip is lifted a small distance (~10 nm) off the device and the second pass retraces the channel topology in the air (or better yet, vacuum) above the sample while the electrostatic potential is recorded (see Figure 2.4.11a). The electrostatic potential data are converted into the OFET surface potential profile by subtracting an appropriate background trace. Thus, KFM measures the full channel potential profile. Since ΔVS and ΔVD are measured directly, calculating the source and drain contributions to the contact resistance can be done without any extrapolation. Figure 2.4.11(b) shows a hypothetical channel potential profile measured by KFM. A clear advantage of KFM over the gated four-probe technique is that the entire channel potential profile is measured experimentally instead of using only two points to extrapolate a linear profile (compare Figures 2.4.10b and 2.4.11b). Thus, other bottlenecks to charge transport (e.g., potential drops at grain boundaries in the channel) can also be visualized using this technique.

2.4.4.4 MEASURED CONTACT RESISTANCE VALUES A collection of experimental OFET contact resistance values is provided in Tables 2.4.1 and 2.4.2. Table 2.4.1 presents contact resistance values for seven different polycrystalline oligomeric semiconductors, including the current benchmark material, pentacene. One must be careful, however, in making comparisons between contact resistance values reported by different groups. It has been proposed that contact resistance is caused by a combination of thermionic emission and carrier diffusion through a depletion region, with the latter dominating in some cases. As a result, contact resistance depends on the level of accumulated charge in the OFET channel, and many experiments have shown that contact resistance is inversely proportional to VG [16,17]. Figure 2.4.12 shows the typical decrease of the source, drain, and channel resistances with increasing gate voltage (carrier density) for a pentacene OFET with gold contacts.


152


TABLE 2.4.1 OFET Contact Resistance Values: Evaporated Oligomer Films Semiconductor Pentacene

PTCDI-C5

PTCDI-C8 PTCDI-C12 PTCDI-C13

Ant-2T-Ant Tet-2T-Tet

Ωcm)b Contact metal TC/BCa RC (Ω Linear acenes (p-channel) Au TC 3 × 104 Au TC 1 × 103 Au TC RS = 3 × 102 RD = 1 × 103 Au BC RS = 4 × 104 RD = 2 × 104 Ag TC 1 × 103 Ag TC 2 × 103 Pt TC 4 × 103 Ca TC 4 × 104 TC 7 × 107

P3HT

Au

F8T2

Soluble oligomer (n-channel) Au BC ~1 × 106

C70-PCBM

Method

Ref.

Four probe

26

R vs. L R vs. L KFM KFM KFM R vs. L KFM R vs. L KFM KFM KFM

13 18 10 10 10 13 10 13 10 10 10

R vs. L

27

a

OFET architecture: top contact (TC) or bottom contact (BC). Total (source resistance, RS + drain resistance, RD) contact resistance, RC, unless individual values are shown. b

Nfree (×1012/cm2) 0

Resistance (Ω)

1011

1.7

4.2

6.7

9.2

VD = −4V

109 RFilm RS RD

107 105 103

V0 0

VT −10

−20

−30

−40

VG (V)

FIGURE 2.4.12 Evolution of the source and drain contact resistances (RS and RD), as well as the channel resistance (RFilm) for a pentacene OFET with increasing carrier density (Nfree), calculated from the capacitance of the Al2O3 dielectric layer and the gate voltage. L = 100 μm, W = 1000 μm.


154


(1 μC/cm2) yields a channel resistance of 1 × 105 Ωcm, which can be compared to the numbers in the tables. Table 2.4.2 presents a collection of contact resistance values measured for a few solution-deposited polymeric and oligomeric semiconductors. These values tend to be higher than in the case of evaporated oligomeric films due to lower mobilities. It has also been observed that, in certain instances, the source contact can constitute a majority of the total contact resistance in polymeric films [10,25].

2.4.5 CONTACT ENGINEERING 2.4.5.1 CHEMICAL MODIFICATIONS As discussed previously, the fundamental reason why contact resistance is generally greater for the bottom contact OFET architecture is the poor semiconductor morphology on top of or near the source and drain contacts that hinders charge injection/extraction. Thus, one of the simplest ways to improve the quality of the semiconductor film on top of bottom contacts is to pretreat the contacts chemically before depositing the semiconductor. In the simplest case, one can use thiol-terminated hydrocarbon molecules to form a self-assembled monolayer (SAM) on top of metallic contacts [14]. The affinity of the thiol group for metal surfaces and SAM formation with such molecules have been well-studied. By forming a CH3-terminated SAM on top of the contacts, the semiconductor layer will no longer “see” the metal, but rather a hydrophobic surface with a different surface energy. As a result, the semiconductor morphology will be modified on top of the contacts — hopefully, in such a way as to improve charge injection/extraction and reduce contact resistance. Taking the idea of SAM formation one step further, one can also use molecules with various chemical functionalities (not only hydrocarbons) for pretreating the contacts. SAMs featuring both electron-withdrawing and electron-donating end groups opposite the thiol-linking ends have been used to alter the charge injection properties of bottom contact OFETs [28,29]. This strategy has the advantage of tuning the energy band line-up at the contact–semiconductor interface in addition to improving semiconductor morphology. Although not as common, there have also been attempts to improve the properties of top contacts by depositing chemical moieties through a shadow mask on top of the semiconductor, immediately prior to contact metal deposition [30]. Two additional examples of manipulating the contact chemistry to achieve lower contact resistance are noteworthy. In the case of the nonmetal contact material poly(3,4-ethylenedioxythiophene) doped with poly(styrenesulfonate), PEDOT:PSS, one group found that low contact resistances were achieved between PEDOT:PSS and the polymer semiconductor F8T2 in a bottom contact OFET [31]. It was posited that the PEDOT:PSS contacts had doped the semiconductor in the vicinity of the contacts, leading to more efficient charge injection/extraction. A second group of researchers found a clever way to make better top contacts to the single-crystal charge-transfer semiconductor DBTTF-TCNQ than with either gold or silver. They used the related metallic charge-transfer material TTF-TCNQ to form top contacts to the semiconductor with success in realizing more efficient charge injection [32].



155

These are only a few examples that show how judicious chemical modifications to the contact–semiconductor interface can often be used to improve contact quality in OFETs.

2.4.5.2 AMBIPOLAR

AND

LIGHT-EMITTING OFETS

While the development of ambipolar (both hole- and electron-transporting) OFETs is still in the early stages, it is certainly an exciting subject within the OFET community. These devices offer not only new possibilities for complementary logic circuit design, but also the potential to control electron–hole recombination within the semiconductor channel to afford light emission. Light-emitting organic fieldeffect transistors (LEOFETs) are particularly intriguing because they possess charge carrier densities in the channel that are orders of magnitude higher than those found in organic light-emitting diodes (OLEDs) [33]. While it is simpler from a fabrication standpoint to deposit the same contact material for both the source and drain contacts (symmetric contacts), one may also consider choosing two different materials for each contact (asymmetric contacts). Based on energy band line-up considerations with the semiconductor HOMO and LUMO levels, depositing two different contact materials at either end of the transistor channel may facilitate more efficient hole and electron injection, respectively. At this point, it is unclear whether separately engineering distinct contacts for hole/electron injection in ambipolar OFETs will prevail over opting for symmetric contacts. However, there will certainly be more reports on this exciting OFET subclass in the next few years.

2.4.5.3 CHANNEL DIMENSIONS: HOW SMALL? At the laboratory scale, OFET channel lengths are typically on the order of 10–100 μm, with W/L ratios ranging from ~10 to over 1,000. Commercialization of organic electronics will lead to a push to make OFET dimensions as small as possible, since the switching speed (cut-off frequency) of an ideal transistor is inversely proportional to L2. As discussed previously, however, one must be careful to avoid becoming contact resistance-limited when shrinking L. For example, from Table 2.4.1, the specific contact resistance for gold top contacts on pentacene can be as low as 1 kΩ·cm. The channel sheet resistance for pentacene devices at high gate voltage (~5 × 1012 carriers/cm2) is about 1 MΩ/sq, assuming a mobility of 1 cm2/Vs. Thus, for a 10-μm channel length pentacene device, the contact resistance is 50% of the total device resistance in the ON state! Clearly, driving the channel lengths smaller will only exacerbate this problem. Thus, ongoing contact engineering efforts to make lower resistance contacts to organic semiconductors will remain important to OFET development and the advancement of organic electronics.

REFERENCES 1. Shen, Y. et al., How to make ohmic contacts to organic semiconductors, Chem. Phys. Chem., 5, 16, 2004.


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2. Kahn, A., Koch, N. and Gao, W., Electronic structure and electrical properties of interfaces between metals and π-conjugated molecular films, J. Polym. Sci. B: Polym. Phys., 41, 2529, 2003. 3. Hill, I.G. et al., Organic semiconductor interfaces: Electronic structure and transport properties, Appl. Surf. Sci., 166, 354, 2000. 4. Seki, K. and Ishii, H., Photoemission studies of functional organic materials and their interfaces, J. Electron Spectroscopy Related Phenomena, 88–91, 821, 1998. 5. Salaneck, W.R. and Fahlman, M., Hybrid interfaces of conjugate polymers: Band edge alignment studied by ultraviolet photoelectron spectroscopy, J. Mater. Res., 19, 1917, 2004. 6. Kera, S. et al., Impact of an interface dipole layer on molecular level alignment at an organic-conductor interface studied by ultraviolet photoemission spectroscopy, Phys. Rev. B, 70, 085304, 2004. 7. Necliudov, P.V. et al., Contact resistance extraction in pentacene thin film transistors, Solid-State Electronics, 47, 259, 2002. 8. Amy, F., Chan, C., and Kahn, A., Polarization at the gold/pentacene interface, Org. Elec., 6, 85, 2005. 9. Wan, A. et al., Impact of electrode contamination on the α-NPD/Au hole injection barrier, Org. Elec., 6, 47, 2005. 10. Bürgi, L. et al., Close look at charge carrier injection in polymer field-effect transistors, J. Appl. Phys., 94, 6129, 2003. 11. Pesavento, P.V. et al., Gated four-probe measurements on pentacene thin-film transistors: Contact resistance as a function of gate voltage and temperature, J. Appl. Phys., 96, 7312, 2004. 12. Pesavento, P.V. et al., Film and contact resistance in pentacene thin-film transistors: Dependence on film thickness, electrode geometry, and correlation with hole mobility, J. Appl. Phys., 99, 094504, 2006. 13. Hamadani, B.H., Natelson, D., Nonlinear charge injection in organic fild-effect transistors, J. Appl. Phys., 97, 064508, 2005. 14. Kymissis, I., Dimitrakopolous, C.D., and Purushothaman, S., High-performance bottom electrode organic thin-film transistors, IEEE Trans. Elec. Dev., 48, 1060, 2001. 15. Dimitrakopolous, C.D., Brown, A.R., and Pomp, A., Molecular beam deposited thin films of pentacene for organic field effect transistor applications, J. Appl. Phys., 80, 2501, 1996. 16. Newman, C. R. et al., High mobility top-gated pentacene thin-film transistors, J. Appl. Phys., 98, 084506, 2005. 17. Zaumseil, J., Baldwin, K.W., and Rogers, J.A., Contact resistance in organic transistors that use source and drain electrodes formed by soft contact lamination, J. Appl. Phys., 93, 6117, 2003. 18. Meijer, E.J. et al., Scaling behavior and parasitic series resistance in disordered organic field-effect transistors, Appl. Phys. Lett., 82, 4576, 2003. 19. Puntambekar, K.P., Pesavento, P.V., and Frisbie, C.D., Surface potential profiling and contact resistance measurements on operating pentacene thin-film transistors by Kelvin probe force microscopy, Appl. Phys. Lett., 83, 5539, 2003. 20. Maltezos, G. et al., Tunable organic transistors that use microfluidic source and drain electrodes, Appl. Phys. Lett., 83, 2067, 2003. 21. Chesterfield, R.J. et al., Variable temperature film and contact resistance measurements on operating n-channel organic thin film transistors, J. Appl. Phys., 95, 6396, 2004.



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22. Chesterfield, R.J. et al., Organic thin film transistors based on n-alkyl perylene diimides: Charge transport kinetics as a function of gate voltage and temperature, J. Phys. Chem. B, 108, 19281, 2004. 23. Gundlach, D.J. et al., High mobility n-channel organic thin-film transistors and complementary inverters, J. Appl. Phys., 98, 064502, 2005. 24. Merlo, J.A. et al., p-Channel organic semiconductors based on hybrid acene-thiophene molecules for thin-film transistor applications, J. Am. Chem. Soc., 127, 3997, 2005. 25. Street, R.A. and Salleo, A., Contact effects in polymer transistors, Appl. Phys. Lett., 81, 2887, 2002. 26. Panzer, M.J. and Frisbie, C.D., Unpublished results. 27. Anthopoulos, T.D. et al., Solution processible organic transistors and circuits based on a C70 methanofullerene, J. Appl. Phys., 98, 054503, 2005. 28. Campbell, I.H. et al., Controlling Schottky energy barriers in organic electronic devices using self-assembled monolayers, Phys. Rev. B, 54, 14321, 1996. 29. Gundlach, D.J., Jia, L., and Jackson, T.N., Pentacene TFT with improved linear region characteristics using chemically modified source and drain electrodes, IEEE Elec. Dev. Lett., 22, 571, 2001. 30. Schroeder, R., Majewski, L.A., and Grell, M., Improving organic transistor performance with Schottky contacts, Appl. Phys. Lett., 84, 1004, 2004. 31. Wang, J.Z., Chang, J.F., and Sirringhaus, H., Contact effects of solution-processed polymer electrodes: Limited conductivity and interfacial doping, Appl. Phys. Lett., 87, 083503, 2005. 32. Takahashi, Y. et al., Tuning of electron injections for n-type organic transistor based on charge-transfer compounds, Appl. Phys. Lett., 86, 063504, 2005. 33. Zaumseil, J., Friend, R.H., and Sirringhaus, H., Spatial control of the recombination zone in an ambipolar light-emitting organic transistor, Nat. Mater., 5, 69, 2006.



3.1

Design, Synthesis, and Transistor Performance of Organic Semiconductors

Abhijit Basu Mallik, Jason Locklin, Stefan C. B. Mannsfeld, Colin Reese, Mark E. Roberts, Michelle L. Senatore, Hong Zi, and Zhenan Bao CONTENTS 3.1.1 Introduction................................................................................................160 3.1.2 p-Channel Organic Semiconductors ..........................................................161 3.1.2.1 Acenes and Non-Thiophene-Based Semiconductor ...................162 3.1.2.1.1 Linear Fused Rings....................................................162 3.1.2.1.2 Fused and Extended Heteroarenes ............................166 3.1.2.1.3 Star-Shaped Oligomers..............................................173 3.1.2.1.4 Oligoaryls ..................................................................173 3.1.2.1.5 Macrocyclics..............................................................175 3.1.2.2 Thiophene-Based Oligomers.......................................................177 3.1.2.3 Polymers ......................................................................................183 3.1.2.4 Solution Processable Semiconductors: The “Precursor Method”.......................................................................................188 3.1.2.4.1 Precusor Polymers and Small Molecules .................189 3.1.2.4.2 Polydiacetylene..........................................................191 3.1.3 N-Channel Organic Semiconductors .........................................................191 3.1.3.1 Fullerenes and Fullerene Derivatives..........................................192 3.1.3.2 Phthalocyanines ...........................................................................192 3.1.3.3 Naphthalene Diimide Derivatives ...............................................194 3.1.3.4 Perylene Diimide Derivatives .....................................................195 3.1.3.5 Quinoid Systems .........................................................................197 3.1.3.6 Thiophene Based n-Channel Oligomers .....................................197 3.1.3.7 Trifluoromethylphenyl-Based Oligomers....................................200 3.1.3.8 Polymeric Systems ......................................................................201 159


160


3.1.4 Outlook and Conclusions ..........................................................................202 3.1.5 Table of Mobilities.....................................................................................203 References..............................................................................................................214

3.1.1 INTRODUCTION Organic materials have been an integral part of the semiconductor industry since 1960 when Kahng and Atalla demonstrated the first metal-oxide-silicon field-effect transistor.[1] Photosensitive polymers have been the defining factor in achieving feature sizes down to and below 45 nm, while other insulating polymers have been vital for packaging chips or improving the performance of low-k interlayer dielectric materials.[2] However, organic materials had not been considered as the active semiconductor layer until 1987, when Koezuka and coworkers demonstrated a polythiophene-based field-effect transistor.[3] Since this initial breakthrough, remarkable progress has been made in organic field-effect transistors (OFETs), due largely to the design of high-performance active-layer materials. Organic materials are attractive for many components of electronic devices, particularly the active semiconductor layer, due to many fundamental advantages over their inorganic counterparts.[4] Simple, solution-based processing allows for unconventional deposition methods,[5] such as inkjet,[6-10] screen,[11,12] and microcontact printing.[5,13,14] The low temperature required for these methods, combined with the mechanical flexibility of organic materials, offers compatibility with plastic substrates,[5] leading to the possibility of flexible integrated circuits[15-17], electronic paper,[18-21] and roll-up displays.[22,23] Although organic materials are not currently suitable for applications requiring high switching speeds, their low material and fabrication costs make them ideal for large-area applications, such as displays or solar cells.[24-26] Disposable electronics is another emerging class of technology based on these advantages, including chemical sensors[27-31] and radio frequency identification cards (RFIDs).[32] Already, the performance of organic materials, such as pentacene[33] and rubrene,[34] has surpassed that of amorphous silicon, while others have been incorporated in commercially available light-emitting diode displays.[35-37] This impressive growth is evident in Figure 3.1.1, which shows a logarithmic plot depicting the evolution of reported field-effect mobility values for p- and n-channel organic semiconductors over the past 18 years. Arguably the most important feature of organic materials, however, is the ability to impart functionality by intelligent material design[5,38-40] and advanced synthetic techniques. This aspect of materials engineering is a necessary component of the success of organic electronics and the development of molecular devices. The design of materials for organic electronics has proceeded through an enormous synthetic initiative, and has attracted theoretical treatment with the prospect of the rational design of high-performance materials. A variety of approaches have given valuable insight to both the relationship between functionality and electronic properties, as well as between electron transport efficiency and molecular superstructure. While existing studies do not currently offer the prediction of performance a priori, the results presented offer the hope for true materials engineering. However, while materials have been the overwhelming focus of the organic electronics field, the



161

101

Mobility (cm2V1S–1)

100

a–Si:H

10–1 10–2 10–3 10–4 10–5

P-channel Pentacene (V) Rubrene (sc) Other small mols. (V) Small mols. (S) Polymers (S) Air-stable n-channel

1986 1988 1990 1992 1994 1996 1998 2000 2002 2004 2006 Year

FIGURE 3.1.1 Evolution of OFET performance with time for various p-channel (pentacene, rubrene, other small molecules, and polymers) and n-channel organic semiconductors. (v): Vacuum deposition; (s) solution deposition; (sc): single crystal. A range of mobilities for hydrogenated amorphous silicon (a-Si:H) is shown for reference. Please refer to Table 3.1.1 for the respective molecule indices.

fabrication of high-performance devices depends not only on the intrinsic properties of the semiconducting material, but also the technique employed to form the active layer. The deposition method and conditions, in addition to any dielectric or postdeposition treatments, will impact device performance by influencing the molecular organization of the semiconductor molecules.[30,41-43] For this reason, it is crucial to understand the various factors that influence the dynamics of thin-film formation. In particular, the initial growth regime — of particular importance for organic thinfilm transistors (OTFTs) — has been well-characterized, and provides a framework within which the key parameters of ordered film growth may be examined. The rate of publication and development of new materials for organic electronics suggests that the coming years will bring the realization of these unique technologies. In this review, we attempt to summarize comprehensively the vast effort toward the development of high-performance organic semiconductors for OFETs. The body of this paper will focus on the extensive collection of semiconducting organic materials synthesized to-date, compiling structural and performance data as reported in the literature.

3.1.2 p-CHANNEL ORGANIC SEMICONDUCTORS The design and synthesis of p-channel organic semiconductors is an area of fervent chemical research. Hundreds of novel molecules with unique optical and electronic properties have been synthesized, almost exclusively in the past ten years. Since the first report of a transistor employing an organic active layer in 1987,[3,44] the performance of small molecule and polymer semiconductors has been steadily


162


improved. In this section we survey the OFET performance of these materials, focusing on two parameters that determine largely the potential speed and efficiency of circuits containing them: the field-effect mobility and the on-off current ratio.

3.1.2.1 ACENES

AND

NON-THIOPHENE-BASED SEMICONDUCTOR

3.1.2.1.1 Linear Fused Rings Polycyclic aromatic hydrocarbons are currently among the most widely studied organic π-functional materials. Particular attention has been paid to linear aromatic hydrocarbons composed of laterally fused benzene rings, called linear acenes or oligoacenes. The highest occupied molecular orbital (HOMO) energy level in linearly condensed [n]acenes significantly increases with n, which facilitates the injection of radical cations (holes) at the interface between the source and semiconductor layer under an applied gate voltage. Furthermore their planar shape facilitates crystal packing and enhances the intermolecular overlap of π-systems. Because of these features, pentance and tetracene, both members of the linear oligoacene series, are among the most promising molecular semiconductors for OFETs. However, it is well documented that preparation of such extended benzenoid molecules is plagued by facile photo-oxidation, insolubility and most detrimental, dimerization and polymerization.[45-47] This section summarizes published data on acene molecules emphasizing the influence of the structural properties and the molecular ordering on OFET performance. Anthracene (1a)(Figure 3.1.2) is the smallest member of the acene series with reported transistor characteristics. Single-crystal OFETs have shown a mobility of 0.02 cm2V–1s–1, but only at very low temperatures.[48] However, rather high mobilities have been observed using oligoanthracene and substituted anthracene derivatives. Substitution at the 2,6–positions on the anthracene backbone is expected to give the most extended π-conjugation and also the highest degree of planarity, due to the retention of linearity and limited steric hindrance. OTFTs fabricated from thermally evaporated thin-films of this type of oligomer gave field-effect mobilities in the order 1b< 1d< 1c< 1e (0.01 cm2V–1s–1, 0.07 cm2V–1s–1, 0.13 cm2V–1s–1, and 0.18 cm2V–1s–1 respectively).[49,50] This observation suggests that the addition of alkyl groups (from 1b to 1c and 1d to 1e) is more effective in improving mobility than the extension of π-conjugation (from 1b to 1d and 1c to 1e), which is also observed for oligothiophenes.[51] Acene and thiophene oligomers represent two of the most heavily studied series of organic semiconducting compounds. Two subsequent examples present anthracene-based materials with thiophene groups centralized on the conjugated core and also on the periphery. In the first example, thiophene units flank an anthracene moiety, as shown in molecules 1f and 1g.[53] The mobility of the dihexyl-substituted oligomer, 1g (0.5 cm2V–1s–1), was nearly an order of magnitude higher than that of 1f (0.063 cm2V–1s–1), presumably due to the self-assembly promoted by the alkyl side chains.[54] OTFTs fabricated using either material showed high stability under ambient conditions. In a second example, anthracene moieties were substituted on a bithiophene core to give molecule 1h.[55] OTFTs employing this compound had a mobility of 0.12



Anthracene based (1a–1i)

163

R

n

R (1a)

n = 0, R = H (1b) n = 0, R = C6H13 (1c) n = 1, R = H (1d) n = 1, R = C6H13 (1e)

S S R

R

S

S

(1h) R = H (1f) = C6H13 (1g)

Tetracene based (2a–2f)

(1i) R

(2a)

R = Cl: R′ = H (2b) = Br: R′ = H (2c) = Cl: R′ = Cl (2d) = Br: R′ = Br (2e)

R′ S S (2f)

FIGURE 3.1.2 Chemical structure of anthracene- and tetracene-based oligoacenes.

cm2V–1s–1 with an on/off ratio of 108 and also exhibited a high degree of thermal stability. Among the two unsubstituted oligomers, 1f and 1h, the latter has higher number of acene moieties and longer effective conjugation length, resulting in a higher mobility. However, the hexyl-substituted oligomer 1g, despite of having lower π-conjugation showed a higher mobility than 1h, further illustrating the importance of alkyl side chains in the molecular packing of thin films. Bis(2-acenyl)acetylene (1i) has also been used as the active layer in OTFTs, achieving a mobility as high as 1.1 cm2V–1s–1 with an on/off ratio of 4.4x105.[56] Tetracene (2a) (Figure 3.1.2) is the next larger homologue of anthracene and therefore, has a slightly higher degree of conjugation. OFETs fabricated from single crystals of tetracene show a mobility of up to 1.3 cm2V–1s–1 with an on/off ratio of 106, [57] whereas thin-films yield a mobility of 0.1 cm2V–1s–1.[58] Semiempirical calculations have shown that substitution of bromo or chloro groups on tetracene lowers HOMO and LUMO levels. The halogen groups also promote cofacial πstacking.[59] Halogenated tetracene derivatives, 2b-e, differ in position of substitution and hence provide a useful system to investigate the effect of molecular packing on charge transport.[60] Single crystal X-ray crystallography showed that the number of halogen substitutions is a key parameter and can be used to control molecular packing. The mono-substituted tetracene derivatives 2b and 2c are shown to pack in a herringbone fashion whereas the di-substituted analogues 2d and 2e π-stack in a slipped face-to-face manner. OFETs fabricated using single crystals of these molecules gave a broad range of mobilities, ranging from 1.4 × 10–4 cm2V–1s–1 (2b)


164


to 1.6 cm2V–1s–1 (2d). The π-stacking structure of 2d, which enhances π-orbital overlap, may be responsible for this high mobility. Similar to the anthracenethiophene oligomer 1h, the tetracene analogue 2f also showed a high mobility (0.5 cm2V–1s–1 with on/off ratio of 108).[55] Pentacene (3a) (Figure 3.1.3) is one of the most widely studied organic semiconductors, with characterization beginning in the 1970s. The highest thin-film fieldeffect mobilities so far have been recorded for pentacene (0.3–0.7 cm2V–1s–1 on SiO2/Si substrates, 1.5 cm2V–1s–1 on chemically modified SiO2/Si substrates, and 3–6 cm2V–1s–1 on polymer gate dielectrics).[57,61-66] Despite the substantial research Pentacene Based (3a-3j) SiiPr 3

SiR 3

X

F (3a)

F R2

R1

X

F

X X

F R2

R1 SiiPr 3

SiR 3

R 1 ,R 2= Me (3b) R= Cl (3c) R 1 = Me; R2 = H (3d) R 1 = C 6H 13 ; R 2= H (3e)

R= Me ( 3f ) R= Et (3g) R= iPr (3h)

X= H (3i) X= F (3j)

Quinones (4a-4c) O

O

CH 3

CH3 (4a)

O CH 3

CH3

(4b)

O

O

O

(4c)

Rubrene derivatives (5a-5b)

( 5b)

(5a )

Polycyclic Aromatic Hydrocarbons (6a-6f) R1

R2

C12H 25O

OC 12 H25

R1-6= H (6a) R1,2,4,5 =H; R 3,6 = C 6H 13 (6b) R3 R1,4=H; R 2,3,5,6 = C 6H 13 (6c) R1-6= C 6H 13 (6d) R1-6= Ph-p-C12 H25 (6e)

R6

R5

R4

(6f)

C12H 25O

OC 12 H25

FIGURE 3.1.3 Chemical structure of pentacene- and rubrene-based oligoacenes.



165

on pentacene devices, relatively few derivatives of pentacene have been synthesized. An obvious advantage of substituted pentacene is improved solubility, which would simplify purification and transistor processing. It has been suggested that the mobility of pentacene could be further enhanced if it is forced to pack in a face-to-face manner rather than the usually observed herringbone structure. One approach to disrupting the herringbone packing is to substitute the four terminal hydrogen atoms. In such a structure, the pentacene core would remain intact while steric interactions would prohibit the herringbone arrangement. Hence, methyl groups (3b) were introduced at the four terminal carbon atoms of pentacene.[67] Thin films of Me4-pentacene (3b) were thermally evaporated and exhibited a mobility of 0.3 cm2V–1s–1 with an on/off ratio of 6.3x103. However, the incorporated methyl groups are not sufficiently large enough to disrupt the herringbone packing and the crystal structure of 3b still resembles that of pentacene. Dimethyl (3d) and dihexyl (3e) pentacene derivatives also exhibited high mobility values of 2.5 cm2V–1s–1 and 0.251 cm2V–1s–1, respectively.[61,68] Tetrachloro pentacene (3c) was also synthesized, but due to its instability, it failed to show any OFET characteristics.[46] Bulky solubilizing trialkylsilyl groups have been substituted at the 6,13-positions of the pentacene molecule (3f-3h)[69,70] in order to impart solubility and disrupt herringbone packing of pentacene core. The substituents were separated from the pentacene core by a rigid alkyne spacer to allow the closest possible contact between the aromatic rings. The functionalized derivatives are very soluble and can be crystallized from common organic solvents. Solid-state packing analysis of triisopropylsilyl (TIPS) pentacene 3h showed that the molecule stacks in a 2D columnar array with significantly larger overlap of the aromatic rings compared to unsubstituted, herringbone-packed pentacene. OTFTs fabricated with TIPS-pentacene (3h) gave a very high mobility of 0.4 cm2V–1s–1 with an on/off ratio of 106, while those fabricated with trimethylsilyl (TMS, 3f) and triethylsilyl (TES, 3g) showed low mobilities of 10–5 cm2V–1s–1 The high field-effect mobility of 3h corroborates its improved thin-film structure compared to 3f and 3g. Fluorine containing derivatives of silylethynylated pentacenes (3i-3j) have also been synthesized[71] in attempt to tune both the electronic properties and thin-film packing of the pentacene derivatives. Fluorine-substituition has been used to lower the HOMO and LUMO energy levels[72-74] while the aryl-fluoroaryl interactions are well-known packing synthons in crystal engineering.[75-77] OTFTs showed pchannel behavior for both fluorine-containing compounds, contrary to expectations based on reports of other fluorinated derivatives found to behave as n-channel semiconductors.[72] Tetrafluoropentacene (3i) showed a mobility of 0.014 cm2V–1s–1, while octafluoropentacene (3j) showed a higher mobility of 0.045 cm2V–1s–1. While the longest commercially available oligoacene, pentacene, also exhibits the highest OTFT performance, very few longer acenes or derivatives thereof have been reported. Recently silylethynylated hexacene and heptacene derivatives were synthesized (Figure 3.1.3), but their transistor characteristics were not reported.[78] The only transistor examples of higher-acenes based semiconductors are molecules containing an electron-rich acene unit and an electron-poor quinoid unit (4a4c).[79] The packing influenced by this structure is such that acene units stack over


166


quinoid fragments in a self-complimentary manner.[80] The highest mobility was obtained for devices fabricated from thermally evaporated thin films of hexacene quinone (4a). A field-effect mobility of 5 × 10–0.05 cm2V–1s–1 and on/off ratio of 106 were reported. Despite the strong electron-withdrawing nature of quinone, these molecules acted as hole-transporting materials. Rubrene (5a) (Figure 3.1.3), a commercially available aryl-substituted linear acene, is soluble in common solvents such as toluene, xylene, and chloroform, unlike its unsubstituted counterpart, tetracene (2a). A tetra-t-butyl derivative of rubrene (5b) has also been reported. OTFTs fabricated from thin films of 5a and 5b show low mobilities in the range of 10–3 cm2V–1s–1.[81] The mobility of thin films of rubrene, 5a can be further improved up to 0.07 cm2V–1s–1 by insertion of a pentacene buffer layer in between rubrene thin films and sapphire substrate.[82] Recently a new method was introduced to fabricate high-performance field-effect transistors using rubrene single crystals, where mobilities as high as 15.4 cm2V–1s–1 have been observed.[34,57] Furthermore a solution-processed semiconductor composite film of rubrene containing a vitrifying additive and a high-molecular weight polymer gave mobilities of up to 0.7 cm2V–1s–1 with an on/off ratio of greater than 106.[83] Recent research reported ambipolar field-effect transitors based on rubrene single crystals. Hole and electron mobilities of 1.8 and 0.011 cm2V–1s–1 were derived from saturated currents.[84] Disc-shaped aromatic hexabenzocoronene (HBC) derivatives (6a-6d) (Figure 3.1.3) have also been tested as an active layer in OFETs.[85] HBC combines structural elements of linear acenes with a disc-shaped, potentially columnar liquidcrystalline core. The mobilities of dihexyl (6b) and tetrahexyl (6c) derivatives (0.011 cm2V–1s–1 and 0.012 cm2V–1s–1 respectively) were ten-fold higher than those of unsubstituted HBC (6a) and the hexahexyl derivative (6d). This good performance was explained by the 2D self-assembling properties of dihexyl and tetrahexyl HBC derivatives. Recently, a magnetic field has also been utilized to oriented thin films of 6e. Such oriented films gave mobilities of up to 0.001 cm2V–1s–1.[86] A new type of HBC derivative, 6f, with nonplanar aromatic core, imposed from steric congestion of its proximal carbon atoms, has been found to be semiconducting.[87] The single crystal structure of 6f confirmed that the aromatic core is severely distorted and the molecule self-assembles into long columns through stacking. Field-effect transistors fabricated from spin-cast films of 6f gave a mobility of 0.02 cm2V–1s–1 and an on/off ratio of 106. 3.1.2.1.2 Fused and Extended Heteroarenes Heteroarenes are an interesting but much less studied class of π-functional materials. Substitution of one of several carbon atoms in oligoacenes with a heteroatom such as nitrogen or sulphur in different valence state may induce unique properties in these π-electron systems. The generally decreased HOMO-LUMO gap is of particular interest for many applications. Several dihydrodiazapentacenes, (7a-7d) (Figure 3.1.4) have been synthesized with nitrogens replacing carbons in the phenyl rings of pentacene,[88] in order to avoid the facile oxidation of pentacene.[47] OTFTs fabricated by thermal evaporation



Dihydrodiazapentacene (7a-7d)

Imidazolylquinoline (8a-8b)

H N N H (7a)

R= H (7b) R= CH3 (7c)

H N

R

R

N H

R

R

N N N H R= H (8a) (8b)

R=

(7d)

R=

Thioacene (9a-9g)

SiR3 S

R

R S R= R= R= R=

S

H (9a) C6 H13 (9b) C12H 25 (9c) C18H 37 (9d)

S

Bis- and oligo-(benzodithiophene) (10a-10f) S S

SiR3

R= Me (9e) R= Et (9f) R= iPr (9g)

S

(10a)

S C 6H 13

S

S

S

C 6H 13

S (10b)

S

S

S S

C 6H 13

S (10d)

S

S

S

C 6H 13

S (10c)

S

S

S

S S

S (10e)

(10f)

Dithienothiophene Derivatices (11a-11f) S

S S

S

S S

S

(11a)

S

S

S

(11d) S

S

(11b)

S S

S S

S C6 H13

S S

S S

S

S

C6H 13

(11c)

S

(11e)

S S

S S

S

Substituted Fused-Bithiophene (12a-12b) S

S S

167

S (12a)

FIGURE 3.1.4 Chemical structure of thioacene-based oligomers.

(12b)

(11f)


168


of 7a and 7b gave field-effect mobilities of 5 × 10–5 cm2V–1s–1 and 3–6 × 10–3 cm2V–1s–1, respectively. Two imidazolylquinoline compounds (8a-8b) (Figure 3.1.4) have also been tested in OTFTs [89,90] with thermally evaporated thin films showing mobilities of 0.038 cm2V–1s–1 and 0.148 cm2V–1s–1 for 8a and 8b, respectively. Anthradithiophene (ADT, 9a) (Figure 3.1.4), a fused heterocyclic compound similar to pentacene is expected to have a higher oxidation barrier than pentacene, which would impart increased stability.[91] OTFTs of the soluble dihexyl-substituted dihexylanthradithiophene (DHADT, 9b) were fabricated by solution-casting at different temperatures. The field-effect mobility was found to be strongly dependent on the solvent evaporation temperature. The highest mobilities, 0.01–0.02 cm2V–1s–1 were obtained for a substrate temperature of 100°C. Films of DHADT (9b) obtained by thermal evaporation were highly polycrystalline in nature and gave mobilities as high as 0.15 cm2V–1s–1. Furthermore, compound 9d has a mobility of 0.06 cm2V–1s–1 even though 70% of its molecular volume is comprised of hydrocarbon chains. These alkylated anthradithiophenes combine a pentacene-like high mobility with greater solubility and oxidative stability. As with pentacenes, silylethynylated anthradithiophene derivatives (9e-9g) (Figure 3.1.4) with substituents on the central aromatic ring have also been synthesized.[92] OTFTs fabricated with the trimethyl (9e) and triisopropyl (9g) derivatives showed negligible transistor characteristics while those with triethyl derivative (9f) deposited by solution gave a mobility of 1.0 cm2V–1s–1 with on/off ratio of 107. The excellent performance of 9f was attributed to the enhanced π-orbital overlap observed in its crystalline thin-film structure. A linearly fused heteroarene, Bis(benzodithiophene) (10a) (Figure 3.1.4), in which benzenes are embedded in a fused-ring system, was synthesized in attempt to limit the conformational freedom and thereby reducing interrupted conjugation that is possible in oligothiophenes.[38,93] The benzodithiophene system was further extended to a fully fused-ring compound with dibenzo[b,b′]thieno[2,3-f:5,4f ′]benzothiophene (DBTBT, 10d), thereby completely constraining conformational freedom.[94] OTFTs comprised of thermally evaporated thin films of 10a and 10d gave mobilities of 0.04 cm2V–1s–1 and 0.15 cm2V–1s–1, respectively. The greater mobility of 10d compared to 10a is expected, due to its extended π-conjugation. However, further substitutions of 10a with hexyl (10b) or thiophene-hexyl (10c) units gave lower mobilities on the order of 10–3 cm2V–1s–1.[38] Another two extended fused-ring compounds, isomer-pure thieno[f,f′]bis[1]benzothiophens,syn and anti with respect to the orientation of thiophenes along the long molecular axis of 10e and 10f were reported. The field-effect mobilities of both regioisomers were as high as 0.12 cm2V–1s–1.[95] With the same goal of less conformational freedom and stronger inter-and intramolecular π-overlap, a new fused compound, bis(dithienothiophene) (BDT, 11a)[96-98] and its fused counterpart pentathienoacene (PTA, 11b)[99] have been synthesized (Figure 3.1.4). PTA is attractive because it combines the molecular shape of pentacene with a thiophene monomer. The linear condensed thiophene backbone not only imparts extended π-conjugation, but is also more planar than BDT. OTFTs of 11a yielded high mobilities of 0.05 cm2V–1s–1 with a high on/off ratio of 108,



169

whereas for a PTA (11b) based OFET, a mobility of 0.045 cm2V–1s–1 and an on/off ratio of 103 were achieved. Oligothiophene molecules with a central fused thiophene dithienothiophene (11c) demonstrate reversible oxidation at low potentials with similar HOMO levels to the equivalent dihexyl-sexithiophene (28i) (Figure 3.1.10). Thermally evaporated films on silicon oxide dielectrics showed a field-effect mobility as high as 0.02 cm2V–1s–1 with an on/off ratio greater than 106.[100] Recently, a series of new organic semiconductors (11d-f) were reported, using dithienothiophene as the core. The compounds exhibit excellent field-effect performance with a high mobility of 0.42 cm2V–1s–1 and an on/off ratio of 5 × 106.[101] A structural combination of fused bithiophene with fluorene (BFTT, 12a) or biphenyl units (BPTT, 12b) has been fabricated (Figure 3.1.4).[102] Thin films of BFTT (12a) and BPTT (12b) showed mobilities of 0.06 cm2V–1s–1 and 0.093 cm2V–1s–1, respectively. The slightly higher mobility of BPTT (12b) may be due to its more planar and straight conformation, which leads to better thin-film packing. A series of novel fused heteroarenes, naphtha[1,8-bc:5,4-b′c′]dithiophene (NDT) and its derivatives (13a-13f) (Figure 3.1.5) have been synthesized and tested in OTFTs.[103,104] These fused heteroarenes are formally isoelectronic with pyrene, but their aromatic character is greatly reduced by the loss of benzene rings in the skeleton. As a result, the HOMO energy levels of these π-systems are raised, whereas the LUMO energy levels are lowered. OFETs fabricated from thermally evaporated thin films of these materials showed mobilities from 10–4–0.1 cm2V–1s–1. Although the incorporation of alkyl chains tends to improve the mobility of semiconductor due to better thin-film formation,[105] NDT with (13a) and without (13b) alkyl chains showed a low mobility of 10–4 cm2V–1s–1. However 13c, the more conjugated version of 13a, gave a higher mobility of 6 × 10–3 cm2V–1s–1. Replacing the thiophene groups with phenyl units led to improved performance. The naphthyl derivative 13f showed the highest mobility of 0.11 cm2V–1s–1 with an on/off ratio of 105. Novel butterfly-shaped pyrene derivatives have also been synthesized and tested in OTFTs. The pyrene derivative 13g does not show any field-effect behavior, while 13h exhibited a mobility of 3.7 × 10–3 cm2V–1s–1 with an on/off ratio of 104.[106] Amorphous thin films of spiro-linked compounds (14a-d) (Figure 3.1.5) have been tested in OTFTs.[107,108] The characteristic structural feature of these materials is the linkage of two hole transporting units by a spiro junction. A mobility of 7 × 10–5 cm2V–1s–1 was obtained with an on/off ratio of 106. A new class of semiconductors based on a tertiary diamine structure, 5,11-disubstituted indolo[3,2-b]carbazole (15a-15d) (Figure 3.1.5) have been reported.[109,110] The derivatives are soluble in organic solvents such as toluene, chloroform, and chlorobenzene. OTFTs using these carbazole derivatives (15a-15d) exhibited p-channel behavior with mobilities of up to 0.12 cm2V–1s–1 and a current on/off ratio of 107. Recently dichloro derivatives (15e-15f) of these carbazoles gave improved device performance with mobility of 0.14 cm2V–1s–1.[111] 2,7-Carbazolevinylene-based conjugated oligomers, 15h-15j, have also been tested in OTFTs.[112,113] The higher mobility of 15h (0.30 cm2V–1s–1 with on/off ratio of 107) compared to 15j (10–6 cm2V–1s–1 with on/off ratio of 103) is attributed to the relative coplanar structure of the former and twisted biphenyl units in latter.


170

Organic Field-Effect Transistors Naphtha[1,8-bc:5,4-b'c']dithiophene and pyrene derivatives (13a-13h) S R

F3C

R

CF3

S

S

R=

S S

S (13a)

C6 H13

S

S

(13b)

(13c)

S F3C

R= (13d)

(13e)

S

CF3

(13g)

(13h)

(13f)

Spiro-linked compounds (14a-14d) Ar

Ar N

N

N

N Ar

Ar Ar=

Ar

Ar N

N

N

N Ar

Ar

(14a)

Ar=

(14b) (14c)

(14d)

Carbazole derivatives (15a-15j) R N

C12H 25

N

CH 3

R1 N

R2

N R= octyl (15a) R R= dodecyl (15b) R= 4-octylphenyl (15c) R= 4-methylphenyl (15d)

C8 H17

R2

N R1

C 12H 25

H3 C

R 1 = Cl, R 2= H (15e) R 1 = H, R 2= Cl (15f)

N C8H 17

(15g)

R2 N R1

R1 N R2

Ph

Ph

R1 = H, R 2= C6H 13 (15h) R1 = C 6H 13, R 2= CH3 (15i)

N C 6H 13

(15j)

Thiazolothiazole and Thiazole derivatices (16a-16q) N

S

S

N

Ar

Ar S

Ar=

O (16f)

(16a) S

S

(16b)

S S S O N

S

(16g)

N

S

O

(16c)

C6 H 13

(16h) O N

S

(16d)

(16i) S

N (16e)

(16j)

OMe

S (16k) S

C6 H13

N

S S

S

S

N

N

S S

S

S S R= H (16o) = C6 H13 (16p)

S

S

S

N

R= H (16l) = C 6H 13 (16m)

R

N

S

R

R

(16n)

C6H13 S S

R

FIGURE 3.1.5 Chemical structure of heteroarene oligomers.

N S (16q)

S S



171

Thiazole rings in thiophene oligomers are known to reduce the steric interaction due to the absence of one hydrogen atom.[114] In addition, thiazole is an electronwithdrawing heterocycle and hence, imparts stability to oxygen.[115] Thiazole oligomers also form π-stacked structures.[116] Thiazolothiazole is an analogue of thiazole and its co-oligomers with thiophene (16a-16c), furan (16d, 16f, 16h), thiazole units (16e, 16g, 16i), phenyl (16j) and benzothiophene (16k) have been tested in OTFTs (Figure 3.1.5).[117-119] These oligomers are donor-acceptor compounds (donor: thiophene; acceptor: thiazolothiazole) with lowered energy gaps. The unsubstituted derivative (16a) failed to show any field-effect, whereas thiophene(16b) and hexyl-thiophene (16c) derivatives showed mobilities of 0.02 cm2V–1s–1 and 0.003 cm2V–1s–1, respectively, with on/off ratios of 104. OTFTs of 16d gave mobilities up to 2 × 10–3 cm2V–1s–1 with an on/off ratio of 104, however 16f failed to show any field-effect mobility. The thiazole (16e, 16g, 16i), phenyl (16j) and benzothiophene (16k) derivatives did not show any field-effect mobility, either. The lower mobility of furan substituted molecule (16d) compared to its thiophene counterpart (16b) was attributed to its smaller size and decreased polarizability, which increases columbic repulsion and decreases intermolecular interactions. Unsubstituted (16o) and hexyl-substituted (16p) bisthiazole-thiophene co-oligomers were also reported, but had rather low mobilities of 1 × 10–5 cm2V–1s–1 and 2 × 10–5 cm2V–1s–1, respectively.[120] With increasing effective conjugation length by incorporation of additional thiophene rings, as in 16l and 16m, mobilities of 0.011 cm2V–1s–1 and 3.5 × 10–4 cm2V–1s–1, respectively, were obtained.[120] The lower performance of thiazole-containing oligomers compared to the corresponding oligothiophenes was attributed to a larger charge injection barrier and less favorable thin-film morphologies. Tetrathiafulvalene (TTF) and its derivatives (Figure 3.1.6) have been intensively investigated as building blocks for charge-transfer salts, producing a multitude of organic conductors and superconductors.[45,121-124] It was recently shown that TTF derivatives can also be used in OFETs. Single-crystal transistors of dithiophenetetrathiafulavalene (DT-TTF, 17c) were prepared by drop casting of a warm, saturated chlorobenzene solution. The maximum mobility observed in these crystals was 1.4 cm2V–1s–1.[125-128] Similarly, the mobility of single crystal (thiophene)(thiomethylene)-tetrathiafulvalene (TTDM-TTF, 17a) was found to be 0.4 cm2V–1s–1. TTF derivatives with fused aromatic rings (benzo-17h, naphtha-17j, pyrazino-17i and quinoxalino-17k) have also been synthesized, with mobilities of 0.06 cm2V–1s–1 and 0.2 cm2V–1s–1 obtained for single crystals of 17h and 17k.[129,130] The high performance of TTF derivatives suggests that bis(1,3-dithiol-2-ylidene) compounds (Figure 3.1.6), π-extended TTF analogues with a conjugated spacer group, may also be attractive candidates for OFETs. Thin films of Bis[1,2,5]thiadiazolo-p-quinobis(1,3-dithole) (BTQBT, 18a)[131-133] show a mobility of 0.2 cm2V–1s–1 with an on/off ratio of 108. BTQBT can be modified by introducing substituents on 1,3-dithiole ring and replacement of fused dithiadiazole with different π-spacers (18b-18j).[134] OTFTs fabricated from these molecules showed zero to negligible mobility (10–7–10–4 cm2V–1s–1). The replacement of sulphur in heterocycles such as thiophene has been attempted with chalcogens from the same group, such as selenium and tellurium.


172

Organic Field-Effect Transistors TTF derivatives (17a-17k) X2 X1

S

S

S

S S

S

S S

X1 = S, X2= H (17a) X1 = H, X2= S (17b) S

S

S

S

S

S (17g)

X2 X1

S

S S

X

S

S

S

S

S

X

S

S S X= C (17h) = N (17i)

X

S

X1 = H, X2 = S (17d) X1 = S, X2 = H (17e)

DT-TTF (17c) S

S

X

S

S

S

S

S

S (17f)

S

X

S

S

X

X

S

S

X

X= C (17j) = N (17k)

Bis(1,3-dithiol-2-ylidene) compounds with conjugated spacer (18a-18j) S

R

S

R R

N

N

S

S

S N

S

S

S

N N S

S

S S

N

S

R S

X

N Se

N S

S

S

S

R R R R= H (18e) = Benzo (18f)

R= H (18c) = Benzo (18d) BenzoBTQBT (18b)

N

X

S

S R

S

S

N S

S N

N S

BTQBT (18a)

R

S

N

N

S

R

S

S R

R= H (18i) S = Benzo (18j) R

X= C (18g) = N (18h)

Chalcogenophenes (19a-19f) Se

X

Se Se

Se

X X= S (19b) X= Se (19c) X= Te (19d)

(19a)

Se R

R Se

R=H (19e) =Ph(19f)

FIGURE 3.1.6 Tetrathiafulvalene and chalcogenophene derivatives.

Because charge transfer is dependent on intermolecular orbital overlap, the addition of atoms such as tellurium or selenium may have a positive impact on the transistor performance. Despite this possibility, there are very few examples of this approach in the literature. OTFTs with quaterselenophene (19a) as an active layer have been reported, demonstrating a mobility of 3.6 × 10–3 cm2V–1s–1.[135] A new class of chalcogen-containing OFET materials, 2,6-diphenylbenzo[1,2-b:4,5-b′]dichalgenophenes (19b-19d) have also been reported.[136] All three chalcogen compounds perform as good hole-transporting materials. In particular, the benzodiselenophene derivative (19c) demonstrates a relatively high OFET performance with a mobility of 0.17 cm2V–1s–1 and an on/off ratio greater than 105. Recently, single crystal FETs of 19c were fabricated and a mobility of 1.5 cm2V–1s–1 was observed.[137] [1]Benzoselenopheno[3,2-b][1]benzoselenophene (19e) and its 2,7diphenyl derivative (19f) were synthesized and tested in OTFTs. Derivative 19f exhibited excellent p-channel field-effect properties with hole mobilities as high as 0.3 cm2V–1s–1 and current on/off ratios of ~106.[138]



173

3.1.2.1.3 Star-Shaped Oligomers Although the major emphasis so far has been on linear-structured organic semiconductors, there are a few of examples of star-shaped tritopic molecules and their use as the active material in OFETs. Compared to their linear counterparts, these materials show somewhat lower mobilities, but they possess better solubility and improved film-forming properties. Aromatic amines are frequently used as xerographic materials and incorporated in OLEDs,[139] while reports of their use as the semiconductor material in OFETs are few. Only very recently triarylamine polymers have been used in OFETs.[140,141] Star-shaped molecules with triphenylamine core and carbazole (20a, 20c-20d) or fluorene (20b) side groups have all been characterized in OFETs (Figure 3.1.7).[142] Devices containing these materials were fabricated using a drop casting technique and showed mobilities on the order of 10–4 cm2V–1s–1 with good stability under ambient conditions. The star-shaped oligothiophene, 21a gave a mobility of 2 × 10–4 cm2V–1s–1 with an on/off ratio of 102 (Figure 3.1.7).[143] A series of star-shaped oligothiophenefunctionalized truxene derivatives (21b-21d) have also been tested in OFETs.[144,145] The large truxene core is expected to extend the π-delocalized system compared to star-shaped polythiophene (21a). Thin films were obtained by spin coating from chloroform solutions, yielding mobilities as high as 1.03 × 10–3 cm2V–1s–1 for 21b with an on/off ratio of 103. Recently, two novel ‘hybrid’ systems (20e-f) consisting of a triphenylamine core carrying π-conjugated bithienyl branches have been synthesized, with 20e displaying a hole mobility of 0.011 cm2V–1s–1.[146] 3.1.2.1.4 Oligoaryls p-Quaterphenyl (p-4P, 22a), p-quinquephenyl (p-5P, 22b), and p-sexiphenyl (p-6P, 22c) are examples of other commercially available semiconductor materials (Figure 3.1.8). OTFTs fabricated from thermally evaporated films of these oligophenyls gave field-effect mobilities ranging from 0.01 cm2V–1s–1 for p-4P (22a) to 0.07 cm2V–1s–1 for p-6P (22c) with on/off ratios from 105 to 106.[147] Rod-like oligo(arylacetylene)s (23a-e) have also been studied for OTFT applications.[148-149] The arylacetylene oligomers contain electron-donating para-substituted trimethyl amine units and differ by the alkyl substituents on the silyl group. Thin films of 23a and 23b showed mobilities of 0.3 cm2V–1s–1 and 4.3 × 10–4 cm2V–1s–1, respectively.[148] The superior device performance of 23a was attributed to its better structural ordering, and hence more effective π-orbital interaction and charge transport, due to the less bulky trimethylsilyl group. Incorporation of acetylene groups was achieved in an oligothiophene molecule using microwave assisted reactions to produce 23c.[150] The thermally evaporated thin films were highly crystalline with large grains, similar to the unsubstituted pentathiophene (27b). However, the electronic performance was shown to be far inferior, with a maximum mobility of 8 × 10–4 cm2V–1s–1 at a substrate deposition temperature of 140°C. Alkylated oligo(arylacetylene)s 23d and 23e showed ambipolar charge transport with hole mobilties reaching up to 0.02 cm2V–1s–1 and maxium electron mobilities around 2 × 10–3 cm2V–1s–1.[151,149]


174


Triarylamines (20a-20f) R N (20a)

R=

R=

(20c)

N

N (20b) N R

R N

S

S

R

R N

n-C6 H13

(20e) S R=

R= S

(20d)

R N

O

O

N

N

n-C6 H 13

S (20f)

S

R

S R

N R

Oligothiophenes (21a-21d) H n

C10H 21

S

S

S C6H 13

C 6H 13 S H n

C 6H13

C6 H13

C 6H 13

C6 H13

n= 1 (21b) = 2 (21c) = 3 (21d)

S

S H n

S S

S (21a)

C 10 H21

S

S S C 10 H 21

FIGURE 3.1.7 Star-shaped aryl amines and thiophenes.

A modified oligo(p-phenylenevinylene) derivative, 1,4-bis[4-(4-octylphenyl)styrl]-benzene(24a) has been reported[152], showing a field-effect mobility of 0.12 cm2V–1s–1 with an on/off ratio of 106 (Figure 3.1.8). Recently, OTFTs have also been fabricated using thermally evaporated films of two oligo-p-phenylenevinylenes, 1,4-bis(4-methylstyryl)benzene (24b) and 1,4-bis(2-methylstyryl)benzene (24c). Compared with the performance of 24c, 24b showed a higher mobility of 0.13 cm2V–1s–1, attributed to better film continuity and molecular linearity compared to 24c.[153] Vinylene groups have also been incorporated in the thiophene backbone



175

Oligo(arylacetylene) (23a-23e)

Oligo(aryl) (22a-22c)

Me3N

SiR 3

n n= 2 (4P) (22a) n= 3 (5P) (22b) n= 4 (6P) (22c)

R= CH 3 (23a) R= iPr (23b) R

S

S

S

S

S

R= H (23c) =n-C 6H 13 (23d) =n-C 10 H21 (23e)

R

(23c)

Oligo(arylvinylene) (24a-24g) C 8H 17 C 8H 17 (24a)

(24b)

S S

(24c)

R2

R2

S S (24d)

S

R1

R2 S

S R2

R2 S

S R2

R2

R1 R2

R 1, R2 = H (24e) R 1= C6 H13, R2= H (24f) R 1= H, R2= C 6H 13 (24g)

FIGURE 3.1.8 Oligoaryls used in thin–film transistor devices.

to give a series of unsubstituted (24d-24e) and substituted (24f-24g) oligothienylenevinylenes.[154,155] It has been shown that these oligomers have the largest effective conjugation and smallest band gap among all classes of π-conjugated oligomers with comparable chain extension.[156] OTFTs fabricated with thin films of unsubstituted thiophene-vinylene oligomers, 24d and 24e, showed field-effect mobilities of 0.012 cm2V–1s–1 and 0.0014 cm2V–1s–1, respectively. In order to study the effect of the substitution pattern of oligothienylenevinylene core on TFT characteristics, oligomers with alkyl groups either at the terminal (24f) or central thiophene rings (24g) were synthesized. Thin films derived from 24g exhibited very low mobilities, in the range of 1.1 × 10–6 cm2V–1s–1 while compound 24f bearing hexyl chains at the end π-positions showed mobilities 0.055 cm2V–1s–1. The low mobility for 24g compared to 24f is attributed to the steric effects of the hexyl substitutents at the 3and 4-positions of the thiophene rings, which maintain intermolecular separation of the conjugated units. 3.1.2.1.5 Macrocyclics Phthalocyanines (Pc) are another of the first reported families of organic semiconductors (Figure 3.1.9).[157,158] Their structures consists of a molecular cage, into which various metals can be introduced. Highly ordered, thermally evaporated thin films of copper phthalocyanine (CuPc, 25a) are p-channel semiconductors, with a


176


Phthalocyanines and Porphyrins (25a-25r)

t

Bu R

N N

M= = = = = = =

N M

N

N N

N N

N

N N

N N

Cu (25a) Sn (25b) H 2 (25c) Zn (25d) Fe (25e) Pt (25f) Ni (25g)

N N

N N

t Bu C8H 17 O O C8H 17 N N N N C8H 1 7O N N N N C8H17O

N

M

C8H17 O O N

C8 H1 7O

N N N

N N N

N O

O

OC8 H1 7 OC 8H 17

O O

C8H17 O C8H17

O O

M= Tb (25k) = Lu (25l)

O

O O O

O

O

N NN N N N N

N N N N N N

O O O

O O M= Eu (25m) = Ho (25n) = Lu (25o)

O O O O O

M

O O O O

N

OC8H17 OC 8H17

C8H 17 O C OH 8 17

M

O O O

N O O

N

R= O(CH2) 10OH; M= Cu (25h) = O(CH2)10OH; M= none (25i) = NH 2; M= none (25j)

N

t Bu

C8 H1 7 C8 H1 7O

N M

O

O O

O O

N O O O O O

N N

Pt N

N

NH N

(25p)

N

HN

NH

HN H2SO 4

HN

NH

HN NH

(25q)

N (25r)

Other Coordination Compounds (26a-26c) H N (26a)

2+

H N

RH 2N

Ni N H

Pt

RH2N

N H

Cl

NH 2R Pt

Cl (26b)

S

S Ni

S

S

RH 2N RH2 N Cl Cl

FIGURE 3.1.9 Macrocyclic phthalocyanines and porphyrins.

NH 2R

Cl

2-

Cl Pt

NH2R

2+

NH2R Pt

Cl Cl

2-

R= (26c)



177

mobility strongly dependent on the thin-film morphology. The best OTFT performance obtained was a mobility of 0.02 cm2V–1s–1, with an on/off ratio of 105. [159] A single–crystal CuPc-based OFET was recently reported, with mobility as high as 1.0 cm2V–1s–1.[160] In a subsequent study, different metallophthalocyanines (25b25g) were tested in OFETs.[11] The nature of metal ions and substrate temperature were found to be especially crucial in achieving high field-effect mobilities. Besides Cu, Zn-Pc (25d) gave the highest mobility of 2.8 × 10–3 cm2V–1s–1 at a 200°C substrate temperature. Recently it was observed that the charge mobility of phthalocyanines can be increased up to 0.11 cm2V–1s–1 for OFETs having source/drain electrodes sandwiched between two layers of CuPc and CoPc.[161] Most devices based on Pc are fabricated by vacuum deposition, i.e. thermal evaporation. However, recent solution-based Langmuir-Blodgett (LB) techniques have also been used. LB-films of 25h-25l gave poor (10–6 cm2V–1s–1 for 25j) to moderate (1.7 × 10–3 cm2V–1s–1 for 25l) mobilities (Figure 3.1.9).[162-164] More recently, high mobilities of 0.24–0.60 cm2V–1s–1 have been observed for LB films of amphiphilic tris(phthalocyaninato) rare earth triple-decker complexes, 25m25o.[165] OFET characteristics have also been reported for porphyrin derivatives, 25p-q with mobilities of 2.2 × 10–4 cm2V–1s–1 and 0.012 cm2V–1s–1, respectively.[166-168] LB-prepared thin films of Cyclo[8]pyrrole (25r), an extended porphyrin-like molecule, also exhibited a high mobility of 0.68 cm2V–1s–1 , but with a low on/off ratio.[169] The coordination compounds bis(4-methyl-1,2-phenylenediamino) nickel (26a) and bis(dithiobenzyl) nickel (26b) have also been used in OFETs. The mobilities reach up to 0.013 cm2V–1s–1.[170] A platinum based polymer with poly-Pt chain structure (26c) was synthesized and fabricated as FETs under ambient conditions[171]. Mobilities of 10–3–10–4 cm2V–1s–1 were obtained with 10–103 on-off ratios. Interestingly, immersion of FETs in water at 90°C for 12 hours did not deteriorate important device characteristics.

3.1.2.2 THIOPHENE-BASED OLIGOMERS For many years oligothiophenes and their alkyl-substituents have been among the most intensely investigated organic semiconductors because of the synthetic versatility of the thiophene heterocycle.[172] The first printed organic transistor was fabricated by Garnier et al. using sexithiophene as the active semiconductor layer.[12,173,174] Recently, they have also been used to demonstrate fast integrated circuits.[16] Because of their ease in functionalization, oligothiophenes provide the opportunity to study systematic variations in molecular structure by controlling the number of repeat units in the conjugated backbone and/or varying the length and functionality of alkyl substituents. One of the most widely studied oligothiophenes is α-6T (27c) (Figure 3.1.10); since the first report, its hole mobility has improved from 10–4 cm2V–1s–1 to greater than 0.07 cm2V–1s–1. [93,105,175] It has also been shown that the orientation and morphology of thermally evaporated thin films of α-6T depend strongly on the substrate temperature.[54] Increasing substrate temperature from –204°C to 280°C improves the thin-film morphology, with mobility increasing from 0.006 cm2V–1s–1


178


Unsubstituted Thiophenes (27a-27d) S

S

S

S

S

S

S

S

4T (27a) S

S

S

S

S

S

5T (27b) S

S

S

S

S

S

S

6T (27c)

S

S

8T (27d)

Alkyl-substituted Thiophenes (28a-28x) S

R

S

S

S

R

R

S

R= C 6H 13 (28a) = C 10H 21 (28b) = cyclohexyl (28c) =C 12 H 25 (28d)

S

S

S

R

S

S

C6 H13

S (28e)

S

R

S

S

S

S

S

S

S

S

S

S

S

S

R= CH3 (28f) = C 10H 21 (28g) = C 6 H13 (28h) R

C6 H 13

S

S

R

R= C 2H 5 (28i) = C 6 H13 (28j) = C 10H 21 (28k) = C 12H 25 (28l) = C 18H 37 (28m)

R

R = C6 H13 (28n) R

SC 4H 9

R

ButMe2 Si

S

S S

S

S

S

S

n

=3 =4

S C 4H 9 S

(28s)

n= 2 R=C5 H 11 (28o) C12 H25

R= C 5 H11 (28p) R= C 6 H13 (28q)

S

S

S

C 6H 13

S

S

S

S

S

S

S

S

S

S

S S

C 12H 25

C 6H 13

SiMe2Bu t

C 12H 25

S

R=C6 H13 (28r)

S

n n= 4 (28t) = 5 (28u) = 6 (28v)

S

(28w)

S S C 12 H25

C6 H13

C6 H13

(28x)

FIGURE 3.1.10 Unsubstituted and alkyl-substituted oligothiophenes.

to 0.025 cm2V–1s–1. Single crystals of α-6T give mobilities as high as 0.1 cm2V–1s–1.[176] Carrier mobilities near 0.2 cm2V–1s–1 have been reported for αoctithiophene (8T, 27d) OFETs with films deposited at a substrate temperatures of 150°C and higher.[105] With oligothiophenes, as well as all other conjugated oligomers, an increase in conjugation length brings about a decrease in solubility. This characteristic makes purification and deposition more difficult using traditional wet techniques, such as column chromatography and spin-coating. One way to increase the solubility of oligothiophenes in organic solvents is to introduce alkyl substituents at the β-position



179

of the thiophene ring. Furthermore, it has been suggested that an introduction of end-substituents at the 2-position of thiophene oligomers provide stability against oxidation and polymerization while minimizing steric interactions that can prohibit co-planarity of the thiophene rings. Properly designed end-substituents may also improve crystallinity, encourage two-dimensional growth, and promote dense intermolecular packing within each layer, thereby improving charge transport. Molecular engineering of thiophene oligomers by alkyl substitution results in a remarkable improvement of the structural organization of molecules at the mesoscopic level thereby creating highly ordered thin films, resulting in enhanced mobilities.[54] Evidence of strong correlation between orientation and electrical transport properties has been given by the comparison between films of unsubstituted linear oligothiophenes (27a-27d) and oligothiophenes substituted by linear alkyl groups at terminal α-positions (28a-28b, 28d-28m) (Figure 3.1.10).[54], [177] The self-organization induced by the alkyl groups results in a preferred orientation of the substituted oligomers, with their long molecular axes roughly perpendicular to the substrate plane, while the unsubstituted molecules 27a-27d may align either parallel or perpendicular. The change in orientation induced by the alkyl chains leads to an order of magnitude increase in the field-effect mobility to 0.13 cm2V–1s–1 for 28i,[54,178] from 0.01 cm2V–1s–1 for the unsubstituted 6T (27c)[175,179,180]. For the substituted oligomers, the positions of the alkyl chains on the conjugated backbone are also important.[54] In the case of the β-substituted 28s, very low mobilities of ~10–7 cm2V–1s–1 were observed for thin films that exhibited poor structural order. Furthermore, a systematic study was performed where both the conjugated units (from four to six thiophene units) in the backbone and length of the side chains (zero, two, six and ten alkyl units) were varied.[105] It was found that OTFT performance depends critically on the length of the side chains, but is relatively insensitive to the length of the conjugated backbone. This has previously been observed for unsubstituted oligothiophenes ranging from 4 to 8 thiophene units (27a-27d).[181] For side chain longer than hexyl, reduced mobility values were observed.[105] In addition to linear alkyl chains, cyclohexyl groups have also been attached as side substituents (Figure 3.1.10).[182] Although slightly bulkier in nature, the mobility for thermally evaporated thin films of dicyclohexyl-quarterthiophene, 28c (0.038 cm2V–1s–1)[182] was found to be higher than that of the dihexyl-substituted counterpart, 28a (0.02 cm2V–1s–1) with the same device dimensions.[183] Owing to the good solubility of 28c, solution deposition was also performed by drop casting, and mobilities as high as 0.06 cm2V–1s–1 were observed. Bis-silylated versions of oligothiophenes (28t-28v) have also been tested in OFETs, exhibiting mobilities on the order of 10–5 cm2V–1s–1.[184] Recently, a novel structure of , α′dihexylpentathiophene-based swivel cruciform (28x) was synthesized and tested in an OFET. It exhibited a field-effect mobility of 0.012 cm2V–1s–1 and on/off ratio greater than 105. [185] The incorporation of ether linkages into oligothiophene side-chains has also been investigated.[186] The solubility of these oligomers increased by a factor of two compared to the corresponding alkyl substituted oligomers of the same conjugation length. The hole mobility of quarterthiophene derivative 29a was 0.003


180


cm2V–1s–1 while that of the sexithiophene derivatives 29b and 29c were 0.033 cm2V–1s–1 and 0.009 cm2V–1s–1, respectively (Figure 3.1.11). Although the solubility of the new α-6Ts was much improved, the incorporation of an oxygen atom appears to diminish the mobility in solution-cast films relative to alkyl substituted thiophene oligomer films. Other polar species, such as phosphonate groups, have also been incorporated as end-substituents (DBP-α-6T, 29d) (Figure 3.1.11) and used in OTFTs with films fabricated from spin-casting demonstating a mobility of 4.9 × 10–3 cm2V–1s–1 and an on/off ratio of 104.[187] OTFTs have also been fabricated with thin films of oligoethyleneoxide functionalized sexithiophene (29e). Mobilities on the order of 10–4 cm2V–1s–1 were observed.[188] Sexithiophenes bearing amide or ester groups (29f-l) were also synthesized and tested in OTFTs. The oligomer bearing the ester functional group separated from the sexithiophene core by an ethylene spacer showed a hole field-effect mobility as high as 0.012 cm2V–1s–1. [189] A series of fluorene-thiophene oligomers with an identical core but having different end functionalities ranging from unsubstituted (30e), n-hexyl (30b) to cyclohexyl (30f) groups has been investigated (Figure 3.1.11).[182,190] Cyclic voltammetry showed that these derivatives have a lower HOMO energy than the corresponding oligothiophene of the same length (–4.95 eV for DH6T, 28i and –5.36 eV for DHFTTF, 30b). Thermally evaporated, thin films of the unsubstituted (30e), n-hexyl (30b) and cyclohexyl (30f) derivatives gave mobilities of 0.08 cm2V–1s–1, 0.14 cm2V–1s–1 and 0.17 cm2V–1s–1 respectively. The higher mobility of the fluorene-thiophene derivatives (30a-f) is a result of molecular alignment perpendicular to the substrate, in contrast to the equivalent fluorenone-substituted oligothiophene (30g), which was shown to grow parallel with the substrate, resulting in a significantly reduced mobility of less than 10–7 cm2V–1s–1.[191] The crystallinity and growth mode of the oligothiophene-fluorene substituted thin-films was also highly dependent on the substrate deposition temperature and end-group substitution. Another series of fluorene-thiophene oligomers with fluorene as the core (30i-q) have been synthesized and tested in OTFTs. A room-temperature mobility of up to 1.0 × 10–3 cm2V–1s–1 was obtained for 30l.[192] In an effort to decrease reactive degradation of the fluorene, 5,5′-bis(9,9′-dialkylfluorene-2-yl)2-2′-bithiophenes (30q-s) were synthesized. The field-effect mobilities range from 10–5 cm2V–1s–1 for an amorphous film of 30s to 3 × 10–3 cm2V–1s–1 for a polycrystalline film of 30r. The high mobility remains constant after 3 months at ambient conditions, which demonstrates the high environmental stability of this class of materials.[193] Thiophene-phenyl cooligomers with various end substitutions have been synthesized by several groups (Figure 3.1.11). Unsubstituted phenylene capped oligothiophenes showed mobilities of 1.4 × 10–3 cm2V–1s–1 (31c), 3 × 10–3 cm2V–1s–1 (31h) and 0.033 cm2V–1s–1 (31l). Biphenyl-capped oligomers (31d, 31j, 31n) were synthesized, where the thin film morphology and corresponding electrical performance was found to be dependent on the conjugation length.[194] Mobility values of 7 × 10–3 cm2V–1s–1, 0.17 cm2V–1s–1, and 0.055 cm2V–1s–1 were observed for 31d, 31j, and 31n, respectively, which correlated to the degree of crystallinity observed in the films. The same group had earlier reported a mobility of 0.66 cm2V–1s–1 for



Polar Substituents (29a-29l) S

R

S

S

R

S

S

R

R=

O

(29e)

O

n

O C4 H 9

N

O (29i)

C4 H 9

O

S

S

S

S

R

S

R= (CH 2) 3OC4 H9 (29b) = (CH2 )3 OC 8H 17 (29c) = (CH2 )4 PO(OEt) 2 (29d) O O O C10H 21 C10H 21 C4 H 9 N N N (29f) H (29g) H (29h) O O C10H 21 O (29l) (29k) O O

O

R 6

S

R

R= (CH2 )3 OC 4H 9 (29a)

(29j)

Thiophene-Fluorene Oligomers (30a-30v)

S S S

R n= 1, R= C 6H 13 (30a) n= 2, R= C 6H 13 (30b) n= 3, R= C 6H 13 (30c)

(30g)

O

OR O

O

(30j)

O

n=8 R=

(30n) O

R=

S

(30p)

S

O S

R1

S

R2

S

R= n-C2 H5 (30r) n-C4 H9 (30s) sec-C 4H 9 (30t)

R

OR

(30q)

R

R

(30o)

S

O

O

O

S

R

S

S

(30k)

O

C 8H 17 C 8H 17 RO

O

O

O

(30m)

O

(30l)

S

(30h) S

O

O

S

S

(30i) O

O

n= 4, R= C 6H 13 (30d) n= 2, R= H (30e) n= 2, R= cyclohexyl (30f) H2n+1 Cn C nH 2n+1

RO

n=3 R=

R

n

R1

S

R 1= C 6 H13, R2 = flourene (30u) R 1= C 6 H13, R2 = flourenone (30v)

Thiophene-phenyl Oligomers (31a-31z) C6 H 13

C 6H 13 n

S

n=1 (31a) n=2 (31b)

R

=

R

S

n=4

C 6H 13 (31k)

R= phenyl (31l) = tolyl (31m) = biphenyl (31n)

S

S

(31f)

= S

C 6H 13 (31o)

=

S

S

C 6H 13 (31e)

=

n

S

R

R

=

F

C 10 H21 (31f')

S

(31g)

S

S S

R

R= H (31q) C 6H 13 (31r) C 10 H21(31s)

R

S

S

S R=C 6H 13 (31w)

R=H (31v)

31t C6 H13 R

C10H 21 (31u) F F

F

R

F S

C 6H 13 S

S

F

S

S S

F =

R= tolyl (31p) R

R

S

R= phenyl (31c) = biphenyl (31d)

R= phenyl (31h) = tolyl (31i) = biphenyl (31j)

n=3

S

R

n

S

S S

R

S S

S

R=H (31x)

F

S S F

F

F

F

(31z)

R=C 6H 13 (31y)

Asymmetric End-capped Thiophenes (32a-32b) R1

S S

R 1 = H, R2 =C6 H13 (32a)

S S

R2

S S

S S (32b)

FIGURE 3.1.11 Thiophene and co-oligothiophene derivatives.

C6 H13

181


182


crystals of 31d that were epitaxially grown on a KCl (001) substrate and transferred to SiO2.[195] The number of thiophene units have been varied from one to four (31a, 31e, 31k, 31o) with highest mobility realized from 31e (0.09 cm2V–1s–1) and 31o (0.09 cm2V–1s–1).[196] Tolyl-substituted oligothiophenes containing three (31i), four (31m), and five (31p) thiophene rings have also been prepared [197] Although different thin-film morphologies were observed for each derivative, nearly identical mobilities were reported for the three (0.03 cm2V–1s–1 for 31i, 0.03 cm2V–1s–1 for 31m, and 0.02 cm2V–1s–1 for 31p). At elevated substrate deposition temperatures, layer-by-layer growth was observed for each molecule, but the molecular tilt angle at the substrate was substantially different. Furthermore, differences in morphology were observed based on odd or even number of thiophene units. By AFM, films with an even number of thiophene units showed faceted grains with sharply defined edges, while those with an odd number of thiophenes displayed elliptical domains with rough step edges. Oligothiophenes have also been substituted at the central position of the core with phenyl based units.[198] The co-oligomers 31r, 31y exhibited improved thermal stability over oligothiophenes of the same length. The oligomers have a broader band gap and lower HOMO energies than hexyl-substituted sexithiophene (28j). Depending on the structure of the central units, the hole mobility from the top-contact device configuration was 0.042 cm2V–1s–1 (31r), 0.049 cm2V–1s–1 (31y). A series of decyl-substituted (31s, 31u, 31f′) were synthesized yielding a hole mobility as high as 0.3 cm2V–1s–1, 0.08 cm2V–1s–1 and 0.4 cm2V–1s–1 .[199,200] Top-contact devices showed a hole mobility of 0.007 cm2V–1s–1 with an on/off ratio of 104.[115] Fluorene substituted phenyl-thiophene cooligmers (31g, 31z) are p-type semiconductors with mobilities of 0.01 and 4 × 10–5 cm2V–1s–1.[201] Unsubstituted and hexyl substituted oligomers with alternating phenylene and thiophene on the backbone have also been reported. Mobilities of 0.02 cm2V–1s–1 and 0.054 cm2V–1s–1 were obtained in the case of thermally deposited thin films of 31v and 31w.[115,196] An asymmetric hexyl-substituted quarterthiophene derivative (32a) has been reported[202] and compared to unsubstituted and di-hexyl substituted analogues (Figure 3.1.11). The authors showed an increase in mobility from 2 × 10–3 cm2V–1s–1 with no side chains (27a), to 0.02 cm2V–1s–1 with one hexyl chain (32a), and 0.06 cm2V–1s–1 for the symmetric hexyl substituted quaterthiophene (28a). Fused benzothiophene end-groups (32b) also improved the mobility (0.01 cm2V–1s–1) compared to unsubstituted 4T (27a), but to a lesser extent.[202] Oligomers composed of alternating thiophene and furan rings, and those having alkyl substituents at both ends of the molecules have been reported (Figure 3.1.12). A mobility of 0.042 cm2V–1s–1 was achieved for a thienyl-furan oligomer composed of five heterocycles and having hexyl groups at the terminal rings (33c)[203] A dithiophene-substituted indenoflurene compound (34c) has also been synthesized; a relatively high hole field-effect mobility of 0.012 cm2V–1s–1 was reported.[204]



183

Furan-Thiophene Oligomers (33a-33f) R

S

O

S

O

S

R

R

S

R

3

2

R= H (33a) = C 2 H5 (33b) = C 6 H13 (33c)

R= H (33d) = C 2H 5 (33e) = C 6H 13 (33f)

Other Thiophene Oligomers (34a-34c)

S

S S

S

(34a)

S C 6 H13

S

S

S

C6 H 13

(34b) S

S (34c)

FIGURE 3.1.12 Core-substituted thiophenes.

3.1.2.3 POLYMERS Polymers are attractive materials for solution processable organic semiconductors because of their high solubility and good film forming properties. One of the first solution-processed organic semiconductors used for field-effect transistors was poly(thiophene) (35a) (Figure 3.1.13). Since then various substituents have been incorporated on the polymer backbone, e.g. to impart functionality, increase solubility, or induce self-assembly. Among them, poly(3-hexylthiophene) (P3HT) is the most widely studied. The 3-alkyl substituents can be incorporated into a polymer chain with two types of arrangements, either head-to-tail (HT) or head-to-head (HH). A polymer with a mixture of HH and HT linkages in P3HT is referred to as regiorandom (35c), while one with only HT linkages is referred to as regioregular (35b). It has been found that highly regioregular P3HT (35b) self-orients into a wellordered lamellar structure with an edge-on orientation of the thiophene rings relative to the substrate.[205] P3HT with a high regioregularity (>95% HT linkages) adopt lamellae with an edge-on orientation giving mobilities of 0.05–0.2 cm2V–1s–1 and on/off ratios close to 106.[206-208] Polymers with a low degree of regioregularity


184


Thiophene-based Polymers (35a-35z'') R

R

S S head-to-tail (HT)

S

n

S R

RR

(35a)

S

S S head-to-head (HH) C12H 25 O

C 6H 13

C 6H 13

C 6H 13

S S S S poly(3-hexylthiophene) P3HT

S

C 6 H13

R= C4 H 9 (35d) = C 8H 17 (35e) = C 10 H21 (35f) n = C 12 H25 (35g)

C6 H 13

Regioregular P3HT (35b) Regiorandom P3HT (35c) O

HO O

O S

S

S

S O O

(35k)

n

S

n

(35h)

C 12 H25

O O

S (35j)

n

(35i)

C6 H13

F

S S

S

C 6 H13

F

S

S

S

F F (35m)

C6 H13

(35n)

C 8H 17

n

S

C 6 H13

S

S (35o)

S S

S

S

R R = C 8H 17 (35p) = C10H 21 (35q) = C12H 25 (35r)

S

3,3"PTT-8 (35s) S

Bu S m

S

S

n

n 3',4'PTT-8 (35t) C12H 25

S

S

C 12 H 25

PQT-12 (35u)

S 2

S S

S

Bu S

C8 H 17

S

S

n

n

C 8 H17

C 8H 17

C 8H 17 C8 H17

S

n C 12H 25

n

(35l)

S

S

n

C12H 25 C12H 25

S

Bu Si Bu

n

N

S

R

C6 H13

Bu Si Bu

m n

m= 2 (35v) = 3 (35w)

n Bu

S

m= 2 (35x) = 3 (35y)

Bu S

S 2

S m

S 2

2

n

R S S

S S

n

R = C 10 H21 (35z) = C12 H25 (35z') =C 14 H 29 (35z'')

R

FIGURE 3.1.13 Thiophene-based polymers.

adopt lamellae having face-on orientation giving mobilities of 10–4 cm2V–1s–1. A variety of solvents have been used for solution deposition of P3HT.[159] The mobility was found to vary by two orders of magnitude, with 1,2,4-trichlorobenzene giving the highest mobility.[208] LB prepared films of regioregular P3HT have shown a mobility of 0.02 cm2V–1s–1.[209] A comparison of poly(3-alkylthiophene)s (P3ATs) with side chains ranging from butyl (35d) to dodecyl (35g) showed a non-monotonic dependence of field-effect mobility on alkyl chain length (Figure 3.1.13).[11,210] The average mobility varied



185

from 1.2 × 10–3 cm2V–1s–1 in butyl (35d) to 2.4 × 10–5 cm2V–1s–1 for dodecyl (35g)substituted P3AT. In addition to the variation of alkyl side-chain length, modification of the chemical nature of the side chains in regioregular substituted polythiophenes have also been investigated. Regioregular polymers having chiral (35h), branched side chains (35i) and terminal carboxylate (35j) units were synthesized to study the nature of the side chains’ influence on morphology and OFET performance.[211] The carboxylic acid side chains were incorporated in an attempt to use hydrogen bonding to aid selforientation at polar surfaces, e.g. SiO2. The polymers are soluble in common organic solvents; however, solution-cast thin-films were non-uniform and essentially amorphous (except for 35h) giving mobilities of 1 × 10–3 cm2V–1s–1, 2.8 × 10–5 cm2V–1s–1 and 2.9 × 10–4 cm2V–1s–1 for 35h, 35i and 35j, respectively. The lower field-effect mobilities of these polymers compared to unbranched regioregular P3HT (35b) is believed to be due to the longer π–π stacking distance. The introduction of branches (methyl substitution), bulky side chains (oxazoline) or carboxylic acid units prohibit the formation of crystalline films through steric interactions. Polythiophenes containing electron withdrawing alkyl carboxylates (35k-35l) have also been synthesized.[212] OTFTs fabricated with the regioregular polythiophene (35k) gave mobilities of up to 0.07 cm2V–1s–1 compared to 0.0063 cm2V–1s–1 for the regiorandom polythiophene (35l).[212] Another electron-withdrawing moiety, tetrafluorobenzene, has been substituted on the polyalkylthiophene backbone (35m). Spin-cast films of 35m from 0.5 wt % chloroform solution showed a mobility of up to 2 × 10–3 cm2V–1s–1.[213] Recently there has been a thrust in design and synthesis of novel solution processable polymeric semiconductor other than P3HT. It has been reasoned that the electrical performance of P3HT is lower in air due to oxygen, which acts as a dopant, and its stability can be improved by increasing its ionization potential (IP).[214,215] The latter is partially dependent on effective π-conjugation length of the polymer backbone, which can be controlled either sterically — by reducing πoverlap between adjacent thiophene rings, or electronically — by introducing less conjugated unit in the backbone. Researchers at Xerox, Canada have demonstrated that proper control of the extended π- — system of regioregular polythiophenes enables solution fabrication of stable OFETs in ambient conditions.[216-218] They have synthesized regioregular alkyl-substituted polythiophenes with the alkyl sidechains strategically placed along the polythiophene backbone, to tune the π-conjugation by controlling both the torsional barriers and rotational freedom of thienylene moieties. Polymers 3,3′PTT-8 (35s)[217] and 3′,4′PTT-8 (35t)[216] differ from each other only in the regiochemistry of the alkyl side-chains. OTFTs were fabricated by spin coating 0.5–1 wt % chlorobenzene solutions of polythiophenes. The two regioisomers 35s and 35t showed mobilities of 0.03 cm2V–1s–1 and 0.01 cm2V–1s–1, respectively. The “spaced-out” arrangement of alkyl side-chains along the polymer backbone in 35s resulted in 3D-lamellar π-stacking and extensive intermolecular side-chain interactions, leading to a higher mobility compared to 35t, where only two-dimensional face-to-face π-stacks exist in thin films. PQT-12 (35u),[218] however, showed a much higher mobility of 0.14 cm2V–1s–1 and improved air stability, probably due to the presence of more unsubstituted thienylene moieties that possess


186


substantial rotational freedom, to us reducing the π-conjugation to some extent. This reduced π-conjugation imparts sufficient oxidative stability to the system, which is reflected in the minimal decrease in mobility and on/off ratio after being stored under ambient conditions for one month. In contrast, regioregular P3HT undergoes a drastic degradation in performance for devices stored under identical conditions. Polymers 35n and 35o have sterically inhibiting alkyl groups on neighboring thiophene rings, while polymers 35p-35r have an aromatic heterocycle thieno[2,3b]thiophene in the backbone that has a central cross-conjugated double bond designed to break the conjugation.[214,215] All the polymers exhibited useful molecular weights and low polydispersities. Thin-film transistors were fabricated either by spin–coating or drop–casting chloroform, dichlorobenzene or xylene solutions. Polymers 35o and 35n gave mobilities of 0.03 cm2V–1s–1 and 6 × 10–4 cm2V–1s–1, whereas mobilities of 0.15 cm2V–1s–1 and 0.12 cm2V–1s–1 were obtained for liquid crystalline polymers 35q and 35r, respectively with on/off ratios of 105.[214,215] The mobility of 35q has been further improved to 0.6 cm2V–1s–1 which is among the highest for solution-processed polymers.[219] The devices were found to be stable to ambient light and air. Monosilanylene-oligothienlene alternating polymers (35v-35y) contain silicon linkages which were anticipitated to provide solubility and elevate the HOMO energy level via its electron donating properties. However, due to the amorphous nature of their thin films, low mobilities in the order of 10–5 cm2V–1s–1 were observed.[220] McCulloch et al reported new semiconducting liquid-crystalline thieno[3,2b]thiophene polymers (35z-z′′′). Good transitor stability under static storage and operation in a low-humidity air environment was demonstrated, with charge-carrier field-effect mobilities of 0.2–0.6cm2V–1s–1 achieved under nitrogen.[221] Recently, fluorene-based co-polymers have attracted attention as promising materials for polymer OTFTs (Figure 3.1.14). A new class of co-polymeric semiconductor having fluorene and thiophene blocks, the AB-type co-polymer, poly(9,9′dioctyl-fluorene-co-bithiophene) (F8T2, 36a) have promising semiconducting characteristics.[222,223] OTFTs fabricated by spin-coating 0.5 wt % chloroform solution of 36a gave mobilities of 3 × 10–3–5 × 10–3 cm2V–1s–1. However, the mobility of 36a was dramatically enhanced to 0.01–0.02 cm2V–1s–1 by depositing the polymer onto a mechanically rubbed, substrate followed by high temperature annealing.[7,222] The chains of 36a were aligned parallel to the rubbing direction of the underlying substrate to facilitate charge transport resulting in a higher mobility. To increase the planarity and rigidity of the polymer backbone, the thieno[3,2-b]thiophene group was introduced into a fluorene-based alternating co-polymer, F8TT (36b).[224] Thin films of 36b gave a mobility of up to 1.1 × 10–3 cm2V–1s–1. Conjugated polymers consisting of dioctyl fluorene units and low band gap donor-acceptor-donor (D-AD) units (36c) have also been tested in OTFTs.[225] The D-A-D segment includes two electron-donating thiophene rings combined with an electron-withdrawing thiadiazolo-quinoxaline unit. Thin-film transistors of 36c gave a field-effect mobility of 3 × 10–3 cm2V–1s–1 and an on/off ratio of 104. Charge-transfer type co-polymers comprised of electron-donating and electronwithdrawing units are considered to be potentially useful as ambipolar semiconductors. With this in mind, co-polymers of thiophene (-electron donor) and thiazole or



187

Poly (9, 9'-dioctyl-flourene-co-bithiophene) (36a–36c) S S

S

S

C8H17 n

SN N

n

C8H17

S

C8H17 F8T2 (36a)

C8H17

Thiophene-thiazole/thiadiazole co-polymer (37a–37b) C9H19 N N N S S S S n (37b) (37a)

n

N N

F8TT (36b)

C8H17

S

C8H17 (36c)

Cyclopentadithiophene based polymers (38a–38c) R1

R1

S

S

R1

R1 R2

n

n

R1 = H, R2 = C8H17 (38a) R1, R2 = C8H17 (38b)

Poly(alkylidene fluorene) (39a–39c) R R

R1

S n S R = C8H17 (38c)

Poly(p-phenylene vinylene) (PPV) (40a–30e) OR1

OR2 O

n

R = C6H13 (39a) = C8H17 (39b) = C10H21 (39c)

R1O

n

R3O

n

R1 = CH3, R2 = C18H37, R2 = H (40a) R1 = CH3, R2,3 = 3,7-dimethyloctyl (40b) R1,2 = C11H23, R3 = C18H37 (40c)

n RO R = CH3 (40d) = 3,7-dimethyloctyl (40e)

Phenoxazine-based polymers (41a–41f) C6H13 CH 6

13

N

N O (41a)

O

n

C6H13 C H 6 13 Polytriarylamines (42a)

x

(41b) (X = 0.25) (41c) (X = 0.50) (41d) (X = 0.75)

1–x n

C6H13 N

N O (41a)

C6H13 (41e)

S

N n

O (41f)

(42a)

n

n

FIGURE 3.1.14 Thiophene copolymers and non-thiophene-based polymers.

thiadiazole (-electron acceptor) were synthesized and tested in OTFTs. The thiadiazole co-polymer (37b) (Figure 3.1.14) indeed showed ambipolar behavior[226], demonstrating an electron mobility of about 5 × 10–3 cm2V–1s–1 and a hole mobility of 3 × 10–4 cm2V–1s–1. The thiazole containing co-polymer (37a), however, showed only a hole mobility of 2.5 × 10–3 cm2V–1s–1.[227] Polycyclopentadithiophenes, a class of solution processable, thiophene based analogues of the polyfluorenes, have also been investigated as charge transport


188


materials in OTFTs.[228-231] Alternating copolymers, 38a and 38b, of cyclopentadithiophene and dioctylfluorene were prepared by Suzuki cross coupling. Thin films of 38c, obtained from spin-casting from a 0.8 wt% chloroform solution gave a mobility of 7.9 × 10–5 cm2V–1s–1, whereas polymers 38a and 38b show field-effect mobilities of 10–6–10–7 cm2V–1s–1. Since the carbon at the 9-position of fluorene-based molecules is sp3-hybridized and hence limits the π-conjugation, polyalkylidene fluorenes (39a-39c) have been synthesized. Polyalkylidene fluorenes (39a-39c) have a sp2-hybridized carbon at the 9-position which beside increasing the effective conjugation also permits the alkyl chains to adopt a coplanar conformation relative to the polymer backbone.[232] OTFTs fabricated by spin-casting 0.4–1 wt% polymer solutions in chloroform gave a mobility of up to 2 × 10–3 cm2V–1s–1 with an on/off ratio of 106. Different alkoxy-substituted poly(p-phenylenevinylene)s (PPVs) have been examined to relate performance to the substitution on the phenyl ring (symmetric: 40b, 40e or asymmetric: 40a, 40c, 40d) and nature of side chain, either linear or branched.[233-235] The PPV polymers were soluble in common organic solvents and OTFTs were fabricated by spin-casting a 0.3 wt % solution of PPV in chlorobenzene. Unsymmetrical alkoxy-substituted PPVs gave mobilities on the order of 10–4 cm2V–1s–1, compared to 0.01 cm2V–1s–1 for symmetrically substituted polymers. Polymers containing phenoxazine (41a-41e)[236] or phenothiazine (41f)[237] moieties (Figure 3.1.14) have recently attracted much research interest because of their unique electro-optical properties and their potential in diverse applications including light-emitting diodes and chemical luminescence. The rigid and planar structure of phenoxazine ring and its low ionization potential also make it a good building block for the design of semiconductors for OTFT applications. Phenoxazine-based conjugated polymers (41a-e) have been synthesized and tested as pchannel organic semiconductors.[236] Thin films of the polymers were made either by spin-casting or drop-casting a 0.2 wt% toluene or chloroform solution. OTFTs based on poly-phenoxazine (41a) showed a mobility of 0.9 × 10–5 cm2V–1s–1 and copolymers 41c and 41e showed mobilities of 3 × 10–4 cm2V–1s–1 and 6 × 10–4 cm2V–1s–1, respectively. The sulfur counterpart to 41a, poly-phenothiazine (41f) has also been reported, with a similar field-effect mobility (0.8 × 10–4 cm2V–1s–1).[237] Polytriarylamines (42a) are another class of highly stable polymeric semiconductors.[140,141] They can be handled in air and OTFTs are stable, with mobilities of 10–3–10–2 cm2V–1s–1 under ambient conditions. Moreover, several members of this family of compounds are amorphous, simplifying their characterization.

3.1.2.4 SOLUTION PROCESSABLE SEMICONDUCTORS: THE “PRECURSOR METHOD” Fully integrated solution-processed organic devices are of interest for large area rollto-roll manufacturing. To achieve this goal, extensive research efforts have been placed on functionalizing organic semiconductors in order to impart solubility. But most chemical modifications interrupt the natural π-stacking tendency of the molecules, which can inhibit charge transport that relies heavily on π-orbital overlap. In spite of a great deal of success in solution-processable thin-film devices with soluble



189

oligomeric and polymeric semiconductors, their low solubility still remains an inherent problem. Therefore, they usually require elevated solvent temperature to render solubility and the resultant films obtained by spin-casting are non-uniform in thickness and morphology. In general, spun-cast soluble semiconductors have lower performance than their vacuum-deposited counterparts. To mitigate these disadvantages, an alternative approach known as the “precursor route” has been developed. The idea is to synthesize a precursor molecule that is soluble in organic solvents but can be converted to the insoluble counterpart by either thermal or chemical treatment after spin-casting a thin film. 3.1.2.4.1 Precusor Polymers and Small Molecules The first semiconductor that was prepared using the precursor method was polyacetylene (43a)(Figure 3.1.15)[238] using the Durham route.[239] The scheme takes advantage of the properties of an intermediate precursor polymer that is non-conjugated and readily soluble in common solvents. Upon heating this precursor polymer to 80–100°C in vacuum for 12 hours, it undergoes an elimination reaction, losing hexafluoro-o-xylene as a volatile fraction and leaving polyacetylene in a fully condensed form.[239-241] OFETs with polyacetylene fabricated by this method have shown a mobility of 10–4 cm2V–1s–1. A thiophene based polymer poly(2,5-thienylenevinylene) (PTV, 43b) has been synthesized using the precursor method.[242] The dimethylformamide solution of the PTV precursor polymer was spun-cast and then heated to 200°C for 5 min under a nitrogen stream containing a small amount of HCl gas. During the heat treatment, HCl effectively acts as a catalyst for the conversion of the precursor polymer to the PTV.[243] Mobilities as high as 0.22 cm2V–1s–1 were reported. The same polymer has been used for fabrication of integrated circuits,[244] with reported mobilities of 10–4–10–3 cm2V–1s–1. The Durham route has also been used to prepare semiconducting tetrabenzoporphyrin 43c.[245] The precursor-porphyrin derivative consists of four ethylene bridged units which imparts solubility, and hence can be purified by column chromatography. Quantitative conversion to 43c is achieved throught heating the soluble precursor between 150–200°C, which results in the elimination of four ethylene molecules. OTFTs fabricated with this method yield a mobility of 0.017 cm2V–1s–1 and an on/off ratio of 105. The preparation of a soluble precursor of pentacene (3a) that can be spun-cast and then heated to form pentacene has been reported.[246-248] The pentacene precursor (Figure 3.1.15) is soluble in dichloromethane and forms a continuous, amorphous film when spun onto transistor substrates. A simple thermal treatment converts the precursor to pentacene via the elimination of tetrachlorobenzene. Thin films of pentacene formed by this technique demonstrated field-effect mobilities ranging from 0.1–0.2 cm2V–1s–1 for films converted at 200°C. Recently a novel approach for a high yield synthesis of another soluble pentacene precursor was demonstrated.[249,250] The synthesis involves a Lewis acid-catalyzed Diels-Alder reaction of pentacene and N-sulfinylacetamide.[251-253] OTFTs fabricated by spin-casting a chloroform solution of the precursor on substrate followed


190

Organic Field-Effect Transistors Precursor Polymers and Small Molecules (43a-e) F3 C

CF3

F3 C n

n soluble polymer precursor

CF3

HCl

+

S

O

N H

n

PTV (43b)

N H N

N

N -4 C 2 H4

H N

soluble ethylenebridged precursor

S

soluble precursor of PTV

polyacetylene (43a)

N

200o C

n

H N tetrabenzoporphyrin (43c)

Cl Cl

Cl

Cl

soluble pentacene precursor

Cl

pentacene (3a)

Cl

Cl Cl

200 oC O 120-200 oC

N S

O

soluble pentacene precursor R2 O R1 S O

S S Si

O H3 C C N S O O

S

R2

soluble precursor 150-200o C S

R1

O

S

S

S

Si S

soluble precursor

+

S

S

S S

S

pentacene (3a)

cat. CH 3ReO 3 CHCl3 , reflux

S S

(43d)

TBAF, pyridine n-butanol, 20 oC

S S

S S (43e)

FIGURE 3.1.15 “Precursor polymers.” Semiconducting polymers can be made by either thermal or chemical treatment of preformed solution-cast films.

by annealing (200°C/15 min or 130°C/25 min) under nitrogen atmosphere converted the adduct to pentacene (3a). The field-effect mobility of this film was as high as 0.89 cm2V–1s–1, ranking among the highest mobilities reported for an OTFTs fabricated by solution processing. The precursor approach has also be applied to thiophene oligomers.[254,255] The synthesis of symmetrically α, ω-ester substituted sexithiophene (precursor to 43d) containing thermally removable solubilizing groups and its incorporation into OTFTs was reported. Bulky and highly soluble ester end groups allowed the oligomers to be solution-cast into thin films at room temperature. A subsequent heating



191

Polydiacetylene (44a-c) R1 C C R2

n

R 1 = -(CH 2 )11CH3 , R 2= -(CH 2)COOH (44a) R 1 ,R 2= -(CH 2) 4OCONHPh (44b) R 1 ,R 2= -(CH 2) 4OCONH(CH2 )2 CH3 ( 44c)

FIGURE 3.1.16 Small molecule precursors.

cycle induced thermolytic removal of the ester solubilizing groups, affording 43d, while permitting the molecules to re-adopt the preferred π-stacking orientation that resulted in high charge mobility. OTFTs showed an increase in hole mobility from 10–5 cm2V–1s–1 with on/off ratios of ~100 before thermolysis to >0.1 cm2V–1s–1 with on/off ratios >105 after thermolysis. Recently, an acene fused-thiophene hybrid, p-channel semiconductor using a soluble precursor in which the solubility originates from trimethylsilyl groups was reported.[256] The elimination of two trimethylsilyl groups from the soluble precursor is accomplished through treatment with tetrabutylammonium fluoride in pyridine to give 43e. OTFTs made by thermally evaporated 43e showed mobility of 10–2 cm2V–1s–1. 3.1.2.4.2 Polydiacetylene Polydiacetylene (44a-c) (Figure 3.1.16) derivatives have been synthesized and tested in OTFTs. A hole mobility of 1.3 × 10–3cm2/vs and an on/off ratio of 104 were reported for 44a, the polydiacetylene prepared from 10, 12-pentacosadiynoic acid.[257]

3.1.3 N-CHANNEL ORGANIC SEMICONDUCTORS N-channel organic semiconductors are important class of materials, due to their necessity in the fabrication of bipolar transistors and complementary logic circuits. However, most of the literature to-date has focused on the design of p-channel semiconductors. Molecules such as unsubstituted oligothiophenes or pentacene are more conducive to the injection of holes than electrons since their ionization potential (~5 eV) matches reasonably well with the work function of typical metals used as source and drain electrodes (such as Au or Ag). Lower work function metals such as Al or Ca, would better facilitate electron injection, but are difficult to work with due to oxidative instability and their tendency to form charge transfer complexes between the organic semiconductor and metal.[40] In order to lower the electron injection barrier between the LUMO level of the organic with respect to the Fermi level of the metal electrodes, strong electron-withdrawing groups are often added to the outer rings of molecules through synthetic design. This has been done successfully with several semiconductor core systems. These groups increase the electron affinity and stabilize the anionic form of the molecule, allowing for the


192


possibility of efficient electron injection and transport. Synthetic design still provides a formidable challenge because of the inherent instability of organic anions, which are susceptible to atmospheric oxidants such as water and oxygen. Recent spectroscopic evidence has also shown that in organic semiconductors without sufficiently large electron affinities (greater than 3.85 eV), mobile electrons can be trapped by silanol groups at the dielectric interface of oxides used in typical transistor test configurations.[258] Nonetheless, recent progress has been made in this direction, with electron mobilities of greater than 0.1 cm2V–1s–1 demonstrated in several different families of materials. The synthetic design of n-channel organic semiconductors remains an intimate balance, which involves optimizing electron transport while limiting the degradation of electronic properties.

3.1.3.1 FULLERENES

AND

FULLERENE DERIVATIVES

Fullerenes (C60 and C60/C70) and fullerene derivatives (Figure 3.1.17) were some of the first n-channel materials studied. C60 (45a) has a solid state electron affinity of ~4.5 eV, and its nearly spherical shape provides isotropic carrier transport in thin films, which is different than the anisotropic behavior observed with linear oligomers (such as pentacene or oligothiophenes). In 1993, n-channel behavior in fullerenes, a mixture of C60 and C70 in a 9:1 ratio, was observed with mobilities as high as 5 × 10–4 cm2V–1s–1.[259] Later, polycrystalline films of C60 with a grain size on the order of 6 nm showed mobilities as high as 0.08 cm2V–1s–1.[260] When the substrate was pretreated with tetrakis(dimethylamino)ethylene prior to C60 deposition, the mobility increased to 0.3 cm2V–1s–1 at the cost of a diminished on/off ratio. The device performance degraded upon exposure to ambient conditions, but could be restored by annealing the films under high vacuum. Since then, mobilities as high as 0.56 cm2V–1s–1 have been reported for C60 films fabricated by molecular beam deposition without breaking vacuum.[261] Soluble methanofullerene derivatives have also been used in n-channel OTFTs. [6,6]-Phenyl C61-butyric acid methyl ester (PCBM, 45b) is a commercially available, solution processable derivative, whose LUMO level lies at approximately 3.7 eV with respect to vacuum. This barrier to electron injection (~1.4 eV with respect to Au) is significantly reduced due to the formation of a strong interface dipole layer at the metal/semiconductor interface.[244] A similar phenomenon has been observed for pristine C60 on Au and Ag.[262] Mobilities as high as 0.02–0.1 cm2V–1s–1 have been obtained for films of 45b measured in ambient conditions, which strongly depend on the work-function of the source and drain electrodes. Similarly high mobilities of 0.2 cm2V–1s–1 were recently demonstrated on polymer gate dielectrics for devices measured under an Argon atmosphere.[263] Soluble fullerene dendron (45b′) has been fabricated as an n-channel OTFT, with field-effect mobility reaching 1.7 × 10–3 cm2V–1s–1 at 300K.[264]

3.1.3.2 PHTHALOCYANINES Unsubstituted phthalocyanine derivatives show p-channel mobility as described in the previous section. However, when substituted with electron withdrawing groups



193

Fullerenes and Fullerene Derivatives (45a-45b') O

O

OMe

O PCBM, (45b)

C 60 , (45a)

N

O O

NH

OMe O

N

HN

NH

O

N

OMe

O

(45b')

O

OMe

Phthalocyanines (45c-45k) R

R

R

R

R

R R

N

R

N N N

R

N

R

N M

R

Cu

N

N N

N

N

N

N

R R

R R= R= R= R= R=

N

N

N

R R

R

R

R

R

R = SO3 Na (45h)

F, M = Cu (45c) F, M = Zn (45d) F, M = Co (45e) F, M = Fe (45f) Cl, Me = Fe (45g) R

R N N

N Cu

N

N N

N

N

R=

R + N

Cl-

N

M

N

N

R

N N

(45i)

N N

N

M = Lu (45j) = Tm (45k)

FIGURE 3.1.17 Fullerenes and phthalocyanines.

at their periphery, such as hexadecafluoro-subsitituted copper phthalocyanine (45c) (Figure 3.1.17), these materials showed n-channel field-effect mobilities as high as 0.03 cm2V–1s–1.[265] Not only do these devices show good electronic performance, but they operate under ambient conditions with excellent stability. Other metals centers and halogenated substitutents have been studied; the highest mobilites were obtained in each case from devices fabricated with a substrate temperature of 125°C or greater. The best mobilities for the other derivatives were 1.2 × 10–3 cm2V–1s–1


194


(45d), 4.5 × 10–5 cm2V–1s–1 (45e), 2.1 × 10–3 cm2V–1s–1 (45f), and 2.7 × 10–5 cm2V–1s–1 (45g). The difference in mobility values was attributed to to morphological differences in the crystallinity and grain size of the films. Ambipolar mobility has also been reported from water soluble copper phthalocyanines substituted with sulfonic acid (45h) and methyl pyridinium groups (45i). Unlike typical transistors with a linear and saturation regions, the drain-source current of these devices increases non-linearly with increasing VDS at a given gate bias. Furthermore, the entire set of I-V curves shifts up with repeated scans while the on/off ratio decreases. The authors attributed this behavior to a mechanism that involved ion-modulated electrical conduction. N-channel mobilities as high as 3 × 10–4 cm2V–1s–1 were observed.[266] Bisphthalocyanines based on rare earth metals lutenium (45j) and thulium (45k) were studied as early as 1990.[267] These compounds had a field-effect mobility of ~10–4 cm2V–1s–1 under vacuum, but the electronic properties degraded rapidly upon exposure to ambient conditions.

3.1.3.3 NAPHTHALENE DIIMIDE DERIVATIVES Naphthalene tetracarboxylic dianhydride (NTCDA, 46a) and its family of imide derivatives are easily synthesized from commercially available starting materials (Figure 3.1.18). When measured in vacuum, the unsubstituted 46a showed n-channel mobility of ~10–4 cm2V–1s–1 when deposited at a substrate temperature of 25°C and 3 × 10–3 cm2V–1s–1 when it was increased to 55°C. Upon exposure to air, the mobility decreased by two orders of magnitude. The morphology of both films showed small crystalline grains with an average size of 200 nm, but the films deposited at a substrate temperature of 55°C had grains with better connectivity, accounting for the higher mobility observed.[268] Alkyl substituted versions of naphthalene tetracarboxylic diimide (NTCDI, 46b) showed mobilities of 0.16 cm2V–1s–1 (R = octyl, 46c), 0.01 cm2/V–1s–1 (R = dodecyl, 46d), and 5 × 10–3 cm2V–1s–1 (R = octadecyl, 46e) when measured under vacuum, but no observable electron mobility when measured in ambient. The incorporation of fluoroalkyl groups on the side chains stabilizes NTCDI thin films and allows for n-channel operation under ambient conditions. Mobilities as high as 0.03 cm2/V–1s–1 were obtained for 46f and 0.06 cm2V–1s–1 for 46g with on/off ratios on the order of 105.[269] An even higher mobility (0.12 cm2V–1s-1) was obtained with a substituted benzyl derivative that contained a CF3 group in the para position (46h). The striking differences in the electron mobility measured under ambient atmosphere between NTCDI derivatives containing fluorine groups in the sidechains and those without was investigated in a later publication.[270] From electrochemistry measurements, the electron withdrawing character of the fluorinated substituents (R = CH2C7F15, 46g) shifts the reduction potential by about 0.14 V in comparison to the nonfluorinated (R= octyl, 46c) NTCDI. The difference was negligible between the octyl-substituted NTCDA and the derivative with a terminal fluoromethyl group (46h). None of the fluorinated compounds had reduction potentials lower than that of NTCDA or C60, neither of which can operate in air. From this data, it was concluded that solid state packing effects must play a major role in explaining the



195

Naphthalene Diimides (46b–46l)

O

O

O

O

N

O

O

N

O

O

O

N

O

O

N

O

O N C3F7

O

O O NT CDA, (46a)

C3F7

R

H

R R = C8H17 (46c) = C12H25 (46d) = C18H37 (46e)

H NTCDI, (46b)

N O

O N

O

O

N

O

O

O N C7F15

O

(46f)

(46g) R1

O F3C

C7F15

N

CF3

O

N

O

O

N

O

(46h)

O

R2 R1 = CH2 (CH2)7CH3 R2 = NH2 (46i) = CH2 (CF2)6CF3 = NH2 (46j) = NH2 (46k) = CH2 (CF2)2CF3 = N(CH3)2 (46l) = (CH2)7CH3

FIGURE 3.1.18 Naphthalene diimide derivatives.

stability, but the origin of this behavior is still unknown. One possible explanation involves the close-packing nature of fluorined side chains, which may provide a kinetic barrier to oxidation. Asymmetric NTCDIs with alkyl and fluoroalkyl sidechains have also been synthesized (46i-46l). The highest mobility obtained with an alkyl version (46i) was 3 × 10–4 cm2V–1s–1 under vacuum and 2 × 10–4 cm2V–1s–1 with a fluoroalkyl derivative (46j).[271]

3.1.3.4 PERYLENE DIIMIDE DERIVATIVES Perylene (47a) (Figure 3.1.19) has been employed as the active-layer in OTFTs, exhibiting low p-channel mobilities and no n-channel mobility, even though theoretical calculation predicts n-channel behavior.[272] Dianhydride and diimide versions of perylene have also been studied in detail by many groups. The earliest report of n-channel mobility of perylene tetracarboxylic dianhydride (PTCDA, 47b), ~10–4 cm2V–1s–1, was obtained under vacuum or in moisture-free air.[273] The substituted perylene tetracarboxylic diimides (PTCDI, 47c-47j) can be easily synthesized in similar fashion to the naphthalene versions of the same molecule. The energy levels of alkyl substituted diimides are similar to unsubstituted versions (3.4 eV and 5.4 eV for electrons and holes referenced to vacuum level).[274,275] The single-crystal structure of 47c reveals π-stacking in a triclinic lattice.[276] The highest reported n-channel mobility for a perylene derivative to date was found from an octylsubstituted PTCDI (47d), reaching a value of 1.3 cm2V–1s–1 under vacuum, when


196

Organic Field-Effect Transistors Perylene Diimide Derivatives (47a-47k) O

O

O

O

O

O PTCDA, (47b)

(47a)

NC O

O

R N

N R

O

O

O

O

R N

N R

O

O CN

R = C5 H11 (47c) = C 8 H17 ( 47d) = C 12H 25 (47e) = C 13H 27 (47f) = C 6 H5 (47g) R= (47h)

R=

(47i)

R = n-CH2 C 3F7 R= n-C 8H 17

(47j)

(47k)

FIGURE 3.1.19 Perylene diimides.

corrected for contact resistance.[277] In the same work, PTCDI substituted with pentyl (47c) and dodecyl (47e) sidechains showed a mobility of 0.06 and 0.5 cm2V–1s–1, respectively. Thin films of these derivatives adopt a molecular packing comparable to the single-crystal structure of 47d. Recently, n-channel mobility of 0.6 cm2V–1s–1 was reported with 47f and complementary inverters were demonstrated (using pentacene as the p-channel material) that showed record gain.[275] Derivatives with phenyl[278] (47g) and cyclohexyl substitutents[182] (47h) have also been reported with mobilities of 1.5 × 10–5 cm2V–1s–1 and 1.9 × 10–4 cm2V–1s–1 respectively. Core-substituted perylene diimides have also been functionalized at the two bay positions (position 1 and 7) with electron withdrawing cyano groups (47i and 47j).[279] Both compounds showed good solubility in organic solvents and have a lower LUMO than 47d, suggesting improved n-carrier stability. Single crystals of 47j were grown by sublimation, and the crystal structure analysis revealed a polycyclic core that is slightly twisted from steric repulsion caused by the cyano groups. Furthermore, the molecules were shown to pack with a face-to-face, slipped πstacked structure with a minimum interplanar spacing of 3.4 Å. Electronic measurements were performed on compounds that were a 1:1 mixture of isomers cyanated at the 1,7- and 1,6-positions of the perylene core. Mobilities as high as 0.1 cm2V–1s–1



197

and 0.64 cm2V–1s–1 were found for 47i and 47j, respectively, at optimized substrate deposition temperatures. Owing to the good solubility, solution processed films of 47j show mobilities of 10–3–10–4 cm2V–1s–1 in the bottom contact device configuration with alkane thiol treated gold electrodes. 47i was also cast from solution onto unmodified substrates, where mobilities of 10–3–10–5 cm2V–1s–1 were observed. These two materials have excellent air stability, with on/off ratios as high as 105 when measured under ambient conditions. N, N′-bis(n-octyl)-dicyanoperylene-3,4:9,10bis-(dicarbpximide) (47k) was also used in an n-type OTFT, with a reported mobility and on/off ratio of 0.14 cm2V–1s–1 and 1.2 × 10–3. [280]

3.1.3.5 QUINOID SYSTEMS Tetracyanoquinodomethane (48a) and 11,11′,12,12′-tetracyanonaptho-2,6,quinodimethane (TCNQ, 48b) are commercially available and n-channel behavior has been observed for both compounds (Figure 3.1.20).[281] Devices fabricated using 48b display a higher mobility (10–3 cm2V–1s–1) than 48a (10–5 cm2V–1s–1), and have better air-stability than NTCDA, although both 48a and 48b have low on/off ratios due to unintentional doping. A new class of dicyanopyrazinoquinoxaline derivatives were also synthesized and tested for their electronic performance. The mobilities obtained for all the compounds (48c-48h) were very low, ranging from 10–6–10–8 cm2V–1s–1, but some of the compounds had a very unique crystal packing arrangement that the authors referred to as a molecular tape structure arising from the C-H…N intermolecular interactions.[282] Dicyanomethylene groups have been incorporated onto the periphery of a terthiophene backbone, stabilizing the all-planar quinoid form of the molecule, with n-Butyl groups were added to the central thiophene moiety for improved solubility (48i).[283]. The bulk crystal structure indicates that the molecules form a face-toface, slipped π-stack structure containing π-dimers that pack in alternating columns. Thin films deposited at elevated temperatures showed an n-channel mobility as high as 0.2 cm2V–1s–1 in high vacuum.[284] Additionally, ambipolar behavior was observed in some cases when the substrate temperature for deposition was greater than 136°C.

3.1.3.6 THIOPHENE BASED N-CHANNEL OLIGOMERS As with perylene cores, perfluoroalkyl substitution has been performed on oligothiophenes (Figure 3.1.21).[39,285,286] In this case, the substitution switches the charge carrier from holes to electrons. With oligomers 49a-49c, the HOMO-LUMO gap is nearly identical among the perfluoroalkyl substituted and the alkyl substituted versions of the same molecules, but in each case, the LUMO is approximately 0.5 eV lower for the perfluoroalkyl versions. The shift of energy levels is quite different from the naphthalene and perylene cores substituted with fluorinated sidechains, in which the perfluoroalkyl chain was not in direct communication with the π-system because of an alkyl spacer. The influence of the halogenated substitutions is also different than the halogenated phthalocyanine system, as the perfluoroalkyl chains


198


Quinoid systems (48a–48i) NC

CN

NC

CN

NC CN

NC

CN

(48a)

TCNQ, (48b)

R2

R1 N

N

CN

N N CN R4 R1 = R2 = R3 = R4 = H (48c) R3

N

N

CN

NC

N

N

N

N

CN

NC

N

N

(48g)

NC NC

(48h) S S

S R

R

CN CN

R1 = R4 = H (48d) R2 = R3 = CH3 R1 = R4 = H (48e) R2 = R3 = OCH3 R1 = R4 = OCH3 (48f) R2 = R3 = H

R = C4H9 (48i)

FIGURE 3.1.20 Quinoid derivatives.

only exhibit a strong σ-inductive electron-withdrawing effect, where as halogens directly attached to the phthalocyanine system exhibit both σ-inductive and resonance donating effects. Electron mobilities of 0.059 cm2V–1s–1, 0.026 cm2V–1s–1, and 1 × 10–3 cm2V–1s–1 were observed for 49a, 49b, and 49c, respectively, with on/off ratios ranging from 104–105. The same type of substitution has been performed with mixed phenylene-thiophene oligomers .[287] All of the oligomers (49d-49g) showed clear n-channel performance, with 49e demonstrating the highest mobility for the series at 0.074 cm2V–1s–1 and an on/off ratio of 6 × 106 when deposited at a substrate temperature of 110°C. The fabrication of transistors based on carbonyl functionalized quarterthiophenes (DHCO-4T (49h) and DFCO-4T (49i)) has been reported.[288] The carbonyl functionality acts as a strong electron withdrawing group that allows the possibility of further chemical functionalization. High n-channel mobilities of 0.1 cm2V–1s–1 for 49h and 0.6 cm2V–1s–1 for 49i were obtained under high-vacuum measurement conditions. Surprisingly, 49h also exhibited p-channel mobilities as high as 0.01 cm2V–1s–1. The authors claim that 49i also exhibits ambipolar behavior only after doping with I2. A more electron deficient 49j and the dioxolane protected 4T core (49k) were also reported to show n-channel mobility an order of magnitude lower than the carbonyl derivatives 49h and 49i.[288] Perfluorophenyl endgroups were attached to a quarterthiophene core (49l) and an n-channel mobility of 0.08 cm2V–1s–1 with an on/off ratio of 105 was reported.[289] However, perfluorophenyl dithiophene (49q) and tertthiophene (49p) exhibited poor mobilities compared with 49l, particularly for 49q. The crystal structure of 49q reveals that only three rings are fully conjugated, making the electronic structure of this molecule closer to lower-mobility terthiophene than quaterthiophene derivatives.[201] Phenacyl groups have been incorporated at the periphery of oligothiophene core (49m) to fabricate OTFTs with high electron mobilities.[290] With the substi-



199

Thiophene-based N-Channel Oligomers (49a-49q) S

C6 F13

S

S

C 8F 13

C 8 F13

C 6F13

S

S

x x = 1 (49a) = 1.5 (49b) = 2 (49c)

x

x = 1 (49d) = 2 (49e) = 3 (49f) = 4 (49g) O S

R

S S

R

S O

R = C6 H13 (49h) = C6 F13 (49i)

O O

O

C6 F13

S

S S

C 6 F13

S

S

C 6F 13

S

S

C6 F13

S

O

O

(49j)

O

O

(49k) F

F

F

F

S

S F

F

(49l)

O

S

S

O

S

F

O

F

F

F

F

F

F

S

S

O (49n)

F

F

S

F

F

F

F (49p)

S F

F

F

S

F F

F

F

F

F

S

(49o)

F

F

F

F

F S

F

F O F (49m) F

F

S

F

O

S

S F

F

F

F O

F

F

F F

F

F

O

F

S

S

F

F F

(49q)

F

FIGURE 3.1.21 n-Type thiophene derivatives.

tution of perfluorophenacyl groups, electron mobilities as high as 0.45 cm2V–1s–1and on/off ratios of 108 were observed with devices tested under argon. Under the same argon conditions, these materials were also cast from solution and a mobility of 0.21 cm2V–1s–1 with on/off ratios of 105 was obtained, accounting for the highest values to date for solution processed n-channel devices. The phenacyl substituted quarterthiophene without fluoro-substitution (49n) showed only p-channel characteristics, with the highest mobility value of 0.04 cm2V–1s–1. Suprisingly, the reduction potential


200


of the perfluorophenacyl derivative was slightly more negative than that of the pchannel derivative, as observed from cyclic voltammetry. The authors were able to grow single crystals of both materials and attributed the different electronic properties to subtle molecular conformational differences in the crystal structure. Careful analysis reveals that the dihedral angle between the carbonyl group and the thiophene core was much greater in the case of the phenacyl derivative with H atoms. The greater conjugation sustained in the perfluorophenacyl derivative should enhance stabilization of a negatively charged core in the solid state. These two molecules provide a good example where molecular energy levels are not the determining factor for the type of charge carrier, and show that even small changes in crystal packing can play a major role. An n-channel, perfluorinated version of pentacene (49o) has recently been synthesized.[72] The crystal structure of both pentacene and perfluoropentacene adopt a herringbone geometry and belong to the same symmetry group (p2gg), but the two molecules in the unit cell exhibit different azimuthal angles. The closest proximity of two intermolecular carbon atoms in pentacene is 3.64 Å while short C-C intermolecular distances ranging between 3.22Å and 3.25Å are observed in perfluoropentacene due to the electrostatic interaction between the electropositive pentacene moieties and electronegative fluorine atoms. Devices fabricated with perfluoropentacene at a substrate temperature of 50°C revealed an n-channel mobility of 0.11 cm2V–1s–1 and an on/off ratio of 106 when measured under vacuum. Inverter circuits were also fabricated using pentacene as the p-channel semiconductor.

3.1.3.7 TRIFLUOROMETHYLPHENYL-BASED OLIGOMERS Recently, the trifluoromethylphenyl endgroup has also been used as an electron withdrawing group in linear oligomers based on phenylene-thiophene (50a), thiazolothiazole (50b, 50c), and anthracene derivatives (50d) (Figure 3.1.22).[118,291] Single crystals of 50d reveal a herringbone packing arrangement similar to usual oligothiophenes, but the oligomers containing the thiazolothiazole unit (50b and 50c) have a columnar π-stacking structure. The best transistor performance reported for this family of compounds occurred for a substrate deposition temperature of 50°C. Compound 50c demonstrated an n-channel mobility of 0.30 cm2V–1s–1 and an on/off ratio of 106. 50b did not show any transistor characteristics, while 50c had an electron mobility as high as 0.30 cm2V–1s–1 and an on/off ratio of 106. The phenylthiophene co-oligomer, 50a, showed a mobility of 0.18 cm2V–1s–1. In a later publication the authors observed an electron mobility of 3 × 10–3 cm2V–1s–1 from the phenyl-anthracene cooligomer, 50d.[292] Recently, transistors fabricated based on thiazole oligomers with trifluoromethylphenyl groups demonstrated very high electron mobility.[293] Within the series of molecules (50e-50i), the electron affinity increases with increasing number of thiazole rings. Compound 50e did not show any mobility but 50f exhibited an nchannel mobility as high as 1.83 cm2V–1s–1 when deposited on OTS treated substrates at room temperature. This is the highest reported room temperature mobility to date for n-channel materials. The longer oligomers also showed good mobilities with 2.8



201

Trifluormethylphenyl-based oligomers (50a–50j) F3C

S

S

N S S N

F3C

CF3

(50a)

F3C

S

CF3

(50b) S

N

N

S

CF3 S

F3C

CF3

(50d)

(50c)

F3C

F3C

N S

S N

S N (50g)

S

N S

S N (50i)

N S S N (50f )

F3C

N S S N (50e)

F3C

CF3 F3C

N S

S N

CF3

CF3

N S

S

CF3

(50h) F3C

S

S

CF3

S

S

S

S

CF3

(50j)

FIGURE 3.1.22 n-Type organic semiconductors with trifluoromethyl endgroups.

× 10–3 cm2V–1s–1 for 50g, 0.085 cm2V–1s–1 for 50h, 0.018 cm2V–1s–1 for 50i cm2V–1s–1 and 0.025 cm2V–1s–1 for 50j.

3.1.3.8 POLYMERIC SYSTEMS The electron mobility has been investigated in ladder type polymers poly(benzobisimidazobenzophenanthroline) (BBL, 51a) (Figure 3.1.23).[294,295] This polymer has high thermal stability, a glass transition temperature greater than 500°C, and is insoluble in aprotic organic solvents. Films can be solution cast or spin coated from methanesulfonic acid (MSA) or Lewis acid (AlCl3, GaCl3, FeCl3)/nitromethane mixtures. Electron mobility in the saturation regime was 0.03–0.05 cm2V–1s–1 for films cast from MSA with an on/off ratio of 103. In the linear regime, mobilities as high as 0.1 cm2V–1s–1 were observed, the same order of magnitude as the best hole mobility obtained in conjugated polymer systems. The non-ladder derivative, BBB (51b), which has an identical optical band gap and absorption spectra, has also been investigated. The highest obtainable mobility was 10–6 cm2V–1s–1 when cast under the same conditions. The large differences in mobility were attributed to morphological differences in thin films between the two polymers. The BBL has more efficient π-stacking and greater intermolecular order reasoned from the high degree of crystallinity. X-ray scattering from BBL (51a) thin films exhibit a Bragg (010) d-spacing of 3.3–3.4 Å due to the — intermolecular stacking of the planar chains.[296] Films of BBB (51b) were shown to be completely amorphous.


202


Polymeric Systems (51a–51e) O

O

N

N

N

N

O

O

N

N

BBL, (51a)

n

C8H17 N

S

C8H17

C8H17

n

(51c)

BBB, (51b)

O C8H17

S

n

N

N

n

COONH4

OR NN

N N RO

O

n

N

(51d)

R = 2' ethylhexyl (51e)

FIGURE 3.1.23 n-Type polymeric systems.

Ambipolar behavior has also been observed in BBL and poly(thiophene-3propionic acid, ammonium salt).[297] 51a showed n-channel mobilities of 0.04–0.06 cm2V–1s–1 and 0.02–0.03 cm2V–1s–1 for p-channel operation. 51c showed values of 0.5–0.7 cm2V–1s–1 for n channel and 1.2–1.7 cm2V–1s–1 for p-channel with on/off ratios between 2 and 50 for devices operated in air. A similar mechanism involving ion-modulated electrochemical conduction, as described above for water soluble phthalocyanines was proposed by the authors. Recently, n-channel transistors of poly(9,9-di-n-octylfluorene-alt-benzothiadiazole) (51d) have been reported using thin films fabricated by spin-coating.[298] Electron mobilities ranging from 6 × 10–4–4.8 × 10–3 cm2V–1s–1 were observed with the use of a 50 nm BCB polymer layer on top of the SiO2 dielectric and Ca electrodes. Alternating copolymers containing 1,3,4-heterodiazoles and fluorenes have also been synthesized and show very low n-channel mobilities (2.2 × 10–8 cm2V–1s–1 for compound 51e).[299]

3.1.4 OUTLOOK AND CONCLUSIONS In this chapter we have reviewed advances in organic semiconductor performance, combining newly synthesized, novel materials and modifications of existing promising materials. In addition to efforts driving new material development, many groups have found that the performance of existing materials can be vastly improved with the optimization of their morphology and structural order. In addition to developing new materials, a significant effort has been focused on improving the stability of materials towards air, moisture and light exposure in addition to the electronic performance. Significant achievements have been made in developing pchannel and, to a lesser degree, n-channel semiconductors. The performance of several organic semiconductors, both n- and p-channel, has already surpassed that of amorphous silicon.



203

The growth of OFET technology has been impressive, especially considering the efforts spent in developing and optimizing the materials for the semiconductor layer alone. Transistor performance has improved rapidly as a result of worldwide activity in research groups worldwide. Applications ranging from radio-frequency identification tags to flexible displays have been proposed and are poised for introduction into commercial products. Significant progress has been made toward the correlation of the molecular electronic properties and their arrangement in ordered superstructures to the macroscopic transport characteristics. Though comprehensive descriptions of such relationships remain elusive, the vast array of empirical data and the isolation of key transport parameters offer a promising look toward the engineering of high-performance semiconducting materials. Despite this impressive set of advances, several issues remain to be addressed before commercialization can begin. A better understanding of effects of the dielectric interface, organic semiconductor growth mechanisms and methods to make ohmic contacts will enable better device performances. Long-term research efforts and innovation are needed to provide new organic semiconductors with improved performance, processability, and environmental stability to oxygen and moisture. In the review, we have limited the scope to the semiconductor material, but there has also been great attention paid to the dielectric layer and conductive electrodes necessary to complete the all-organic electronic devices. The successful development and adaptation of these new materials will require increased multidisciplinary partnerships among physicists, chemists, and engineers. As more sophisticated and versatile methods currently developed in the laboratory make their way into the manufacturing environment, we can expect that organic electronic circuits will have a profound impact on future technological advances.

3.1.5 TABLE OF MOBILITIES Table 3.1.1 contains field-effect mobilities and on/off ratios.

TABLE 3.1.1 Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 1a [48] 1b [49] 1c [49] 1d [49] 1e [49] 1f [53] 1g [53]

Deposition methodb

Mobility (cm2V–2s–1)

sc v v v v v v

0.02 0.013 0.13 0.072 0.18 0.063 0.5

ION/IOFF 104 104 104 105 104 8.7 × 105 2.8 × 107 Continued


204


TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 1h [55] 1i [56] 2a [58] 2a [57] 2b [60] 2c [60] 2d [60] 2e [60] 2f [55] 3a [62] 3a [249,250] 3b [67] 3c [46] 3d [61] 3e [68] 3f [69] 3g [69] 3h [69] 3i [71] 3j [71] 4a [79] 4b [79] 4c [79] 5a [34] 5a [81] 5a [83] 5b [81] 6a [85] 6b [85] 6c [85] 6d [85] 6e [86] 6f [87] 7a [88] 7b [88] 7c [88] 7d [88] 8a [89,90] 8b [89,90] 9a [91] 9b [91]

Deposition methodb


ION/IOFF

v v v sc sc sc sc sc v v p v v v v s s v v v v v v sc v s v v v v v s s v v v v v v v v

0.12 1.1 0.1 1.3 1.4 × 10–4 0.3 1.6 N.A.c 0.5 6 0.89 0.3 N.O.d 2.5 0.251 1 × 10–5 1 × 10–5 0.4 0.014 0.045 0.05 N.A. N.A. 15.4 10–3 0.7 10–3 1.4 × 10–4 0.011 0.012 9.5 × 10–4 0.001 0.02 5 × 10–5 0.006 0.001 5 × 10–4 0.038 0.148 0.09 0.15

108 4.4 × 105 106 106 103 102 105 N.A. 108 106 107 6.3 × 103 N.O. 2 × 106 3.15 × 103 N.A. N.A. 106 N.A. N.A. 106 N.A. N.A. 106 N.A. 106 N.A. 3.3 × 102 8.2 × 104 5.7 × 104 8.7 × 102 104 106 103 5 × 103 7 × 102 3 × 103 N.A. N.A. N.A. N.A.



TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 9b [91] 9c [91] 9d [91] 9e [92] 9f [92] 9g [92] 10a [93] 10b [38] 10c [38] 10d [94] 10e [95] 10f [95] 11a [96–98] 11b [99] 11c [100] 11d [101] 11e [101] 11f [101] 12a [102] 12b [102] 13a [103,104] 13b [103,104] 13c [103,104] 13d [103,104] 13e [103,104] 13f [103,104] 13g [106] 13h [106] 14a [107] 14b [107] 14c [108] 14d [108] 15a [109,110] 15b [109,110] 15c [109,110] 15d [110] 15e [111] 15f [111] 15f [111] 15g [110]

Deposition methodb


s v v s s s v v v v v v v v v v v v v v v v v v v v v v v v v v v v v v v s v v

0.02 0.14 0.06 N.A. 1 1 × 10–4 0.04 1.6 × 10–2 10–3 0.15 0.011 0.12 0.05 0.045 0.02 0.42 0.12 0.14 0.06 0.09 3 × 10–4 1 × 10–4 0.006 0.057 0.036 0.11 N.O. 0.0037 7 × 10–5 7 × 10–5 6 × 10–5 4 × 10–5 0.003 1.2 × 10–4 0.12 1 × 10–5 0.01 0.002 0.14 0.001

ION/IOFF N.A. N.A. N.A. N.A. 107 103 105 N.A. N.A. 106 4.0 × 104 1.6 × 105 108 103 106 5 × 106 5 × 105 104 7.3 × 104 7 × 104 30 10 6 × 102 3 × 104 3 × 104 105 N.O. 104 106 106 N.A. N.A. N.A. 103 107 N.A. 105 104 107 105 Continued

205


206


TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 15h [112] 15i [112] 15j [113] 16a [117,118] 16b [117,118] 16c [117,118] 16d [117,118] 16e [117,118] 16f [119] 16g [119] 16h [119] 16i [119] 16j [119] 16k [119] 16l [120] 16m [120] 16n [115] 16o [120] 16p [120] 16q [115] 17a [125–128] 17b [125–128] 17c [125–128] 17d [125–128] 17e [125–128] 17f [125–128] 17g [125–128] 17h [129,130] 17i [129,130] 17j [129,130] 17k [129,130] 18a [131–133] 18b [134] 18c [134] 18d [134] 18e [134] 18f [134] 18g [134] 18h [134] 18i [134] 18j [134]

Deposition methodb


ION/IOFF

v v v v v v v v v v v v v v v v s v v s sc sc sc sc sc sc sc v v v v v v v v v v v v v v

0.30 0.054 1 × 10–6 N.O. 0.02 0.003 0.002 N.O. N.O. N.O. 0.0007 N.O. N.O. 6 × 10–9 0.011 3.5 × 10–4 1 × 10–4 1 × 10–5 2 × 10–5 0.007 0.4 0.015 1.4 0.0018 1.4 × 10–4 0.062 0.0012 0.06 3.3 × 10–5 0.42 0.2 0.2 3 × 10–4 N.O. 2 × 10–6 N.O. 8 × 10–5 N.O. 6 × 10–7 N.O. 2 × 10–4

107 106 103 N.O. 104 104 104 N.O. N.O. N.O. 102 N.O. N.O. 102 104 104 8 × 10 104 104 104 N.A. N.A. N.A. N.A. N.A. N.A. N.A. 104 105 6 × 103 106 108 104 N.O. 50 N.O. 104 N.O. 103 N.O. 104



TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 19a [135] 19b [136] 19c [136] 19c [137] 19d [136] 19e [138] 19f [138] 20a [142] 20b [142] 20c [142] 20d [142] 20e [146] 20f [146] 21a [143] 21b [144,145] 21c [144,145] 21d [144,145] 22a [147] 22b [147] 22c [147] 23a [148] 23b [148] 23c [150] 23d [149] 23e [149] 24a [152] 24b [153] 24c [153] 24d [154] 24d [154] 24e [154] 24f [155] 24g [155] 25a [159] 25a [160] 25b [11] 25c [11] 25d [11] 25e [11] 25f [11]

Deposition methodb


v v v sc v v v s s s s v v s s s s v v v v v v v v v v v v s v v v v sc v v v v v

0.0036 0.081 0.17 1.5 0.0073 N.A. 0.3 1 × 10–4 1 × 10–4 1 × 10–4 3 × 10–4 1.1 × 10–2 10–5 2 × 10–4 0.00103 6.5 × 10–4 2.2 × 10–4 0.01 0.04 0.07 0.3 4.3 × 10–4 8 × 10–4 0.0019 0.0019 0.12 0.13 0.13 0.012 0.0014 0.0014 0.055 1.1 × 10–6 0.02 1.0 0.0034 0.0026 0.0028 6.9 × 10–4 1.5 × 10–4

ION/IOFF N.A. 2 × 103 105 104 2 × 103 N.A. 106 104 105 103 105 2 × 102 N.A. 102 103 102 102 105 106 106 105 N.A. 10 N.A. N.A. 106 5.8 × 105 9.4 × 102 N.A. 1.5 × 102 N.A. N.A. N.A. 4 × 105 104 N.A. N.A. N.A. N.A. N.A. Continued

207


208


TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 25g [11] 25h [162] 25i [162] 25j [164] 25k [163] 25l [163] 25m [165] 25n [165] 25o [165] 25p [167] 25q [168] 25r [169] 26a [170] 26b [170] 26c [171] 27a [39] 27b [150] 27c [105] 27c [176] 27d [105] 28a [300] 28b [105] 28c [182] 28c [182] 28d [177] 28e [39] 28f [150] 28g [105] 28h [39] 28i [105] 28j [105] 28k [105] 28l [186] 28m [186] 28n [181] 28o [177] 28p [177] 28q [39] 28r [39] 28s [301] 28t [184]

Deposition methodb


ION/IOFF

v s s s s s s s s s s s v v s v v v sc v v v v s v s v v v v v v v v v v v v v v v

5.4 × 10–5 4.0 × 10–4 4.8 × 10–5 1 × 10–6 6.4 × 10–4 0.0017 0.60 0.40 0.24 2.2 × 10–4 0.012 0.68 0.013 2.0 × 10–5 10–4–10–3 0.014 0.078 0.07 0.1 0.2 0.23 0.2 0.038 0.06 3.2 × 10–5 4 × 10–4 0.0092 0.5 0.054 1.1 1 0.5 0.016 0.0013 0.03 N.A. 0.0048 0.012 0.064 10–4 8 × 10–6

N.A. N.A. N.A. N.A. 5.5 × 104 100 N.A. N.A. N.A. 104–105 N.A. 8 × 104 10 100 102–103 20 102 102 N.A. N.A. N.A. 105 4 × 105 105 N.A. 4 × 104 103 105 104 104 104 105 N.A. N.A. N.A. N.A. N.A. 4 × 105 7 × 103 5 × 104 103



TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 28u [184] 28v [184] 28w [177] 28x [185] 29a [186] 29a [186] 29b [186] 29b [186] 29c [186] 29c [186] 29d [187] 29e [188] 29f [189] 29g [189] 29h [189] 29i [189] 29 [189] 29k [189] 29l [189] 30a [190] 30b [190] 30c [190] 30d [190] 30e [190] 30f [182] 30g [191] 30h [191] 30i [192] 30j [192] 30k [192] 30l [192] 30m [192] 30n [192] 30o [192] 30p [192] 30q [192] 30r [193] 30s [193] 30t [193] 30u [198]

Deposition methodb


v v v s v s v s v s s s v v v v v v v v v v v v v v v v v v v v v v v v v v v v

3 × 10–4 4 × 10–5 3.9 × 10–5 0.012 0.008 0.003 0.033 0.01 0.009 0.001 0.0049 10–4 N.O. N.O. 6 × 10–6 N.O. 7 × 10–4 6 × 10–4 0.014 0.012 0.14 0.025 0.023 0.08 0.17 N.O. N.O. 1.1 × 10–5 1.2 × 10–5 1.7 × 10–5 1.1 × 10–3 5 × 10–5 5.4 × 10–5 2.9 × 10–5 7.5 × 10–4 5 × 10–4 3 × 10–3 2 × 10–4 10–5 0.01

ION/IOFF 103 103 N.A. 105 N.A. N.A. N.A. N.A. N.A. N.A. 1.6 × 104 103 N.O. N.O. N.O. N.O. 102 10 103 103 2 × 104 102 102 104 8 × 105 N.O. N.O. N.A. N.A. N.A. N.A. N.A. N.A. N.A. N.A. N.A. 106 104 104 104 Continued

209


210


TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 30v [198] 31a [196] 31b [196] 31c [194] 31d [194] 31d [195] 31e [196] 31f [201] 31f′ [200] 31g [201] 31h [115] 31i [197] 31j [194] 31k [196] 31l [115] 31m [197] 31n [194] 31o [196] 31p [197] 31q [198] 31r [198] 31s [199] 31t [115] 31u [200] 31v [115] 31w [196] 31x [201] 31y [198] 31z [201] 32a [202] 32b [202] 33a [203] 33b [203] 33c [203] 33d [203] 33e [203] 33f [203] 34a [202] 34b [198] 34c [204] 35b [208]

Deposition methodb


ION/IOFF

v s s s v sc s v v v v v v s s v v s v s v v v v s s v v v v v v v v v v v v v v s

0.002 0.01 0.018 0.0014 0.0077 0.66 0.09 0.04 0.4 0.005 0.003 0.03 0.17 0.054 0.033 0.03 0.055 0.09 0.02 5 × 10–4 0.042 0.3 5 × 10–4 0.08 0.02 0.054 0.03 0.049 4 × 10–5 0.02 0.011 4.6 × 10–4 0.042 0.014 0.0049 0.012 0.0097 0.011 0.067 0.012 0.12

106 3 × 103 1.5 × 104 104 N.A. 105 4 × 104 106 105 104 4 × 103 107 N.A. 7.7 × 103 1.7 × 103 107 N.A. 3.5 × 103 107 5.9 × 102 104 105 590 103 5 × 104 4 × 104 105 104 104 1.3 × 102 3 × 102 N.A. N.A. N.A. N.A. N.A. N.A. 102 104 N.A. 106



TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 35d [210] 35e [210] 35f [210] 35g [210] 35h [211] 35i [211] 35j [211] 35k [212] 35l [212] 35m [213] 35n [214] 35o [214] 35p [215,219] 35q [215,219] 35r [215,219] 35s [217] 35t [216] 35u [218] 35v [220] 35w [220] 35x [220] 35y [220] 35z [221] 35z′ [221] 35z″ [221] 36a [222,223] 36b [224] 36c [225] 37a [227] 37b [226] 38a [229,230] 38b [229,230] 38c [229,230] 39a [232] 39b [232] 39c [232] 40a [233–235] 40b [233–235] 40c [233–235] 40d [234]

Deposition methodb


s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s s

0.0012 3 × 10–4 8.5 × 10–5 2.4 × 10–5 0.001 2.8 × 10–5 2.9 × 10–4 0.07 0.0063 0.002 6 × 10–4 0.03 N.A. 0.15 0.12 0.03 0.01 0.14 3.4 × 10–5 4.6 × 10–5 6.9 × 10–5 6.7 × 10–5 0.3 0.3 0.63 0.02 0.0011 0.003 0.0025 3.4 × 10–4 5 × 10–6 1 × 10–7 7.9 × 10–5 4 × 10–4 0.001 0.002 1 × 10–4 0.001 0.01 4 × 10–4

ION/IOFF 6.4 × 102 40 1 × 103 3.7 × 102 60 2 5 1 × 105 1 × 104 N.A. 4 × 104 2 × 102 N.A. 1 × 105 1 × 105 1 × 106 1 × 105 2 × 107 103 103 103 103 106 106 107 N.A. 1.4 × 104 104 2 × 102 104 N.A. N.A. N.A. 104 104 106 N.A. N.A. N.A. N.A. Continued

211


212


TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 40e [234] 41a [236] 41b [236] 41c [236] 41d [236] 41e [236] 41f [237] 42a [140,141] 43a [238] 43b [242,243] 43c [245] 43d [254,255] 43e [256] 44a [257] 44b [257] 44c [257] 45a [261] 45b [263] 45b′ [264] 45c [265] 45d [265] 45e [265] 45f [265] 45g [265] 45h [266] 45i [266] 45j [267] 45k [267] 46a [268] 46c [269] 46d [270] 46e [270] 46f [269] 46g [269] 46h [270] 46i [271] 46j [271] 46k [271] 46l [271] 47a [272] 47b [273]

Deposition methodb


ION/IOFF

s s s s s s s s p p p p v p p p v s s v v v v v s s v v v v v v v v v v v v v v v

1 × 10–3 0.9 × 10–5 0.6 × 10–5 3 × 10–4 2 × 10–4 6 × 10–4 8 × 10–5 0.01 1 × 10–4 0.22 0.017 0.1 0.01 0.0013 7.8 × 10–7 N.A. 0.56 0.2 0.0017 0.03 1.2 × 10–3 4.5 × 10–5 2.1 × 10–3 2.7 × 10–5 3 × 10–4 3 × 10–4 10–4 10–4 0.003 0.16 0.01 0.005 0.03 0.06 0.12 3 × 10–4 2 × 10–4 N.O. 10–5 N.O. 10–4

N.A. 102 102 104 104 104 103 N.A. N.A. N.A. 105 105 N.A. 104 N.A. N.A. 108 1,000 N. A. 3 × 105 N.A. N.A. N.A. N.A. N.A. N.A. N.A. N.A. N.A. N.A. 4 × 105 N.A. 105 105 N.A. N.A. N.A. N.A. N.A. N.A. N.A.



TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 47c [277] 47d [277] 47e [277] 47f [275] 47g [278] 47h [182] 47i [279] 47i [279] 47j [279] 47j [279] 47k [280] 48a [281] 48b [281] 48c [282] 48d [282] 48e [282] 48f [282] 48g [282] 48h [282] 48i [284] 49a [39] 49b [39] 49c [39] 49d [287] 49e [287] 49f [287] 49g [287] 49h [288] 49i [288] 49j [288] 49l [289] 49m [290] 49m [290] 49n [290] 49o [72] 49p [201] 49q [201] 50a [291] 50b [291] 50c [291]

Deposition methodb


v v v v v v v s v s v v v v v v v v v v v v v v v v v v v v v v s v v v v v v v

0.06 1.3 0.5 0.6 1.5 × 10–5 1.9 × 10–4 0.1 10–3–10–5 0.64 10–3–10–4 0.14 10–5 0.001 3.6 × 10–6 1.0 × 10–8 2.1 × 10–7 2.5 × 10–7 2.2 × 10–6 9.6 × 10–7 0.2 0.059 0.026 0.001 0.0002 0.074 0.0039 0.0025 0.1 0.6 10–2–10–3 0.08 0.45 0.21 0.043 0.11 0.003 0.00001 0.18 N.O. 0.3

ION/IOFF 106 106 106 107 N.A. 100 105 N.A. 104 N.A. 1.2 × 103 N.A. N.A. 103 10 102 102 102 102 >106 105 105 104 1 × 103 6 × 106 5 × 105 8 × 106 107 107 N.A. 105 108 105 106 106 104 100 3 × 105 N.A. 106 Continued

213


214


TABLE 3.1.1 (Continued) Field-Effect Mobility and On/Off Ratio Found in the Literature for Each Organic Semiconductor Described in the Text Molecule[ref.] 50d [292] 50e [293] 50f [293] 50g [293] 50h [293] 50i [293] 50j [293] 51a [294] 51b [294] 51c [297] 51d [298] 51e [299]

Deposition methodb


ION/IOFF

v v v v v v v s s s s s

3 × 10–3 N.O. 1.83 0.0028 0.085 0.018 0.025 0.1 10–6 0.5–0.7 4.8 × 10–3 2.2 × 10–8

104 N.A. N.A. N.A. N.A. N.A. N.A. 2 × 103 N.A. 2–50 N.A. N.A.

a

Highest field-effect mobility measured for each organic semiconductor as reported in literature. b sc: single crystal; v: vaccum; s: solution; p: precursor. c N.A.: transistor effect “not available.” d N.O.: transistor effect “not observed.”

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3.2

Dielectric Materials: Selection and Design

Ashok Maliakal CONTENTS 3.2.1 Introduction................................................................................................229 3.2.2 Fundamentals and Figures of Merit ..........................................................231 3.2.2.1 Factors Affecting the Dielectric Constant...................................232 3.2.2.2 Thickness.....................................................................................233 3.2.2.3 Dielectric Roughness...................................................................233 3.2.2.4 Film Morphology ........................................................................234 3.2.2.5 Importance of Interface between Dielectric and Semiconductor.............................................................................234 3.2.2.6 Processing....................................................................................234 3.2.2.7 Reliablity .....................................................................................235 3.2.3 Major Classes of Dielectric Materials.......................................................235 3.2.3.1 Inorganic Dielectrics ...................................................................235 3.2.3.2 Polymer Dielectrics .....................................................................237 3.2.4 Alternative Gate Dielectric Strategies .......................................................240 3.2.4.1 Gate Dielectrics through Anodization of Thin-Metal Films ......240 3.2.4.2 Surface Treatment of Inorganic Materials ..................................241 3.2.4.3 Self-Assembled Monolayers/Multilayers....................................242 3.2.4.4 Nanocomposite and Nanostructured Dielectrics ........................245 3.2.5 Summary and Conclusions ........................................................................248 References..............................................................................................................248

3.2.1 INTRODUCTION The rapid growth in research and development efforts in organic/flexible electronics underscores the importance of this new science and technology [1–4]. The field of organic electronics focuses primarily on the issues relating to the organic thin-film transistor device. This device forms the basis of organic integrated circuits. A schematic of the prototypical organic thin-film transistor (TFT) is illustrated in Figure 3.2.1. This chapter will focus on the selection and design of the gate dielectric, which is the insulating material separating the active semiconducting material from the

229


230


Semiconductor Source Drain Dielectric Gate Substrate

FIGURE 3.2.1 Schematic of bottom gate organic thin-film transistor.

gate electrode. This review is intended to bring researchers interested in the area of organic electronics up to speed on the factors involved in selecting a gate dielectric material. The review will discuss the basic issues involved, provide an overview on existing materials used as gate dielectrics, and finally present the newest developments in gate dielectrics. As can be seen in Figure 3.2.1, the dielectric layer in a typical thin-film transistor device is sandwiched between the gate electrode and the organic semiconductor. This geometry creates two interfaces, which must be considered when selecting a suitable dielectric material. When considering the dielectric–gate interface, one must address the issue of preventing static charge or dynamic charge injection into the dielectric, which can have adverse effects on threshold voltage (VT) as a function of time. Furthermore, complete coverage of the gate electrode is necessary to prevent leakage currents through pinhole defects. The dielectric-semiconductor interface is equally important to control, if not more so. This interface is where the conducting channel is formed. The quality of this channel is determined by interface roughness, surface energy, and charges at this interface [5]. Furthermore, in most TFT structures, the semiconductor is grown or deposited on top of the dielectric. In these instances, the dielectric serves to organize the semiconductor, especially in its first few layers. The structure of these layers is thought to be critical to the performance of organic semiconductors [6]. The basic issues in choosing a gate dielectric for an organic TFT are control of interfaces, leakage, dielectric constant, processability, stability, and reliability. Achieving high capacitances in organic TFTs is of great importance, considering the low mobilities (typically, >300 nm to achieve sufficient insulation, at which point capacitances drop to values of approximately 20 nF/cm2. PVP and PS at 20 nm thickness have very high leakage currents preventing the measurement of meaningful capacitances for these films. These ultrathin cross-linked films are generated by spin coating the substrate with the mixture of PVP and C6 and curing in air or under vacuum at ~110°C for 10–15 min. Chemical cross-linking (see Figure 3.2.5) of PVP and C6, as well as covalent linkages to the gate electrode, make this dielectric robust for further solution-based deposition steps. Furthermore, the substrates are shown to be patternable. In the case of evaporated gold masks, RIE or BOE can etch exposed polymer, leaving behind a patterned dielectric on removal of gold masks. The patterned structures exhibit excellent dielectric properties identical to nonpatterned films. Cross-linking of ultrathin PVP with C6 also renders this dielectric film less hygroscopic (a problem affecting PVP reproducibility). Low losses are reported ( 0.1 and on/off ratio ~ 104. Devices tested after 13 months show negligible degradation in performance. These cross-linked ultrathin polymer films show excellent promise, and further investigation in polymers beyond PVP will no doubt expand the utility of this dielectric design concept (see Table 3.2.4). Majewski et al. have looked at ultrathin SiO2 on metallized (aluminum metal) Mylar [32]. These systems involve 60-nm aluminum on Mylar as the gate electrode, with SiO2 3.5 nm (~ three times the limit predicted by theory and experiments on SiO2/Si). In these ultrathin SiO2 dielectrics, capacitances as high as 1 μF/cm2 are achieved. Addition of an OTS (octyltrichlorosilane) SAM to this system reduces the

TABLE 3.2.4 Dielectric Data for Ultrathin Cross-Linked Polymers, SAMs, and Anodized Films

Dielectric

K

Cross-linked PVP Anodized Ta2O5 Anodized alumina Anodized TiO2 with p(AMS)

6.4 23 9 21 for TiO2

Ultrathin SiO2 on aluminum Octadecyltrichlorosilane SIP/ROMP

3.9 — —

Minimum thickness (nm)

Capacitance (Ci (nF/cm–2)

Leakage (Acm–2)

10–20 50 130 8 nm TiO2 + 10 p(AMS) 3.5 2.8 nm 1,200

300 ~400 ~60 228; ~2400 for TiO2 alone 300 450 3

10–8 10–8 10–9

10–8 —

Ref. 20 31 37 36 32 19 27



245

capacitance to 300 nF/cm2. OFETs of pentacene and rr-P3HT on this dielectric perform with high mobility (0.12 cm2/Vs for pentacene and 0.01 for rr-P3HT). Excellent breakdown voltages are achieved (greater than 7 × 108 V/m). VT is ~–1 V. The device operates below 2 V and at this voltage produces currents ID ~ 1.8 μA. Without surface treatment, much lower (up to 10×) mobilities are observed, with larger leakage currents.

3.2.4.4 NANOCOMPOSITE

AND

NANOSTRUCTURED DIELECTRICS

The challenging materials requirements for a suitable gate dielectric for flexible organic thin-film transistors are not easily met by a single material. As has been demonstrated in previous sections, the use of multilayers comprising different materials is often required to balance the strengths and weaknesses of individual polymer and inorganic materials. For example, high-K ceramics offer excellent dielectric constants, but suffer from poor interfaces with organic semiconductors. Alternatively, polymer dielectrics typically have low capacitances, but excellent film-forming properties. In this section, we will explore recent efforts to blend ceramic and polymeric materials on the nanoscale to generate a nanostructured dielectric with an intelligent blending of the desirable properties of its constituents. The general strategy has been to load a polymeric material with high-capacitance nanoparticles. Table 3.2.5 summarizes the different composite materials described in this section. Early examples of polymer ceramic nanocomposites generated using micron and larger size particles of titanium oxide and polystyrene, in which composites were made through mechanical mixing, have been reported. Poor adhesion and air pockets presented a problem to this approach [63]. With increasing interest in nanotechnology and greater understanding of the control of nanoparticle synthesis, efforts to generate polymer dielectric films loaded with high-K dielectric materials appeared in the literature. An early example involved the sol-gel synthesis of bismuth titanate nanoparticles within a polyacrylate matrix [64]. Roughly 50-nm diameter particles of amorphous bismuth titanate were observed in the polyacrylate matrix using transmission electron microscopy (TEM). Effective dielectric constants as high as 10 were reported; however, detailed characterization of the material was not reported.

TABLE 3.2.5 Dielectric Properties of Nanoparticle Based Dielectrics Material

K

Thickness (nm)

Capacitance (Ci (nF/cm2)

Leakage (A/cm2)

Bismuth titanate-polyacrylate BaTiO3-PVA BaTiO3 Luxprint/smoothing layer TiO2-PVA TiO2-PS SiO2-PDDA LbL

10 9–12 — 5.4 8.2 6

— 170 10–15 μm 600 870 300

— 62.5 ~1 — 8 —

— 10–8 — — —

Ref. 64 66 54 65 35 26


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Efforts at loading titanium oxide nanoparticles in PVA (commercially available from Nanophase) have been reported [65]. In this report, titanium nanoparticles are dispersed in an aqueous solution of PVA with poly(melamine-co-formaldehyde). The solution is spun onto substrate and heated to generate a cross-linked polymernanoparticle dielectric. A modest enhancement of dielectric constant is achieved for 600-nm thick films. Thin-film transistors using this composite show excellent pentacene mobility (> 0.2 cm2V–1s–1) and reasonable on/off ratios ~ 103. VT ~ –7V is reasonably high, suggesting static charge at the dielectric–semiconductor interface. Efforts combining high-K materials such as barium titanate with lower K polymers to achieve processable flexible OFETs has been attempted by Zielke et al. [54]. In this effort, a BaTiO3-filled polymer (Luxprint 8153E) was used as a spin on dielectric to generate a top gate structure over offset printed source drain contacts, and spun on semiconductor (PTAA). The offset printed source drain electrodes were extremely rough (rms roughness of 75 nm with peaks extending as high as 3 μm). For this reason, a thick gate dielectric was required (10–15 μm thick) to smooth the roughness in the films. Barium titanate (high K) allows such thick films still to have reasonable capacitance, but BaTiO3 is highly polar and is thought to induce traps in the semiconductor, thereby reducing mobility dramatically [5]. In order to insulate the semiconductor from the BaTiO3, a second insulating layer of butylene copolymer (P105, 100–600 nm in thickness) that has a K of 2.2 is used. This low-K insulating film improves mobilities. Devices with Luxprint alone showed mobilities ~ 4 × 10–4, while mobilities in devices with the low-K interlayer increased to 3 × 10–3. The lack of control in the Luxprint structure creates a very ill-defined dielectric material, with widely varying threshold voltages (0 ± 8V) and significant hysteresis effects. However, this study clearly demonstrates the ability of high-K films to overcome problems such as rough electrode morphology. Low-K polymeric materials are not able to operate at such high thickness since the voltages required would be too high to turn on devices. Flexibility of the films is also demonstrated by repeated flexing of aluminum foil devices where no loss of function is observed. Another report of barium titanate nanoparticle composite dielectrics describes the combination of commercially available barium titanate nanopowders with PVA and also PVA-co-poly(vinyl acetate)-co-poly(itaconic acid) (PVAIA) in aqueous solution to generate a well dispersed nanocomposite that can be spun on as gate dielectric [66]. Pentacene-based transistors using this barium titanate/PVA nanocomposite show promising effective dielectric constants (11–12), low threshold voltages (–0.8–1.2 V), and high field-effect mobilities (~0.4 cm2V–1s–1). In this report, the level of loading of barium titanate was not reported. It was observed that higher dielectric constant films (~12) showed somewhat lower mobility than films with a K ~ 9 (mobility ~ 0.4 cm2V–1s–1). Prior attempts at nanocomposite dielectrics made no attempt to control the nanostructure of the polymer/dielectric material through controlling the chemistry between the nanoparticle and the dielectric. The first example of the use of a well characterized core-shell nanostructured dielectric material appeared in 2005 [35]. Researchers at Bell Labs used narrowly dispersed anatase phase titanium oxide nanoparticles (rod-shaped; ~15 × 4 nm; K = 31) as the high-K core material (see



247

Flexible high-K dielectric film Core TiO2

TiO2

TiO2

TiO2

TiO2 Polystyrene shell

TiO2 TiO2

TiO2 TiO2

100

TiO2

TiO2

TiO2

TiO2 TiO2

TiO2

TiO2

TiO2

TiO2

TiO2

Particle size distribution of TiO2-oleic

Volume

80 60 Particle size distribution of TiO2-PS

40 20 0 0

20

40 60 80 Particle diameter (nm)

100

FIGURE 3.2.6 Titanium oxide core–polystyrene (TiO2–PS) shell nanoparticle gate dielectrics. Bottom graph shows particle size distribution of as-synthesized TiO2–oleic, and polystyrene functionalized TiO2–PS particles as determined by DLS.

Figure 3.2.6). Narrowly dispersed polystyrene (synthesized by atom transfer radical polymerization [polydispersity < 1.1]) was end functionized with a phosphonate moiety that binds strongly to titanium oxide. The combination of narrowly dispersed titanium oxide and narrowly dispersed phosponate-terminated polystyrene generates a narrowly dispersed core-shell architecture as measured by dynamic light scattering, which can be spun into dielectric films. The covalent coating of polystyrene around titanium oxide is helpful at preventing aggregation of the nanoparticles in organic dispersion and in thin films. The resulting films show effective dielectric constants of ~8 with very low frequency dependence. The loading level for these core-shell films is much higher (almost 20 vol% titanium dioxide) as compared to other nanoparticle-polymer films. The low dielectric constant polystyrene shell (K = 2.5) allows for an excellent interface formation with evaporated organic semiconductors. Pentacene TFTs generated using these dielectrics show excellent mobilities (approaching 0.2 cm2V–1s–1) in unoptimized devices. On/off ratios are ~500 for unpatterned devices. The presence of titanium dioxide in these dielectric films reduces their breakdown voltage from >100 V/μm for polystyrene to ~50 V/μm for ~10 vol% films, and ~25 V/μm for 18 vol% films. Controlling the covalent chemistry between different components in a nanostructured composite dielectric is anticipated to be a general strategy to significant improvement in the dielectric material and its interface with organic semiconductors. Continued efforts to develop this core-shell technique with other high-K ceramic nanoparticles and through better tuning of polymer shells is under way at Bell Labs.


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An alternative nanostructured gate dielectric involving a self-assembly process using SiO2 nanoparticles has been reported [26]. In this procedure, a layer-by-layer approach is used to build up a silica nanoparticle top gate dielectric over a pentacene FET. The polyelectrolyte, poly (dimethyldiallylammonium chloride) (PDDA), is used with a thin layer of polystyrene to generate the first layer of a multilayer structure. SiO2 nanoparticles are then deposited by dipping of the substrate into a dispersion of these particles. The procedure is repeated several times to build up multilayers of silica nanoparticles and PDDA. After several layers, 300-nm films were produced that contained a silica volume fraction of ~70% with 10% PDDA and 20% air-filled pores. The dielectric constant for this silica-PDDA multilayer structure is found to be ~6 as compared with 3.9 for silica.

3.2.5 SUMMARY AND CONCLUSIONS This review has sought to illustrate a variety of different strategies employed to achieve high-capacitance, low-leakage, stable, and reliable dielectric films. The importance of the dielectric layer to the performance of organic TFT devices cannot be overstated. The dielectric is responsible for effective charge injection into the semiconductor and organization of the channel at the dielectric interface, and is critical to electrical parameters such as threshold voltage, leakage current, on/off ratio, and mobility. Various approaches involving inorganic, polymer, SAMs, anodization, and nanocomposites have been discussed. Many of these different approaches will find practical applications, depending on the requirements of the particular application and its prerequisite processing steps. Continued efforts will no doubt generate new approaches and improve existing strategies, generating a robust library of dielectric materials to enable the development of organic electronics, especially with regard to the development of mass printable organic and flexible thin-film transistor devices.

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8080_C004.fm Page 253 Tuesday, April 3, 2007 1:54 PM

4.1

Grazing Incidence X-Ray Diffraction (GIXD)

Tae Joo Shin and Hoichang Yang CONTENTS 4.1.1 Introduction................................................................................................253 4.1.1.1 Two Possible Geometrical Setups...............................................255 4.1.1.2 Bragg Peaks.................................................................................257 4.1.1.3 Bragg Rod Profile .......................................................................257 4.1.2 Interpretation of the Diffraction Data .......................................................258 4.1.2.1 Calculation of Structure Factor...................................................261 4.1.2.2 Calculation of Angle between a- and b-Axes.............................263 4.1.3 Examples....................................................................................................264 4.1.3.1 Poly(3-hexyl thiophene) (P3HT) ................................................264 4.1.3.2 Pentacene.....................................................................................269 4.1.3.3 Oligo Acene-Thiophene ..............................................................271 4.1.4 Concluding Remarks..................................................................................272 References..............................................................................................................273

4.1.1 INTRODUCTION X-rays, interacting weakly with matter, have long been used as an essential characterization tool to study the structure of bulk crystalline materials owing to their negligible multiple scattering and significant penetration depth. Recently, with the benefit of very intense x-ray sources such as synchrotron radiations, it has become possible to obtain surface and/or interface information selectively. One experimental technique is using the grazing incidence geometry (Figure 4.1.1). X-rays are known to be totally reflected on the surface of flat substances and strongly attenuated at small incident angles relative to the surface plane less than a critical angle [1]. Murra et al. used this x-ray total reflection phenomenon to develop a new technique for the structures of crystal surfaces and overplayed interfaces [2]. This method is referred to as grazing incidence x-ray diffraction (GIXD), which is also called grazing incidence x-ray scattering (GIXS), and has been commonly used to analyze “in-plane” crystal structures in a range from a few nanometers to several hundred nanometers beneath the surface on solids in air. Its applications have been reviewed in an article by Fuoss, Liang, and Eisenberger [3].

253


254


qz

GIXD

Z qtot

kin

θ

qxy

α 2θhor kout

θ

REFL

FIGURE 4.1.1 Scattering geometry for grazing incidence X-ray diffraction (GIXD) (and Xray reflectivity). The angle of incidence, θ, of the X-ray beam is less than the angle of total external reflection from the substrate. kin and kout are the wave vectors of the incident and reflected beams. The scattering vector, qxy ≈ 4πsinθhor/λ, is parallel to the substrate plane and qz = 2πsinαf /λ is perpendicular to it.

Theoretically, Vineyard described GIXD with a distorted-wave approximation in the kinematical theory of x-ray diffraction [4]. In terms of the ordinary dynamical theory of Ewald [5] and Laue [6], Afanas’ev and Melkonyan [7] worked out a formulation for the dynamical diffraction of x-rays under specular reflection conditions and Aleksandrov, Afanas’ev, and Stepanov [8] extended this formalism to include the diffraction geometry of thin surface layers. Subsequently, the properties of wave fields constructed during specularly diffracted reflections have been discussed in more detail by Cowan [9] and Sakata and Hashizume [10]. Meanwhile, a geometrical interpretation of GIXD based on a three-dimensional dispersion surface has been proposed by Höche, Brümmer, and Nieber [11]. Synchrotron radiation is very effective for GIXD measurements because of the parallel and intense beams and is often used for characterizing thin films on solids. Because of the limited availability of synchrotron radiation, conventional x-ray sources also have been adapted for the grazing incidence geometry [12]. Saito et al. have applied this method to the structural study of anodic oxide films and the thermally oxidized films on steels and obtained information on the thickness and density of the oxide films. The aim of this chapter is to systematically describe the steps involved in the elucidation of the in-plane order of thin films of organic semiconductors as a whole, not of the film surface alone. The principles of grazing incidence x-ray diffraction are described and a few examples of results obtained using GIXD are given. Organic thin films of submonolayer or monolayer are weak x-ray scatterers because they are composed of two-dimensional crystallites — namely, crystallites randomly oriented on the surface. Therefore, GIXD yields x-ray diffraction data that do not allow for a direct determination of the atomic positions as in three-dimensional crystals. Nevertheless, by making use of fixed atomic coordinate models of organic molecules, it is possible to extract valuable information on the molecular packing and orientation in the two-dimensional crystallites.



255

The general approach for the analysis of GIXD experiments is based on the kinematic theory. This theoretical framework is based on two simplified assumptions: a spherically symmetrical scattering electron density of atoms and insignificant contribution from multiple scattering events. The latter assumption is justified since, for typical surface diffraction studies, the in-plane coherence of the diffracting objects is small enough to preclude multiple scattering. In addition, due to the low atomic number of the organic materials, the incident beam suffers a very small attenuation by absorption. In the general GIXD geometry, the angle of incidence of the x-ray beam is kept below the critical angle αc. This limits the penetration depth of the beam to that of the evanescent wave, and the scattering due to the subphase is efficiently eliminated. This permits a reliable measurement of the weak diffraction signal originating from the crystalline monolayer. When the substrate is disordered and thus scatters in all directions, the angle of incidence of the beam should be below the critical angle of the substrate in GIXD geometry. Conversely, if substrates are crystalline, as long as one stays away from the substrate Bragg peaks, no particular effort is needed to have the incident beam at less than the critical angle. Furthermore, this geometry increases the number of incident photons by ensuring that the whole thin film participates in the diffracted signal. From the GIXD profiles, the in-plane lattice parameters can be derived simply by using the Bragg’s law. Figure 4.1.1 shows schematic representation of grazing incidence diffraction from a horizontal surface. The diffraction condition for two-dimensional crystals lying in the xy plane is that the component of the scattering vector in the horizontal plane, labeled qxy (qxy = q|| ≈ (4π/λ)sinθhor (where 2θhor is the in-plane angle between the incident and the diffracted beam), must coincide with a reciprocal lattice vector qhk = 2π(ha* + kb*), where a* and b* are the reciprocal in-plane lattice vectors and h and k are the linear coefficients for the corresponding lattice point. There is no similar selection rule or restriction on the component of the scattering vector along the normal to the film defined as qz, whose magnitude qz = q⊥ = (2π/λ)sinα, where α is the angle between the diffracted beam and the substrate surface, shown in Figure 4.1.1. Therefore, the GIXD patterns from two-dimensional crystals (crystallites) arise from a two-dimensional array of rods, called Bragg rods (BRs), which extend parallel to qz.

4.1.1.1 TWO POSSIBLE GEOMETRICAL SETUPS As shown in Figure 4.1.2(a), the collection of the diffracted radiation by means of a one-dimensional position-sensitive detector (1-D PSD) is made by scanning the detector over a range along the horizontal scattering vector qxy (≈4πsin θhor/λ) and integrating over the whole qz window of the 1-D PSD to yield the Bragg peaks. Simultaneously, the scattered intensity recorded in channels along the 1-D PSD, but integrated over the scattering vector qxy in the horizontal plane across a Bragg peak, produces qz-resolved scans called Bragg rods. The GIXD method has been commonly used to analyze in-plane crystal structures of thin surface layers. However, measurements of “out-of-plane” structures under a grazing-incidence angle can also be performed. If two-dimensional (area) detectors


256


Side View

α f

1-D PSD

α f α

αi

f X-rays

Incident footprint

2θho r

Top View

Soller collimator

1-D PSD

(a)

Two-dimensional (area) Detector

qz α X-rays

α f

f

αi

r

2θho

qxy

(b)

FIGURE 4.1.2 (a) Top and side view of GIXD geometry. The 1-D PSD has its axis along the vertical. Only the cross-beam are contributes to the measured scattering. The Soller collimator defines the horizontal resolution of the detector. (b) 2-D GIXD geometry under which both in-plane (qxy) and out-of-plane (qz) reflections can be measured simultaneously. (From Kaganer, V.M. et al., Rev. Mod. Phys. 71, 779, 1999. With permission.)

like CCD camera and image plates are used, simultaneous measurements of both inplane and out-of-plane crystal structures become possible (Figure 4.1.2b). Minor and diffuse diffraction arcs can also be recognized more easily by using two-dimensional (area) detectors. However, the terminology “grazing-incidence x-ray diffraction technique” requires only that the incidence x-ray beam impinge the sample under a grazing angle; there is no restriction on the angle of the diffracted beam. Therefore, more information can be obtained if more generalized diffraction angles are measured. Such measurements can be easily realized with two-dimensional detectors such as image plates and a spherical-type goniometer [13] in combination with a grazing-incidence-angle synchrotron radiation beam, as shown sche-



257

matically in Figure 4.1.1. In this configuration the Seemann–Bohlin geometry (out of plane: 2θout) or, similarly, grazing-incidence-angle asymmetric Bragg geometry [14] and GIXD geometry (in plane: 2θin) can be used simultaneously without scanning detectors. The average incident angles, θin, are around 0.3–0.5°, depending on the critical angles of the surface layers of the samples. Since the incident angles that can achieve a total reflection condition are usually very low, in most cases conventional divergent slits cannot effectively restrict the vertical width of the incident beam, and the beam unavoidably spills over the samples. Therefore, the sample width towards the beam direction must be carefully chosen as a compromise between spatial resolution and diffraction intensity, depending on the purpose of the analysis.

4.1.1.2 BRAGG PEAKS The reflections of Bragg peaks can be indexed by two Miller indices, hk. Their angular positions, 2θhor, yield the lattice plane spacing dhk = 2π/qhk for the twodimensional lattice structure. It is possible to have an estimate of the dimensions of the domains (DL) by analyzing the full width at half maximum (FWHM) of a GIXD peak (Δexp), corrected from the instrumental broadening with the Scherrer equation:

DL =

0.94λ 2 1/ 2 , where Δ = ( Δ 2exp − Δ slit ) Δ cos θ

and θ corresponds to the angular peak position [15,16]. The square of the molecular structure factor |Fhk|2, integrated along the Bragg rod over the window of qz seen by the detector, determines the integrated intensity in the peak. The structure factor Fhk(qz) is given by Fhk (qz ) =

∑fe j

iqhk ⋅rj iqz z j

e

(4.1.1)

j

where the sum is over atoms labeled j in the unit cell of dimensions a, b. The jth atom has a scattering factor fj and is located at the lateral vector position rj = xja + yjb and at the vertical position zj (in angstroms). Because the sample consists of a two-dimensional powder, the observed intensity Ihk for _a_ given qhk (or 2θhk) and qz position contains contributions from both the (h,k) and (h,k) reflections, with possibly several h, k integer values.

4.1.1.3 BRAGG ROD PROFILE The variation of the intensity Ihk(qz) along the Bragg rod as a function of qz is given by 2

−2 I hk (qz ) = KLPAABCD Acell V (qz ) Fhk (qz ) DWhk (qz )

(4.1.2)


258


The observed Bragg rod intensity Ihk(qz) is actually a sum over those (h,k) reflections whose Bragg rods coincide at a particular horizontal 2θhor angle or qxy position. In the upper equation, the most important variation is due to the molecular structure factor amplitude |Fhk|2. The Debye–Waller factor DWhk(qz) = exp[–(qhkUxy+ qz2Uz), where Uxy and Uz are, respectively, the mean square displacements in each horizontal direction x, y and along the vertical direction z. They account for the thermal motion of the atoms in the molecule, as well as for possible ripples on the substrate, which lead to roughness of the interface. The grazing geometry factor V(qz) describes the interference of rays diffracted upward with rays diffracted down and subsequently reflected back up by the interface. V(qz) differs from unity value only in the vicinity of q = qc/2, where it contributes a sharp peak. Corrections for the crossed-beam area (AABCD ∝ 1/sin2(2θhk)), Lorentz (L ∝ 1/sin2(2θhk)), and polarization (P = cos2(2θhk)) have been inserted; Acell is the area of the unit cell. The factor K scales the calculated to the observed intensities. For the simple linear surfactant molecules considered here, the square of the molecular structure factor |Fhk|2 is bell shaped as a function of qz and reaches its maximum when the scattering vector q = (qhk, qz) is orthogonal to the molecular axis. Thus, when the molecules are vertical or tilted in a plane perpendicular to qhk, the maximum intensity along the Bragg rod will occur at the horizon for qz = 0 Å–1. For molecules tilted otherwise, we expect the Bragg rod maximum at a finite qz, dependent upon both the magnitude and direction of the tilt relative to the in-plane scattering vector. The width of the bell-shaped Bragg rod profile is inversely proportional to the length of the molecule.

4.1.2 INTERPRETATION OF THE DIFFRACTION DATA In this section, we summarize the general approach to analyzing the GIXD data from monolayers, following Kaganer et al. [17]. In GIXD experiments, the scattered intensity is monitored as a function of two angles: the angle between incident and scattered beams in the water plane and the angle between scattered beam and the substrate. A periodicity of the molecular arrangement in the monolayer manifests itself in a peak in the distribution of the scattered intensity. There is as yet no way of controlling the mosaicity of Langmuir monolayers; in other words, the monolayers are powders within the plane. The diffraction pattern is always averaged over all domain orientations in the monolayer plane (“powder averaging”) in the area illuminated by the incident x-ray beam. As a result, of the three components qx, qy, and qz of the momentum transfer vector q, only the vertical component, qz, can be measured independently. It is not possible to determine the in-plane components qx and qy individually, but only the combination qxy qxy = (qx2 + qy2)1/2. Because lattice fluctuations at room temperature cause the peak intensities to decay rapidly with increasing momentum transfer, the first-order peaks, which correspond to the distances between neighboring molecules, are usually the most intense and frequently the only observed ones. Two first-order peaks sharing the same qxy are an indication of hexagonal packing, with equal distances between the molecules; three distinct values of qxy point to a rectangular unit cell and four peaks are due to



259

an oblique unit cell. The available diffraction data are obviously not sufficient to perform structural analysis in a classical crystallographic sense. Additional knowledge of the packing of the organic semiconductor molecules in crystals may also be necessary because the packing sometimes leads to extinctions in the diffraction intensity and the aforementioned rule of thumb for first-order peaks would not apply (see discussion in Section 4.1.3). If the incident beam is coming in at less than the critical angle for total reflection, the specular beam does not carry useful information. However, if there is any lateral order at the surface (or in a monolayer at the surface) there will be diffraction peaks. The diffracted beam need not be observed only at grazing exit angles, but may be detected at any angle, resulting in a q that has both horizontal and vertical components. As shown in Figure 4.1.3b, the reciprocal lattice of a two-dimensional lattice, such as a thin film or a crystal surface, is a two-dimensional lattice of rods (since there is no periodicity in the z-direction, all qz are equivalent as far as the Bragg condition is concerned). The reciprocal lattice (and thus the real lattice) can be determined by analyzing the qx and qy spacings between these lattice rods. Notice that in the grazing incidence geometry it is never possible to go to qz = 0 Å–1 because this would require that the incident and diffracted photons make zero angles to the surface. However, a two-dimensional system is forgiving: the reciprocal lattice is in the form of rods, so the diffraction peaks can just as easily be located at slightly nonzero qz. (The scattering intensity will vary strongly along the rods if the lattice is not composed of points but of molecules, and this intensity variation contains information regarding the orientation of the molecules.) In the simplest model, a domain of the monolayer is treated as a two-dimensional crystal consisting of uniformly oriented rigid molecules. The scattering pattern in reciprocal space is then given by the product of two factors: the structure factor reflecting translational order of the molecular centers in the plane of the monolayer and the form factor of the individual molecule (see Figure 4.1.3b). The structure factor of a two-dimensional lattice consists of a set of delta-function discontinuities along lines (Bragg rods) normal to the monolayer plane. The form factor of a long rod-like molecule (such as alkane or acene) is large only on a plane normal to its long axis, which will be called the reciprocal disk of the molecule. The intersections of the first-order Bragg rods with the reciprocal disk give rise to six diffraction maxima (Figure 4.1.3c–f). If the molecules do not tilt, the reciprocal disk and hence all the peaks lie in the plane of the monolayer (Figure 4.1.3c). In a hexagonal phase, all six first-order wave vectors q have equal length and overlap completely in the powder pattern. Because of this degeneracy, the sixfold symmetry can only be concluded from the failure to see any other first-order peaks. The degeneracy is lifted in cases where the lattice is distorted from the exact symmetry hexagonal. For example, as a result of ordering of the backbone planes of the molecules, distinct peaks at different values of q are observed. If the unit cell is stretched or is compressed in the direction of the nearest-neighbor molecule, the unit cell becomes centered rectangularly. There are then two distinct first-order wave vectors on the powder averaging: one pair with ±qn and the other two with ±qd (the subscripts n and d denote nondegenerate and degenerate peaks).


260


Incident beam

Real space

Diffracted beam

Reciprocal space

2D lattice Rod-like molecule

Specular beam

Reciprocal disk

(a)

(b)

Untilted

NN tilt

NNN tilt

I

Kz

I

Kxy

Kxy Kz

Kz

Kxy (c)

θ

I

Kxy

Kxy Kz

Kxy (d)

Intermediate tilt

θ

θ

I

Bragg rods

Kxy (e)

Kxy (f )

FIGURE 4.1.3 (a) Schematic diagram of a GIXD experiment; (b) ‘Rods’ formed in reciprocal space by a 2-D lattice of points, and the ‘reciprocal disk’ of an extended molecule, which combines to give the diffraction pattern from a 2-D array of extended molecules; (c) realspace and reciprocal-space views and characteristic diffraction patterns of the monolayer in untilted phase; (d) NN-titled phase; (e) NNN-titled phase; (f) intermediate titled phase. (From Kaganer, V.M. et al., Rev. Mod. Phys. 71, 779, 1999. With permission.)

If the unit cell stretches in the direction of the nearest-neighbor (NN) molecule, then |qn| > |qd|; the opposite inequality indicates that the unit cell shrinks in that direction. The degeneracy may also be lifted by molecular tilt. In this case, the peaks move out of the monolayer plane by a distance qz, which depends on both the tilt magnitude and its azimuth, or direction. Since the only points of the reciprocal disk to remain in the monolayer plane are those on the line normal to the tilt direction, the only diffraction peaks that originate from Bragg rods on that line remain in the monolayer plane (see Figure 4.1.3d). This occurs for one pair of peaks when the molecules tilt towards the NN. The other four peaks move out of the plane: two upwards and two downwards (naturally, diffraction beams going into the substrate cannot be observed); see Figure 4.1.3(d). The wave vectors of the two visible out-



261

of-plane peaks have equal qz components and are thus degenerate in the powder pattern. The tilt angle θ is given by tan θ = qdz/[qdxy2–(qnxy/2)2]1/2. When the molecules tilt towards a next-nearest neighbor (NNN) molecule, all the wave vectors move out of the plane. The two distinct values of qz are in the ratio qnz:qdz = 2:1, and the tilt angle is given by tan θ = qnz/qnxy (Figure 4.1.3e). In these symmetrically tilted phases, the distinction between degenerate and nondegenerate peaks is unequivocal. The ratio qnz:qdz can only be 0:1 or 2:1. In an untitled phase (Figure 4.1.3c), the distinction is not as easy. In the idealized model, where rod-like molecules are represented as cylinders of uniform electron density, the integrated intensity of the degenerate peak should be twice as large as that of the nondegenerate one, but experimentally there are often significant departures from the “ideal” 2:1 intensity ratio. Leveiller et al. calculated the molecular structure factors and intensity ratios from atomic scattering factors, assuming an all transconformation of the molecules and ideal (zero-temperature) packing; they found that the intensity ratio depends on the orientations of the backbone planes of the molecules and symmetry constraints [18]. If the tilt azimuth is intermediate between NN and NNN (Figure 4.1.3f) or if the distortion of the unit cell is asymmetrical, there are three distinct first-order peaks and no indexing problem. Each peak is described by two components of the momentum transfer, giving six measured values in all. Since the monolayer model is completely described by five parameters — three for the in-plane lattice and two for the tilt (magnitude and direction) — the measured values cannot be completely independent. The relationship between them is readily shown to be q1z + q2z = q3z, where peak 3 is the one with largest qz. The two qz ratios 0:1 and 2:1 for the symmetric tilts follow from this more general relationship as particular cases in which one of the qz values is repeated and correspond to 0 + 1 = 1 and 1 + 1 = 2, respectively. Once the three first-order peaks have been assigned, the geometry of the unit cell in reciprocal space is completely determined, and the corresponding real-space lattice is easily determined. In the literature, the peaks are sometimes labeled in terms of either a hexagonal or a centered rectangular unit cell; the latter notation is more common [19]. Denoting the basic translations of the centered rectangular unit cell containing two molecules by [10] and [01] (when the lattice is hexagonal, the length of the vector [01] is √3 times larger than [10]), one finds that, if the two molecules in the rectangular unit cell are actually equivalent (i.e., the rectangular unit cell is not the smallest unit cell describing the lattice), the reflections (01) and (10) are forbidden, as are all (hk) reflections where h + k is an odd number. The lowest order reflections are (02), which is nondegenerate, and two reflections (11) _ + (11), which have equal length and so degenerate in the powder average. In case of the hexagonal unit cell, all three reflections possess equal wave-vector magnitudes.

4.1.2.1 CALCULATION

OF

STRUCTURE FACTOR

Generally, the intensity at a particular value of qz in a Bragg rod is proportional to the square of the molecular structure factor, |Fhk(qz)|2. Thus, Bragg rod intensities can be calculated using the atomic coordinates in the unit cell. The form factor is given by


262


F (q) =

∑f e j

− iqr j

where fj is the scattering factor (or form factor) of the jth molecule in the unit cell at the position rj. If two molecules per a centered rectangular unit cell were related by translation symmetry (f1 = f2 = f) — that is, the same molecules are located at (0,0) and (a/2,b/2), then F(qx, qy) is given by F (q x , qy ) = f + f ⋅ e

i / 2 ( q x a + qy b )

At the peak positions, qx =

2π h a

qy =

2π k b

and

F(q) = F(h, k) = f + f ⋅ eiπ(h+k) vanishes whenever h + k is odd. When two molecules in the unit cell are not translationally equivalent, F(h,k) is given by F(h,k) = f1 + f2 ⋅ eiπ(h+k), where it can be observed, depending on the f1 and f2, even if h + k is odd. As discussed previously, Durbin et al. observed the (21) reflection in herringbone packing due to f1 ≠ f2 in the (21) reflection, but (10) or (01) reflections are not observed because the two form factors are identical [10]. For a unit cell with different dimension along the a, b, and c-axes, when a molecule stands with height (d) along the c-axis and width (w) along the a-axis, the scattering factor is given by

f (q x' , qz' ) =

sin(q x' w / 2 ) sin(qz' d / 2 ) − iqz' d / 2 ⋅e q x' w / 2 qz' d / 2

(4.1.3)

If the molecule is tilted, the sample coordinate and laboratory coordinate are related by q ' = R -1 ⋅ q

(4.1.4)

where R is Euler rotation matrix expressed by the tilt angle (θ), tilt azimuth (ϕ), and the rotation angle to the molecular axis (χ). q′ and q are scattering vectors in the sample coordinate and laboratory coordinate, respectively. The molecule is



263

located at the sample coordinate and the unit cell with the molecules is located at the laboratory coordinate. Finally, from the equations 4.1.3 and 4.1.4, a form factor can be calculated. A structure factor of the herringbone motif can be calculated by applying –χ and +χ to each molecule.

4.1.2.2 CALCULATION

OF

ANGLE

BETWEEN A- AND B-AXES

If one assumes the real space lattice vector a = (ax, 0) and b = (bx, by), then the angle γ between the two vectors is defined as the angle from a to b and is related as follows: γ = tan −1 (by / bx )

(4.1.5)

A two-dimensional lattice can be thought of as the x–y plane of a monoclinic lattice, a* = 1 / (| a | sin( γ )) ⋅ b ⊥ / | b ⊥ |

(4.1.6)

b* = 1 / (| b | sin( γ )) ⋅ a ⊥ / | a ⊥ |

(4.1.7)

Since |a| = ax > 0, sin(γ) = by/|b|, and by ≥ 0, then a* = 1 / ( a x by ) ⋅ (by , −bx )

(4.1.8)

b* = 1 / ( a x by ) ⋅ (0, a x )

(4.1.9)

since the peak position of the (h,k) plane is defined as following q hk = 2π( h ⋅ a* + k ⋅ b* )

(4.1.10)

When the molecules in the lattice are tilted, a* and b* will have z component and qhk also should have qz component. However, the in-plane components of qhk are not dependent on the molecular orientation in the lattice at all. Thus, we can write | q hk |xy ≡ q xyhk =

2π a x by

h 2 (bx2 + by2 ) + k 2 a x2 − 2 hka x bx

(4.1.11)

Thus, if three peaks are resolved, the lattice vector could be calculated by the 11 12 linear algebra. For example, when q11 xy , q xy and q xy are resolved, they could be calculated as follows:


264


2 by = 2 π 3 / (q12 xy ) +

1 11 2 3 11 2 (q xy ) − (q xy ) 2 2

(

11 2 2 (q11 xy ) − (q xy ) 2 (2 π)2 11 2 11 2 a x = 2 π 2 / (q xy ) + (q xy ) − − 8(2π)2 by2

bx =

2 11 2 (q11 xy ) − (q xy )

4(2π)2

a x by2

(4.1.12)

)b 2

2 y

(4.1.13)

(4.1.14)

4.1.3 EXAMPLES 4.1.3.1 POLY(3-HEXYL

THIOPHENE)

(P3HT)

Self-organization in many solution-processed, semiconducting conjugated polymers results in complex microstructures in which ordered microcrystalline domains are embedded in an amorphous matrix [20,21]. This has important consequences for electrical properties of these materials: Charge transport is usually limited by the most difficult hopping processes and is therefore dominated by the disordered matrix, resulting in low charge-carrier mobilities (≤10–5cm2V–1s–1) [22]. Sirringhaus et al. have studied the microstructure of 70- to 100-nm, spin-coated, regioregular poly(3-hexylthiophene) (P3HT) films by GIXD on parts of the SiO2/Si substrates on which field effect transistor (FET) devices were fabricated [20]. Grazing-incidence x-ray diffraction measurements were performed under inert He atmosphere to minimize air scattering and beam damage. A grazing incidence angle below the critical angle of total reflection from the substrate, but above the critical angle for the film, was chosen to enhance the sensitivity to the thin polymer film. The intensities plotted versus the total scattering vector are corrected for polarization and geometric factors [23]. Depending on processing conditions, the π-conjugated crystalline planes can adopt two different orientations — parallel and perpendicular to the substrate (the mobilities of which differ by more than a factor of 100) — and can reach values as high as 0.1 cm2V–1s–1 [24]. Two different orientations of the microcrystalline P3HT domains with respect to the FET substrate are shown in Figure 4.1.4. The coexistence of two different phases is evident from the different intensity distributions of the (100) reflections due to interdigitation of hexyl side chains and the (010) reflections due to π–π-interchain stacking [25]. In samples with high regioregularity (>91%) and low molecular weight, the preferential orientation of ordered domains is with the -axis normal to the film and the -axis in the plane of the film (Figure 4.1.4a). In contrast, in samples with low regioregularity (81%) and high molecular weight, the crystallites are preferentially oriented with the -axis in the plane and the -axis normal to the film (Figure 4.1.4b) [26]. The cause of this surprising orientational



265

change could be understood as a dynamic phenomenon occurring during the rapid growth of spin-coated films, affected by regioregularity or molecular weight. In films prepared by drop casting from a dilute solution, the -axis is normal to the film for all polymers (gray trace in Figure 4.1.4c and d). The ability to induce different orientations allows one to establish a direct correlation between the direction of π–π-stacking and the in-plane FET mobility. At room temperature, the highest mobilities of 0.05–0.1 cm2V–1s–1 are observed for the sample with the highest regioregularity (96%) and the largest size of crystallites with in-plane orientation of the -axis (~95Å, as calculated from GIXD analysis for (010) peak shape) [23]. For spin-coated samples with 81% regioregularity (-axis normal to the film), the mobility is only 2 × 10–4 cm2V–1s–1 in spite of pronounced in-plane crystallinity along the -axis with a grain size of 130 Å. This is consistent with the FET mobility being limited by π–π-interchain, rather than intrachain, transport. In the case of the low-regioregularity polymers, it is possible to compare mobilities directly for the in-plane (cast films) and out-of-plane orientation (spin-coated films) of the π–π-stacking direction (Figure 4.1.4d). In cast films of the 81% regioregular polymer, the mobility is higher by more than an order of magnitude than in spin-coated samples and only slightly lower than that of the highly regioregular polymers (Figure 4.1.4e). The large mobility anisotropy caused by different preferential orientations of the ordered, microcrystalline domains is clear evidence that the transport is no longer dominated by the remaining amorphous regions of the polymer film. Rather, it is starting to reflect the transport properties of charge carriers in ordered polymer domains. McGehee and coworkers observed that the charge mobility values of spin-cast RR P3HT films changed up to four orders of magnitude depending on the MW [27,28]. They observed that the crystalline morphology of low MW films (Mn < 4 kDa) showing low mobility had highly ordered short nanoribbons, while highmobility, high MW films (Mn > 30 kDa) had featureless crystalline structure. Because the charge transport in organic thin film transistors (OTFTs) occurs in the plane of the substrate, GIXD was used to characterize the in-plane π-stacking. However, because in-plane GIXD measures crystals with lattice planes perpendicular to the substrate (within about 0.7°), only a small portion of reciprocal spaces samples and it is likely that crystals not measured by in-plane GIXD can be important for charge transport. In contrast, rocking-curve x-ray experiments, where the sample is rotated in the x-ray beam while keeping the scattering angle (2θ) constant, provide a means for measuring the distribution of crystal orientations in the films throughout reciprocal space. The rocking curve x-ray experiments for different MW P3HT thin films was used to explain change in charge mobilities through dielectric surface treatment with hexamethyldisilizane (HMDS) and octadecyltrichlorosilane (OTS) and postannealing [29]. Figure 4.1.5 shows the specular diffraction obtained for the (100) peak of spin-cast films with various MWs on OTS-treated substrates. The peak intensity, width, and position vary with MW. The trend of reduced d(100) spacing in the films with decreasing MW is related to either a change in molecular tilt angle or the amount of side-chain interdigitation [28,30]. Also, the reduced peak width indicates


266


s

a

s s

s s

s

a

s s

b

Intensity (a.u.)

5 b

4.5 4 3.5 3 2.5

(010) (300)

1.5

(200)

1

(100) (010)

(100) (b)

γ

(300)

Side view (200)

0.4 0.2 0 0.0

(010)

Detector

α (100)

Intensity (a.u.)

(a)

0.6

0.5

1.5

1.0

2.0

q (Å−1)

(010)

(c)

2θ

81% cast

Top view (200)

(100)

Intensity (a.u.)

Detector 0.4

0.2

81%

96%

0 0.0

0.5

1.0

1.5

2.0

q (Å−1) (d)

μsat (cm2 V−1 s−1)

10−1 10−2 10−3 10−4 70

80 90 % head-to-tail (e)

FIGURE 4.1.4

2

2.5

0.5



267

0.06

Counts (arb. units)

2θ Low-MW OTS 0.04

Medium-MW OTS 0.02

High-MW OTS

0.00 0.25

0.30

0.35

0.40

0.45

q (Å−1)

FIGURE 4.1.5 Specular diffraction of the peak. Out-of-plane diffraction peaks analyzed with rocking curves shown in Figure 4.1.6 for the P3HT films with low, medium, and high molecular weights on OTS-treated substrates. The inset shows the specular diffraction geometry. (From Kline, R.J. et al., Nat. Mater. 5, 222, 2006. With permission.)

that the domain size perpendicular to the films is larger for low MW films than it is for medium and high MW films. In addition, Figure 4.1.6 shows rocking curves for the orientation of the (100) crystalline planes relative to the substrate normal. It was suggested that the sharp peaks in the rocking curves, Δω ~ 0.03°, corresponding to the instrumental resolution are related to P3HT crystals located near the substrate (with r.m.s. roughness of about 1 Å) rather than the film surface (with r.m.s roughness of about 6 Å). The contribution of the crystals at the air–film interface (with high surface roughness) should have increased at least an order of magnitude greater than the width of the measured rocking-curves peak. Specifically, the different crystalline orientation in the films cast on OTS- and HMDS-treated substrates strongly supports their interpretation for the location of the highly oriented crystals. As a result, the fact that the low MW film cast on HMDS has few highly oriented crystals probably explains why the mobility (~10–6 cm2V–1s–1) is so much lower than for the other films. The anisotropy of PHT crystals limits charge transport to two crystal directions (

FIGURE 4.1.4 (See figure, facing page.) Two different orientations of ordered P3HT domains with respect to the FET substrate: (a), (b), 2-D GIXD distribution from spin-coated, 70-100 nm thick P3HT films with regioregularity of 96% (a) and 81% (b) on SiO2/Si substrates. The vertical (horizontal) axes correspond to scattering normal (parallel) to the plane of the film. The insets show schematic orientations of the microcrystalline grains with respect to the substrate. (c), (d), The change of orientation is confirmed by high-resolution synchrotron X-ray diffraction measurements for constant, grazing-incidence angle with outof-plane (c) and in-plane (d) scattering geometry. (e), Charge carrier mobility of P3HT-based field-effect transistors with different microstructures. (From Sirringhaus, H. et al., Nature 401, 685, 1999. With permission.)


268


10−4

0.0002

2θ

Counts (arb)

Counts (arb)

0.0003

10

LowMW OTS

−5

MediumMW OTS

ω

0.0001

10

−6

HighMW OTS −0.04 −0.02 0.00 0.02 0.04 ω (°)

0.0000 −1.5 −1.0 −0.5

s s s s s s s s s s s s s s s s

0.0

0.5 1.0 ω (°)

1.5 (d)

(a)

Counts (arb)

0.1

0.01

MediumMW OTS

0.001 Medium-MW HMDS

−0.2

−0.1

(e)

0.0 ω (°)

0.1

0.2

(b)

Counts (arb)

1 Annealed lowMW OTS

(f )

0.1

0.01

Low-MW OTS

Annealed lowMW HMDS

Low-MW HMDS

−0.2

−0.1

0.0 ω (°)

0.1

0.2 (g)

(c)

FIGURE 4.1.6 Rocking-curve measurements on the (100) peaks and resulting crystal orientations: (a) Rocking curves for the films (with low, medium, and high molecular weights) on OTS-treated substrates. (Left inset: the rocking curve geometry showing the angle relative to the sample normal ω. Right inset: a magnification of the ω = 0 region plotted on a logarithmic scale.) Logarithmic–scale rocking curves on films with (b) medium molecular weight on both HDMS and OTS, and (c) low molecular weight before and after annealing for both HMDS and OTS. Schematics showing the highly oriented and misoriented crystals in films with (d) low molecular weight on OTS, (e) medium and high molecular weights on OTS, (g) low molecular weight on HMDS before annealing, and (f) after annealing (drawn on a smaller scale to accommodate the larger crystals). Lines correspond to the (100) plane. The bottom circles in (e) and the first and third circles in (f) denote crystals measured by specular diffraction and the other circles denote those that are not. (From Kline, R. J. et al., Nat. Mater. 5, 222, 2006. With permission.)



269

and ), with the insulating hexyl chains preventing charge transport in the direction. A large distribution of crystal orientations increases the likelihood of having poor electronic overlap along these high-mobility directions between neighboring grains.

4.1.3.2 PENTACENE Pentacene, consisting of five fused benzene rings, has emerged as a viable candidate for use in OTFTs because of its relatively high hole mobility and high on/off ratio [31,32]. However, little is known about the crystal structure in the first few layers of a vacuum sublimed pentacene film, the most important for charge transport. It is also widely recognized that the transport properties of crystalline organic films depend strongly on the intermolecular overlap of electronic wave functions within the semiconductor layer, which is very sensitive to the molecular packing in the crystal [33]. Fritz et al. reported preliminary GIXD data for a monolayer-thick pentacene film grown on amorphous silicon dioxide (a-SiO2) [34]. The data confirmed that the monolayer is crystalline and has a structure that differs from that of bulk pentacene, which has important implications for carrier transport in pentacene-based OTFTs. The crystal structure of bulk pentacene consists of layers of pentacene molecules arranged in a herringbone packing motif with an interlayer spacing of d001 = 14.1 Å [35]. Three thin film multilayer phases with different d001 values of 14.4, 15.0, and 15.4 Å have been identified by wide-angle x-ray diffraction [36,37]. The selectivity toward these phases appears to be governed by a variety of factors, including substrate material, substrate temperature during deposition, rate of deposition, and film thickness [36,38]. The different d001 values imply dissimilar packing of the pentacene molecules in the ab plane, which is the basal plane in crystalline vacuum deposited pentacene films on inert substrates and is therefore the crystal plane crucial for charge transport. Characterization of this pentacene monolayer by GIXD at room temperature yielded a diffraction pattern (Figure 4.1.7) that could be indexed as a near rectangular in-plane unit cell with dimensions a = 5.92 Å, b = 7.59 Å, and γ = 89.95°. These values differ from the corresponding lattice parameters reported for bulk pentacene (a = 6.27 Å, b = 7.78 Å, and γ = 84.68˚), but are consistent with a packing motif resembling the (001) layers in the bulk (i.e., the pentacene molecules adopt a nearvertical orientation in the monolayer on the a-SiO2 substrate). A model of the monolayer structure was constructed using the room-temperature single-crystal structure of bulk pentacene as a starting point. The a and b lattice parameters of the bulk structure were adjusted to the monolayer values, and the interlayer d001 spacing was expanded to an arbitrarily large distance to mimic an isolated monolayer. The simulated diffraction pattern produced by this energy-minimized monolayer structure is in reasonable agreement with the GIXD data. Some of the discrepancy can be attributed to contributions from the a-SiO2 background, particularly in the lower qxy region. We anticipate that a more complete data set, particularly data collected at qz > 0, will permit precise determination of the molecular tilt and faster refinement of the thin-film structure [39]. Nevertheless, the GIXD data demonstrate unequivocally that the pentacene monolayer on the a-SiO2 dielectric layer is highly crystalline and exhibits a structure distinguishable from the bulk.


270


Bulk pentacene a b

a

Intensity (a.u.)

Expanded 3x

(1 1)

1

2

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

(1 3), (2 2)

(2 1)

(1 2)

(2 0)

Expanded 3x (0 2)

0

3

z b Pentacene monolayer d e

Calculated

Experimental

c

a

4

z

f

z b

a

b

z b

a

qxy (Å−1) (a)

(b)

FIGURE 4.1.7 (a) GIXD pattern (bottom of (a)) for a pentacene monolayer and a diffraction pattern (top of (a)) calculated for an energy-minimized crystal structure model based on the GIXD lattice parameters and the (001) layer motif of bulk pentacene as the starting point; (b) Normal views of the ab planes of bulk pentacene and the model monolayer structures (a and d) and the respective side views (b and e, e and f). The z-axis is the normal to the ab plane. (From Fritz, S.E. et al., J. Am. Chem. Soc. 126, 4084, 2004. With permission.)

The charge carrier mobility of organic semiconductors is sensitive to the dielectric layer surface on which it is deposited. Kelley et al. obtained very high field effect mobility (~3 cm2V–1s–1) with pentacene multilayers consisting of large twodimensional grains on a 1-phosphoneo-hexadecane-treated alumina dielectric layer [40]. The molecular orientation and grain morphology of the first layer of pentacene molecules depend on balance between pentacene–substrate interactions, which can be shifted by modification of the dielectric surface in an OTFT with a self-assembled monolayer (SAM) [41]. They have studied the correlation between pentacene ultrathin film morphology and the overall OTFT device performance and observed a direct correlation between the crystalline structure of the initial submonolayer of a pentacene film and the mobility of the corresponding thick film. The terrace-like multilayered pentacene films, grown on single crystal-like faceted islands in the first layer, have shown much higher field-effect mobility than those grown on polycrystalline dendritic islands. Figure 4.1.8 shows the two-dimensional GIXD patterns and tapping mode atomic force microscopy (TM-AFM) topographies for 60-nm thick pentacene films deposited on HMDS- and OTS-treated surfaces [42]. The occurrence of many reflection spots in the direction of qz (out of plane) at a given qxy (in plane) strongly suggests growth of three-dimensional islands. Specifically, diffraction patterns of 2-ML pentacene films yielded a pseudocentered rectangular unit cell (a = 5.90 ± 0.01 Å, b = 7.51 ± 0.01 Å, and γ = 89.92 ± 0.01°) with a herringbone packing and molecules tilted along the b-axis by 4° with respect to the surface normal.



(a)

μ = 3.4 ± 0.5 cm2/Vs

271

(b)

μ = 0.5 ± 0.15 cm2/Vs

1 μm

1 μm

(0, 2) (1, 2) (2, 1)

qz

qxy (Å−1)

(1, ±1)

(2, 0)

qz

qxy (Å−1)

FIGURE 4.1.8 Tapping mode AFM topographs and 2-D GIXD patterns of 60-nm-thick pentacene films on (a) HMDS- and (b) OTS-treated SiO2/Si substrates. (From Yang, H. et al., J. Am. Chem. Soc. 127, 11542, 2005. With permission.)

Despite minute differences such as a small portion of differently oriented crystals (indicated by black arrows in Figure 4.1.8b) and smaller grain size in the OTS sample, two-dimensional GIXD supports the fact that the HMDS and OTS samples have a similar vertical conducting path for top contact devices. Mobility measurements, however, show drastically different mobilities: μ = 3.4 ± 0.5 cm2V–1s–1 on HMDS-treated surface and 0.5 ± 0.15 cm2V–1s–1 on OTS-treated surface, using 60nm thick pentacene films in the top contact OTFT. This fact suggests that the minute differences in grain sizes may be important.

4.1.3.3 OLIGO ACENE-THIOPHENE Frisbie et al. have reported the structural and electrical characterization of two new p-channel organic semiconductors: 5,5′-bis(2-tetracenyl)-2,2′-bithiophene (1) and 5,5′-bis(2-anthracenyl)-2,2′-bithiophene (2) [43]. GIXD was performed to determine the thin-film unit cell parameters of 1 and 2. Figure 4.1.9 shows the GIXD patterns (intensity vs. qxy) for 1 and 2 from which the in-plane lattice parameters (a,b) and the angle between them (γ) were determined. Films of 1 exhibited two slightly different phases as evidenced by the occurrence of several peaks in the diffraction patterns. Using the first peak of each pair of doublets, the peaks are indexed to a rectangular in-plane unit cell with dimensions a = 5.99 Å, b = 7.66 Å, and γ = 90.0°. For 2, a similar in-plane unit cell was found with a = 5.98 Å, b = 7.64 Å, and γ = 90.0°. There are two molecules per unit cell similar to the pentacene bulk crystal structure [35]. The calculated densities of 1 and 2 are 1.41 g/cm3 and 1.39 g/cm3, respectively. These density values are comparable to other common organic semiconductors such as pentacene (1.36 g/cm3) and sexithiophene (1.55 g/cm3) [35,44]. Several of the peaks seen in the spectrum of 2 (denoted by arrows in Figure 4.1.9b) are not compatible with this lattice parameter assignment and appear to be a second phase. The peaks at qxy = 0.698, 0.932, 1.396, and 1.864 Å–1 (denoted by arrows) are weaker


272


Log intensity (arb. units)

(11) (02) (12) (20)

(23)

(21)

0.5

1.0

1.5

2.0

(14)

(22)

2.5

3.0

3.5

qxy (Å−1) (a)

Log intensity (arb. units)

(11)

(02) (12) (20) (22)

(23) (14)

(21)

0.5

1.0

1.5

2.0

2.5

3.0

3.5

qxy (Å−1) (b)

FIGURE 4.1.9 GIXD patterns for (a) a 350 Å film of 5,5′-bis(2-tetracenyl)-2,2′-bithiphene on PS-SiO2 and (b) a 350 Å film of 5,5′-bis(2-anthracenyl)-2,2′-bithiphene on bare SiO2. (From Merlo, J.A. et al., J. Am. Chem. Soc. 127, 3997, 2005. With permission.)

in intensity and correspond to d-spacings of 9.00, 6.74, 4.50, and 3.37 Å, which correspond to the values of (003), (004), (006), and (008) reflections as obtained from thin-film XRD. Therefore, these peaks suggest that some fraction of the molecules is lying down (i.e., long axis of the molecule parallel to the surface). For thicker films (~750 Å) of 2, the (110) peak was observed in thin-film XRD spectra, which supports this hypothesis. The fact that a fraction of the molecules are lying down is likely to have a detrimental effect on the electronic transport properties of films of 2, since the preferred transport direction occurs within the ab plane. With a full set of qxy and qz data, the complete thin-film unit cell can be determined. Both films have triclinic cells. The d001-spacing calculated from these unit cells was 31.83 Å for 1 and 27.03 Å for 2, which is in good agreement with the values from wideangle XRD (31.83 and 26.97 Å, respectively).

4.1.4 CONCLUDING REMARKS Over the past decade, various aspects of organic semiconductor thin films have been investigated, such as film processing parameters and dielectric surface chemistry.



273

Recent work has been focused on understanding the assembly of semiconducting materials in thin films. Understanding the bulk crystalline structure of these materials has gradually expanded toward ultrathin film architectures — the so-called thin film phase — near the dielectric. For these semiconducting thin films, GIXD profiles displayed by both the horizontal and vertical scattering summation of the crystals distributed in the films can provide a deep insight into the state of the crystal packing and orientation in the films with the thickness ranging from the monolayer in direct contact with the substrate through the subsequent multilayer. Monitoring the in situ crystallization of organic semiconductors through combining vacuum evaporator with GIXD setup can provide more detailed information for the crystal nucleation and vertical growth mechanism of the molecules on the substrates under various processing parameters. In semiconducting thin films, determination of crystalline structures with nearly atomic resolution can be achieved in favorable cases by a combination of GIXD experimental and computational techniques. While the ultimate analysis will usually involve the more complicated methods, the simpler considerations described here are useful in offering a general understanding and in providing suitable initial conditions for least-squares fittings and other parameters.

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31. Lin, Y.-Y., Gundlach, D.J., Nelson, S.F., and Jackson, T.N., Stacked pentacene layer organic thin film transistors, IEEE Electron Device Lett. 18, 606, 1997; Dimitrakopoulos, C.D. and Malenfant, P.R.L., Organic thin film transistors for large area electronics, Adv. Mater. 14, 99, 2002; Karl, N., Charge carrier transport in organic semiconductors, Synth. Met. 133, 649, 2003; Horowitz, G., Organic field-effect transistors, Adv. Mater. 10, 365, 1998. 32. Baude, P.F., Ender, D.A., Haase M.A. et al., Pentacene-based radio-frequency identification circuitry, Appl. Phys. Lett. 82, 3964, 2003. 33. Cornil, J., Calbert, J.P., and Brédas, J.L., Electronic structure of the pentacene single crystal: Relation to transport properties, J. Am. Chem. Soc, 123, 1250, 2001; Cheng, Y.C., Silbey, R.J., da Silva, D.A. et al., Three-dimensional band structure and bandlike mobility in oligoacene single crystals: A theoretical investigation J. Chem. Phys. 118, 3764, 2003. 34. Fritz, S.E., Martin, S.M., Frisbie, C.D. et al., Structural characterization of a pentacene monolayer on an amorphous SIO2 substrate with grazing incidence x-ray diffraction, J. Am. Chem. Soc. 126, 4084, 2004. 35. Holmes, D., Kumaraswamy, S., Matzger, A.J., and Vollhardt, K.P.C., On the nature of nonplanarity in the [N]phenylenes, Chem. Eur. J. 5, 3399, 1999; Mattheus, C.C., Dros, A.B., Baas, J., Meetsma, A., de Boer, J.L., and Palstra, T.T.M., Polymorphism in pentacene, Acta Crystallogr. C57, 939, 2001. 36. Mattheus, C.C., Dros, A.B., Jakob, B. et al., Identification of polymorphs of pentacene, Synth. Met. 138, 475, 2003. 37. Minakata, T., Imai, Ozaki, M., and Saco, K., Structural studies on highly ordered and highly conductive thin films of pentacene, J. Appl. Phys. 72, 5220, 1992. 38. Bouchoms, I.P.M., Schoonveld, W.A., Vrijmoeth J., and Klapwijk, T.M., Morphology identification of the thin film phases of vacuum evaporated pentacene on SiO2, Synth. Met. 104, 175, 1999; Knipp, D., Street, R.A., Volkel, A., and Ho, J., Pentacene thin film transistors on inorganic dielectrics: Morphology, structural properties, and electronic transport, J. Appl. Phys. 93, 347, 2003. 39. Kuzmenko, I., Rapaport, H., Kjaer, K. et al., Design and characterization of crystalline thin film architectures at the air-liquid interface: Simplicity to complexity, Chem. Rev. 101, 1659, 2001. 40. Kelley, T., Boardman, L.D., Dunbar, T.D. et al., High-performance OTFTs using surface-modified alumina dielectrics, J. Phys. Chem. B 107, 5877, 2003. 41. Ruiz, R., Choudhary, D., Nickel, B. et al., Pentacene thin film growth, Chem. Mater. 16, 4497, 2004; Laquindanum, J.G., Katz, H. E., Dodabalapur, A., and Lovinger, A. J., n-Channel organic transistor materials based on naphthalene frameworks, J. Am. Chem. Soc. 118, 11331, 1996; Fritz, S.E., Kelley, T.W., and Frisbie, C.D., Effect of dielectric roughness on performance of pentacene TFTs and restoration of performance with a polymeric smoothing layer, J. Phys. Chem. B 109, 10574, 2005; Nickel, B., Barabash, R., Ruiz, R. et al., Dislocation arrangements in pentacene thin films, Phys. Rev. B 70, 125401, 2004; Verlaak, S., Steudel, S., Heremans, P., Janssen, D., and Deleuze, M., Nucleation of organic semiconductors on inert substrates, Phys. Rev. B 68, 195409, 2003; Heringdorf, F.J.M.Z., Reuter, M.C., and Tromp, R.M., Growth dynamics of pentacene thin films, Nature 412, 517, 2001; Kasaya, M., Tabata, H., and Kawai, T., Scanning tunneling microscopy and molecular orbital calculation of pentacene molecules adsorbed on the Si(100)2 × 1 surface, Surf. Sci. 400, 367, 1998; Pratontep, S., Brinkmann, M., Nüesch, F., and Zuppiroli, L., Nucleation and growth of ultrathin pentacene films on silicon dioxide: Effect of deposition rate and substrate temperature, Synth. Met. 146, 387, 2004; Lim, S.C. and Kim, S.H., Surfacetreatment effects on organic thin-film transistors, Synth. Met. 148, 75, 2005.


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4.2

Near-Edge X-Ray Absorption Fine Structure (NEXAFS) Spectroscopy

Dean M. DeLongchamp, Eric K. Lin, and Daniel A. Fischer CONTENTS 4.2.1 Introduction................................................................................................277 4.2.1.1 The Importance of Structure .......................................................277 4.2.1.2 NEXAFS Background.................................................................280 4.2.1.3 NEXAFS for Organic Electronics ..............................................282 4.2.2 Experimental Considerations.....................................................................284 4.2.3 Data Analysis for Orientation....................................................................287 4.2.4 Examples of Applied NEXAFS Spectroscopy..........................................289 4.2.4.1 Pentacene.....................................................................................289 4.2.4.2 Poly(3-Hexyl Thiophene)............................................................292 4.2.4.3 NEXAFS of Oriented Liquid Crystalline Polymers...................294 4.2.4.4 NEXAFS for Molecular Electronics...........................................295 4.2.5 Future Horizons for NEXAFS Spectroscopy............................................295 References..............................................................................................................296

4.2.1 INTRODUCTION 4.2.1.1 THE IMPORTANCE

OF

STRUCTURE

Organic electronics components may soon enable the low-cost, high-volume manufacture of electronic circuitry to support emerging applications such as flexible displays, radio frequency identification (RFID) tags, and plastic photovoltaics. Highperformance organic semiconductors are required to realize these possibilities, and they have proven to be the most difficult component of organic circuitry to develop and optimize. A prime challenge of organic semiconductor development has been to understand the development of their crystalline microstructure during film

277


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formation and the impact of that microstructure on the electronic properties of the film. This understanding would enable the intelligent design of molecules and processing strategies to produce organic semiconducting layers with robust and predictable performance for targeted applications. Organic semiconductor molecules typically feature a primary chemical structure with extended conjugation. These molecules often crystallize well but lack covalent bonds between molecular repeats, and their crystals are held together by weak van der Waals forces. The strength and geometry of these intermolecular interactions determine the degree of electronic overlap and therefore the efficiency of carrier transport between molecules within a crystal. The anisotropy of molecular packing also renders the carrier mobility within a single crystal strongly anisotropic. Furthermore, practical organic thin-film transistors (OTFTs) have channels with micronscale dimensions that often appear to be filled with many hundreds or thousands of semiconductor crystals. The orientation of crystals within the transistor channel with respect to each other and to the substrate plane will influence the field effect mobility that can be achieved in a particular OTFT. Grain boundaries, gaps, and amorphous regions will also influence carrier mobility. The size, orientation, and distribution of organic semiconductor crystals within a thin film is typically strongly dependent on the purity and processing of the semiconductor, and differences in these parameters are largely responsible for the widespread, general increase in reported organic semiconductor performance, even for materials with the same nominal primary chemical structure. The conjugated network of an organic semiconductor molecule is its most important structural motif because it is within this network that carriers reside and travel. Within this conjugated network, the extent of inter- and intramolecular overlap of the π-orbital electron clouds correlated to a carrier band is the primary concern of most structural characterization methods. At the molecular level, there are essentially three types of structure considerations: (1) the orientation of molecular conjugated planes with respect to each other; (2) the orientation of molecular conjugated planes with respect to the device plane; and (3) the orientation of the conjugated long axis for molecules, such as polymers, with significantly long axes. The first structure consideration concerns crystal packing style. The packing of small molecule organic semiconductors typically falls between a “herringbone” edge-to-face interaction [1–3] and cofacial π-stacking [4–8]. Cofacial π-stacking is considered to be the superior motif because the increased physical overlap of the π-orbitals may lead to increased electronic overlap and more efficient carrier hopping. Even so, some of the highest mobility small molecule semiconductors have been of the herringbone variety, most notably pentacene [9–11] and rubrene [12–14]. Synthetic control over this packing style and the purposeful enforcement of cofacial π-stacking has become a focus of primary chemical structure studies in organic semiconductors, with some recent successes based on the addition of bulky side groups that block the edge-to-face interaction [4–8]. Notably, semiconducting polymers with reasonable performance such as the regioregular poly(3alkyl thiophenes) (P3ATs) are believed to pack naturally in a cofacial π-stacked manner that is enforced by the segregation of their side chains into aliphatic lamellae [15,16].


Near-Edge X-Ray Absorption Fine Structure

279

The second structural consideration concerns the alignment of packed crystals with respect to the device plane that can be induced by chemical interactions, anisotropic crystal geometry, or specific processing steps. Single crystals typically exhibit anisotropic mobility [12–14,17–19], and it is believed that the highest mobility within a single crystal is in the direction of the greatest electronic overlap of the conjugated planes of individual molecules. Thus, for device geometries of OTFTs, where carriers move horizontally across the device plane, the preferential orientation of conjugated planes is edge-on with respect to the substrate. This orientation allows π-overlap from side to side so that carriers may travel more efficiently in the device plane. For the device geometries of organic photovoltaic (OPV) cells, where large exciton diffusion lengths through the thickness of the device plane may be advantageous, the preferential orientation of conjugated planes may be in a flat or planeon orientation (most often, the semiconductors in organic light-emitting diodes [OLEDs] are made amorphous to avoid excimer formation). The semiconductors most commonly used in specific applications appear to naturally adopt the preferred device-relative conjugated plane orientation atop common substrates. For example, pentacene and P3ATs naturally adopt an edge-on orientation atop typical dielectrics in OTFTs. The third structural consideration of long-axis orientation applies primarily to polymer semiconductors, where the orientation of the conjugated polymer long axis may result in anisotropic carrier mobility because transport along the polymer chain may be preferred to interchain hopping [20–23]. Long-axis orientation is typically imposed by processing; the methods used are borrowed from liquid crystal alignment methods and are often applied to polymer semiconductors in a liquid crystalline state. These methods may include mechanically rubbing the semiconductor film directly or casting it upon a mechanically rubbed substrate. Typically, measurable but modest increases in mobility anisotropy are observed. Long-axis orientation of electroluminescent polymers is an emerging theme in the OLED community, where it has been shown to lead to polarized light emission [24–28]. Other structural considerations exist at length scales larger than molecules. Defects such as grain boundaries, dewetting or film gaps due to materials’ incompatibilities, and chemical contamination or degradation must also be considered. These types of defects become more common when organic semiconductors are processed from a fluid, which allows bottom-up device fabrication as an alternative to photolithography. Upon application to a substrate from a fluid, a solid semiconductor layer is formed in a dynamic assembly process at the substrate interface that may include the nucleation, growth, and orientation of crystalline material. Chemistry, chemical interactions, or heat may be employed to optimize these processes. In an OTFT, the microstructure and defects that develop near the dielectric and electrode interfaces will determine the semiconducting layer’s field-effect mobility, trap depth and distribution, and injection properties. Characterizing these many aspects of microstructure is necessary to establish relationships between primary chemical structure, processing, and performance. Currently, the most commonly used methods are scanning probe microscopy techniques such as atomic force microscopy (AFM) or kelvin probe force microscopy


280


(KPFM) and crystallographic techniques such as grazing-incidence x-ray diffraction (GIXD). The microstructure information collected by these techniques can sometimes be correlated to fundamental electronic properties or device performance. AFM and KPFM provide a local probe of nanometer-scale morphology or electrical heterogeneity below the light diffraction limit, while GIXD provides a whole-film measurement of unit cell dimensions and preferential orientation for crystalline regions. We describe here a complementary technique — near-edge x-ray absorption fine structure (NEXAFS) spectroscopy, which can be used to quantify chemical composition, molecular orientation, and defects for broad classes of organic semiconductor films.

4.2.1.2 NEXAFS BACKGROUND NEXAFS was developed in the mid-1980s with a focus on the structure and chemistry of molecules physisorbed or chemisorbed at interfaces [29]. Early applications of NEXAFS were primarily for catalysis [30]. NEXAFS measures the absorption of linearly polarized soft x-rays within 30 eV of the K-shell threshold into resonant electron excitations. This x-ray absorption results in the “fine structure,” which is a collection of absorption peaks near the core shell ionization edge. The technique is element sensitive by energetic selection of the K-shell that is accessed. Most often, NEXAFS is performed on the K-edges of low-Z elements with binding energies less than 1 keV, such as carbon (285 eV), nitrogen (400 eV), oxygen (535 eV), and fluorine (685 eV) [31]. Because the NEXAFS of low-Z elements requires a tunable monochromatic incident beam in the soft x-ray range with significant flux, it must be practiced at synchrotron facilities within an ultrahigh vacuum sample environment. The excitation of K-shell (1s) electrons by soft x-rays can result in either a bound state or a continuum state. Continuum state excitations create the photoelectrons that are analyzed in the well-known x-ray photoemission spectroscopy (XPS) technique. In contrast, the bound state excitations of NEXAFS occur as either Rydberg transitions or resonant excitations of 1s electrons to unfilled (typically antibonding) molecular orbitals, which may have either π- or σ-symmetry. A potential energy diagram is shown in Figure 4.2.1(a) to illustrate these transitions. A 1s → π* transition typically occurs at energies less than the vacuum level, and a 1s → σ* transition typically occurs at energies greater than the vacuum level. Among 1s → σ* transitions, the resonance position varies systematically with σ-bond length; lower energy resonance positions are expected for longer σ-bonds. The quantized energy separation between transitions creates the fine structure of near-edge x-ray absorbance and allows precise determination of the intramolecular bonding present within complex molecules. Spectrum interpretation is best illustrated with an example. In Figure 4.2.2, the carbon K-edge NEXAFS spectrum for poly(3-hexylthiophene) (P3HT) is shown. There are four distinct peaks. The lowest energy peak, near 286 eV, is the 1s → π* peak associated with the carbon–carbon double bonds of thiophene [30] (for P3HT, they are delocalized over the backbone). The second resonance at 288 eV is a combination of carbon–hydrogen, carbon–sulfur, and Rydberg excitations [30]. The


Near-Edge X-Ray Absorption Fine Structure

281

Continuum states σ∗ Rydberg states

π

Vacuum

π∗

σ hν

1s

(a)

Single bond

σ∗

Double bond

Conjugated system

σ∗

σ∗

π∗

σ∗

σ∗

π∗

π∗ π∗

σ∗

(b)

FIGURE 4.2.1 a) Band energy diagram depicting a NEXAFS resonant excitation. Incident soft x-ray photons excite 1s electrons to unfilled molecular orbitals such as the π* or σ*. A 1s → π* transition, as shown here, can occur at energies below the ionization edge. b) Directional resonances are dependent on the spatial location of the final state orbital, and can be expressed as vectors or planes. (Adapted from Stöhr, J., NEXAFS spectroscopy, Springer–Verlag, Berlin, 1992.)

edge jump, where continuum states are accessed, appears at ≈289 eV. This edge jump causes a step change in the absorbance, so all energies above the edge jump exhibit a constant absorbance that has been normalized to 1 in Figure 4.2.2. The peak near 293 eV is associated with the 1s → π* of carbon–carbon single bonds, while the peak near 303 eV is associated with the 1s → σ* of carbon–carbon double bonds. Both peaks are present in conjugated systems [32]. This peak hierarchy follows Figure 4.2.1(a). The intensities of these peak shapes are equivalent to the stoichiometric bond densities within the sampled volume.


282


Partial electron yield

3.0

σ∗C-S/C-H

2.5 2.0 1.5

σ∗C-C σ∗C=C

π∗C=C

S

n

1.0 0.5 0.0 280

290

300

310

320

330

Photon energy (eV)

FIGURE 4.2.2 NEXAFS spectrum of poly(3-hexylthiophene) (P3HT) with labeled resonances. PEY standard experimental uncertainty is > ξ ⇒ g(R) → 2σ2. The treatment, so far, has considered static spatial variations only, but film growth is dynamic and temporal fluctuations must also be considered [30]. Family and Vicsek [48] considered the evolution of the lateral and perpendicular fluctuations in their “dynamic scaling” theory. According to dynamic scaling, the rms roughness evolves with time as σ(t) ∝ tβ if the average film height is proportional to time (〈h〉 ~ t). Combining this with the lateral fluctuations yields: ⎛ t ⎞ σ ( L, t ) = LH F ⎜ H /β ⎟ ⎝L ⎠

(5.1.21)

where F(x) → 1 as x → ∞ and σ(t) ∝ tβ as x → 0. Central to Equation 5.1.21 is a scaling of the correlation length with time ξ ∝ t1/z, where z = H/β is known as the dynamic scaling exponent. The scaling exponents have been calculated theoretically using a variety of assumptions about the ratelimiting factors [49]. The numerical estimate by Kardar–Parisi–Zhang (KPZ) ignores the effects of surface diffusion and estimates H, β, and z as 0.385, 0.24, and 1.61, respectively [49]. Other estimates, more applicable to PVD growth, incorporate diffusion and yield values of 2/3, 1/5, and 10/3, respectively [49]. Durr et al. demonstrated the first application of scaling laws to organic systems for the deposition of diindenoperylene (DIP) on SiO2 [50]. They utilized x-ray reflectivity, AFM, and diffuse x-ray scattering to analyze the temporal and spatial fluctuations of the surface roughness. To do this, they used an extended version of


364


Equation 5.1.20 that includes a steepening exponent λ: g(R) ∝ t2λR2H [50]. Fitting the scaling equations to their samples yields 1/z = 0.92 ± 0.20, H = 0.684 ± 0.05, β = 0.748 ± 0.07, and λ = 0.17 or 0.07 depending on how H was calculated. The values presented do not match the existing theories mentioned previously; instead, the behavior of the film growth is referred to as rapid roughening. Models are being developed to explain this phenomenon.

5.1.4 EFFECTS OF THE SUBSTRATE The morphology and OTFT performance of the active layer are dependent on several external parameters in addition to vacuum pressure, deposition rate, and substrate temperature. Specifically, the nature and strength of the interaction potential between the adsorbate and the substrate will greatly influence the arrangement of anisotropic molecules when deposited and even control, to some extent, whether the islands will nucleate and grow in as two or three dimensions. The surface roughness of the substrate as well as its surface tension, γ, will greatly influence morphology (as mentioned in Section 5.1.1.2) and device performance.

5.1.4.1 EFFECT

OF

SURFACE ENERGY

The effects of surface tension and the interaction potential between the substrate and pentacene are demonstrated in Figure 5.1.15. The left column shows AFM micrographs for pentacene deposited on poly(imide-siloxane) with varying amounts of siloxane in the polymer backbone as reported by Yang et al. [51]. The variation in siloxane shifts the surface energy from 48 mJm–2 (a) to 30 mJm–2 (c) without varying the surface roughness appreciably. As can be seen, the grain size decreases with surface energy as expected from the thermodynamic explanation of film growth (see Section 5.1.1.2). The authors also show that the mobility follows the opposite trend (not shown). The right column of Figure 5.1.15 shows the dependence of the molecule–substrate interaction as discussed originally in Verlaak et al. [15] and continued in Steudel et al. [52]. The value of the molecule–substrate interaction is plotted for a variety of common OTFT substrates and pentacene in Figure 5.1.15d [52]. In Figure 5.1.15e, the effects on the growth mode of pentacene films is shown for pentacene growth on OTS and on (heptadecafluoro-1,1,2,2-tetrahydrodecyl)trichlorosilane (FDTS). As shown, the lower interaction potential between pentacene and FDTS forces the growth mode to three dimensions.

5.1.4.2 EFFECT

OF

SURFACE ROUGHNESS

Several groups have explored the effects of surface roughness on OTFTs. Figure 5.1.16 shows the effects of increasing surface roughness on pentacene film morphology. The pentacene grain sizes for growth on SiO2 with an rms roughness of 0.2 nm (a) and rough SiO2 with an rms roughness of 1.5 nm (b) are dramatically different [53]. Figure 5.1.16c and d show the morphology of pentacene films grown on OTScoated SiO2 [54]. The last two panes in Figure 5.1.16 show the effects on pentacene film morphology of the surface roughness of SiN substrates (prepared by plasma


Vacuum Evaporated Thin Films

365

(a) (d)

Φ = 0.25Å/s

Pentacene

ψmol-sub [eV]

0.12 OTS, PVP

BUTS, UETS

0.08

CUTS, PTS, PETS 0.04

FDTS

0.00 (b)

300

320

340

360

Tsub [K] 10 nm

(e)

(b)

(c) 10 nm

100 μm

(a)

1 μm

(c)

FIGURE 5.1.15 Effects of the substrate’s surface energy as it decreases from (a) to (c) for pentacene deposited on poly(imide-siloxane). (Yang, S.Y. et al., Adv. Func. Mat., 15, 1806, 2005.) The interaction potential between the molecule and the substrate for various substrates (d) and the dependence on the growth mode on the magnitude of the interaction (e). The twodimensional connected regime is for pentacene grown on OTS and the three-dimensional portion is grown on FDTS. (Steudel, S. et al., Appl. Phys. Lett., 85, 5550, 2004.)

enhanced chemical vapor deposition) [55]. The standard SiN surface had an rms roughness of 0.5 nm (Figure 5.1.16f), while the SiN in Figure 5.1.16e was atomically smooth. The second group in Figure 5.1.16 shows the effects of surface roughness for pentacene deposited on OTS-coated SiO2.

5.1.5 OUTLOOK In the preceding sections we outlined some models from the inorganic film growth community, and early attempts to adapt and apply them to organics. This process is still at its infancy because the primary identifying characteristic of organics — namely, their anisotropic shape — has not been taken into account. This will require


366


Increased roughness (a)

(b)

(c)

(d)

SiO2

OTS-coated SiO2

1 μm (e)

(a) 2.5 μm

(b) (f )

2.5 μm

SiN

Sample B 70°C, 0.5 Å/s, smooth

Sample A 70°C, 0.5 Å/s, standard

FIGURE 5.1.16 Effects of substrate surface roughness on pentacene film morphology. Growth on smooth (a) and rough (b) SiO2 (Fritz, S.E. et al., J. Phys. Chem. B, 109, 10574, 2205); smooth (c) and rough (d) OTS-coated SiO2 (Steudel, S. et al., Appl. Phys. Lett., 85, 4400, 2004); and smooth (e) and rough (f) SiN. (Knipp, D. et al., Appl. Phys. Lett., 82, 3907, 2003.)

a generalization of growth models and will lead not only to better control of the morphology of organic films, but also to new and exciting growth in physics.

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24. Ruiz, R., Mayer, A.C., Malliaras, G.G., Nickel, B., Scoles, G., Kazimirov, A., Kim, H., and Headrick, R.L., Structure of pentacene thin films, Appl. Phys. Lett., 85, 4926, 2004. 25. Marra, W.C., Eisenberger, P., and Cho, A.Y., X-ray total-external-reflection–Bragg diffraction: A structural study of the GaAs-Al interface, J. Appl. Phys., 50, 6927, 1979. 26. Bauer, E., Low energy electron microscopy, Rep. Prog. Phys., 57, 895, 1994. 27. Meyer Zu Heringdorf, F.-J., Reuter, M.C., and Tromp, R.M., Growth dynamics of pentacene thin films, Nature, 412, 517, 2001. 28. Balluffi, R.W., Allen, S.M., and Carter, W.C., Kinetics of materials, Wiley Interscience, New York, 2005. 29. Cohen, P.I., Petrich, G.S., Pukite, P.R., Whaley, G.J., and Arrott, A.S., Birth–death models of epitaxy: I. Diffraction oscillations from low index surfaces, Surf. Sci., 216, 222, 1989. 30. Krim, J. and Palasantzas, G., Experimental observations of self-affine scaling and kinetic roughening at submicron length scales, Int. J. Mod. Phys. B, 9, 599, 1995. 31. Ruiz, R., Papadimitratos, A., Mayer, A.C., and Malliaras, G.G., Thickness dependence of mobility in pentacene thin-film transistors, Adv. Mater., 17, 1795, 2005. 32. Wang, G.Z., Luo, Y., and Beton, P.H., High-mobility organic transistors fabricated from single pentacene microcrystals grown on a polymer film, Appl. Phys. Lett., 83, 3108, 2003. 33. Sanvitto, D., De Seta, M., and Evangelisti, F., Growth of thin C60 films on hydrogenated Si(100) surfaces, Surf. Sci., 452, 191, 2000. 34. Park, D.S., Kang, S.J., Kim, H.J., Jang, M.H., Noh, M., Yoo, K.-H., Whang, C.N., Lee, Y.S., and Lee, M.H., Characteristics of perylene-based organic thin-film transistor with octadecyltrichlorosilane monolayer, J. Vac. Sci. Technol. B, 23, 926, 2005. 35. Möbus, M. and Karl, N., The growth of organic thin films on silicon substrates studied by x-ray reflectometry, Thin Solid Films, 215, 213, 1992. 36. Zinsmeister, G., A contribution to Frenkel’s theory of condensation, Vacuum, 16, 529, 1966; Zinsmeister, G., Theory of thin film condensation. Part B: Solution of the simplified condensation equation, Thin Solid Films, 2, 497, 1968. 37. Ehrlich, G. and Hudda, F.G., Atomic view of surface self-diffusion: Tungsten on tungsten, J. Chem. Phys., 44, 1039, 1966. 38. Schwoebel, R.L. and Chipsey, E.J., Step motion on crystal surfaces, J. Appl. Phys., 37, 3682, 1969. 39. Pratontep, S., Brinkmann, M., Nüesch, F., and Zuppiroli, L., Correlated growth in ultrathin pentacene films on silicon oxide: Effect of deposition rate, Phys. Rev. B, 69, 165201, 2004. 40. Pratontep, S., Nuesch, F., Zuppiroli, L., and Brinkmann, M., Comparison between nucleation of pentacene monolayer islands on polymeric and inorganic substrates, Phys. Rev. B, 72, 085211, 2005. 41. Amar, J.G., Family, F., and Lam, P.-M., Dynamic scaling of the island-size distribution and percolation in a model of submonolayer molecular-beam epitaxy, Phys. Rev. B, 50, 8781, 1994. 42. Stroscio, J.A. and Pierce, D.T., Scaling of diffusion-mediated island growth in ironon-iron homoepitaxy, Phys. Rev. B, 49, 8522, 1994. 43. Ruiz, R., Nickel, B., Koch, N., Feldman, L.C., Haglund, R.F., Jr., Kahn, A., Family, F., and Scoles, G., Dynamic scaling, island size distribution, and morphology in the aggregation regime of submonolayer pentacene films, Phys. Rev. Lett., 91, 136102, 2003.



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5.2

Solution Deposition of Polymers

Hoichang Yang CONTENTS 5.2.1 Introduction................................................................................................371 5.2.2 Solution Casting and Spin-Coating to Form Self-Assembled Polymers.....................................................................................................374 5.2.2.1 Molecular Structure.....................................................................374 5.2.2.2 Effect of Solvent .........................................................................382 5.2.2.3 Processing Condition...................................................................387 5.2.3 Concluding Remarks..................................................................................392 References..............................................................................................................395

5.2.1 INTRODUCTION Various organic thin film transistors (OTFTs) containing semiconducting polymers with π-conjugated backbones for driving light-emitting diodes and prototype electronic polymer circuits are presently being demonstrated [1–3]. Unlike vacuum deposition, which is more costly and requires long pump-down time, solutiondeposition methods for semiconductor thin films can be done under ambient conditions on short timescales. As a result, they are compatible with large-area thin film fabrication and result in lower production cost per device. Two approaches have been used to obtain organic semiconductor thin films from solution: deposition of a soluble precursor from solution and subsequent conversion reaction direct deposition of a soluble semiconductor An example of a conversion reaction from a precursor to a polymer is shown in Figure 5.2.1 [4]. The conversion reaction for the precursor materials required for achieving high mobility is performed at a high temperature above 125°C, making it incompatible with low-cost plastic substrates, which typically have glass transition temperatures below 125°C. The additional conversion step is also time consuming. Accordingly, direct deposition of soluble semiconductors is more desirable.

371


372


HCl

H2 CH C S OCH3 Precursor polymer

n

Δ

S

C H

C H n

Poly(2, 5-thienylenevinylene)

FIGURE 5.2.1 Conversion reaction from a precursor polymer to poly(2,5-thienylenevinyle) (PTV).

Electrochemical polymerization has been used to synthesize and deposit polythiophenes directly into device channels [5,6]. In this case, deposited films tend to show low on/off ratios and field-effect mobility. Another interesting example involves solution deposition of an n-type ladder polymer using a proper Lewis acid acting as a chelating and a solubilizing group. After film fabrication, the Lewis acid is removed by extensive washing [7,8], resulting in highly ordered polymer thin films [8]. However, the remaining trace amount of Lewis acid yields unusual electrical characteristics and has to be completely removed in order to obtain high mobility. Accordingly, direct deposition of soluble semiconductors is more desirable. Since most semiconducting polymers with π–π-conjugated backbones are insoluble in common organic solvents, side chains have been incorporated into the molecular structure through various chemical substitutions [9–14]. For example, remarkably good solubility has been achieved for polythiophenes by substituting alkyl side chains at the 3-position of each thiophene ring because these side groups mitigate the strong π–π-interchain interaction between the polythiophene chains. Until now, various poly(3-alkylthiophene)s (PATs) [15–18], polythiophenes with end-functionalized side chains [14,19], block copolymers [20–22], combined with these polymers have been used for solution processable OTFT applications. Among these polymers, regioregular head-to-tail poly(3-hexylthiophene) (HT PHT) films have shown higher charge mobility (0.01 ~ 0.1 cm2V–1S–1) and reasonable on/off ratios (>100 in air and 106 in an inert atmosphere) in OTFTs [17,18,23,24]. The Langmuir–Blodgett (LB) technique [25,26] is another possible method to fabricate organic semiconductor thin films [27–33]. The molecular density at the air–water interface is manipulated with movable barriers in an LB film trough. By moving the barriers, the film density increases (compression) or decreases (expansion), which changes the surface tension between liquid and vapor phases, γlv, at the air–water interface. The film can then be transferred onto a substrate either as a monolayer or as multilayers. The molecules in the LB film are usually well aligned, but PHT is not amphiphilic, and it was difficult to obtain highly aligned mono- or multilayers on an air–water interface. Surfactant molecules, as well as nonamphiphilic polymers decorated with alkyl chains, can be incorporated to improve the film ordering, but at the same time these insulating molecules disrupted charge transport [29]. Conversely, amphiphilic polythiophenes [32–34] can self-assemble into π-stacked conjugated chains that form a very stable monolayer with a local structure optimized for high electrical conductivity. Generally, polymeric semiconductor thin films for OTFT applications can be fabricated through various direct solution-processing methods: spin-casting [15,17,18,35,36], drop-casting [15,23], screen printing [37,38], and ink jet printing



373

[39–42], which allow for a single-step deposition and patterning of the active materials [43]. Both spin-casting and drop-casting are commonly used for the film deposition. The spin-casting method yields fast solvent evaporation, allowing less time for molecular ordering but more uniform films compared to drop-casting. These solution-deposited films can be annealed [15,44–48] or melt crystallized [24] afterwards to improve molecular ordering. Electrical performance of the OTFT devices is affected by the microstrucuture in these semiconductor thin films [17,23,24,35,36,44]. It has been recognized that intrachain and interchain (π-conjugated plane) transport takes place in polymer films, whereas transport in amorphous portion and between grain boundaries has a high energy barrier [49,50]. Therefore, it is important to achieve greater control of the solution deposition process that takes these polymers from a soluble state into a polycrystalline thin film. Specifically, the desire to make nanoscale devices for future electronic applications further emphasizes a preferential orientation of π-conjugated crystal planes as the main hole-transporting path of p-type semiconductors. Understanding the crystal growth mechanisms of these solution-processable polymers is required to control the self-assembled structure during film formation. Most studies on solution-deposited semiconducting polymer thin films for OTFTs pertain to PATs [6,10,15,23,24,27,30,35,36,44,46–48,51–81]. The key factors affecting self-assembly of these polymers during solution processing are as follows: molecular properties: π-conjugated polymer backbone, side chain type, regioregularity, number average molecular weight ( M n , MW), molecular weight distribution ( M w / M n , PDI) solvent nature: volatility (boiling point), solubility of the polymer processing condition: substrate treatment, solvent evaporation rate, thermal annealing (Figures 5.2.2 and 5.2.3), and nature of substrate surface Mesoscale crystalline morphology, crystallinity, and molecular orientation in these deposited thin films strongly depend on molecular properties [17,18], chemical nature of the solvent, and processing condition, resulting in very different field-effect mobilities [15,23,36]. Specifically, due to heterogeneous surface-induced (epitaxy) crystal growth as a nature of semicrystalline polymers, fine control of substrate properties and solvent evaporation rate tends to yield favorable molecular orientation of these polymers (i.e., edge-on structure with respect to dielectric substrates) in solution-deposited films [24,66]. In this chapter, we will discuss the formation of various mesoscale crystalline morphologies and molecular orientation of HT-PATs, specifically HT-PHTs in solution-deposited films. Also, we will introduce poly(3,3″′-dialkylquaterthiophene) [82,83] and poly(2,5-bis(3-alkylthiophene-2-yl)thiono[3,2-b]thiophene [84], which can grow into highly oriented crystallites with respect to the substrates and show high charge mobilities of 0.2 ~ 0.6 cm2V–1s–1 achieved under nitrogen. Finally, we will discuss nucleation and growth of these crystalline polymers during solution deposition and present various approaches to control mesoscale and nanoscale formation of these self-assemblies.



Meltcrystallization Annealing (at 210) (180°C, 12h)

Spin-cast

No treatment

Drop-cast

374

(a) CHCl3

(b) Toluene

(c) THF

(d) CH2Cl2

FIGURE 5.2.2 Tapping mode atomic force microscopy topographs of poly(3-octyl thiophene) thin films cast on SiO2/Si substrates with various solution processing techniques from different solvents: (a) chloroform (CHCl3); (b) toluene; (c) tetrahydrofuran (THF); (d) methylene chloride (CH2Cl2). (From Yang, H., unpublished data, 2006. With permission.)

5.2.2 SOLUTION CASTING AND SPIN-COATING TO FORM SELF-ASSEMBLED POLYMERS 5.2.2.1 MOLECULAR STRUCTURE In order to improve solubility of unsubstituted polythiophenes, which are insoluble in organic solvents, various side chains have been incorporated into these backbones. In this case, these substituted side chains may interfere with molecular ordering or increase cell dimensions. Specifically, side-alkyl stacking (d(h00)) and π–π-stacking (d(010)) distance in crystals tends to increase with an increase in side-chain length, resulting in lower crystallinity and charge-carrier mobilities [51–55,58,74,77,78,85]. Figure 5.2.4 illustrates typical two-dimensional grazing incidence x-ray diffraction (GIXD; a detailed description is given in chapter 4.1) patterns of HT-PAT films dropcast from methylene chloride in closed jars. In all drop-cast HT-PAT films, these molecules tend to maintain preferentially an edge-on structure with respect to the SiO2/Si substrates, regardless of the alkyl side chains: hexyl (–C6H13), octyl (–C8H17), and dodecyl (–C12H25). The values of d(100) and d(010) increase from 16.1 to 25.5 Å and from 3.74 to 3.90 Å, respectively, with an increase in alkyl chain length from hexyl to dodecyl [51,77]. Semiconducting polythiophene derivatives have been designed to assemble into large crystalline domains on crystallization and to possess an extended, planar


375

Meltcrystallization Annealing (at 210) (180°C, 12h)

Spin-cast

No treatment

Drop-cast


(a) CHCl3

(b) Toluene

(c) THF

(d) CH2Cl2

(a)

1.0

1.5

(b)

(200) (100) 0.0 0.5 1.0 1.5 qxy(Å−1)

(c) (400)

(300)

0.0

0.5

qz (Å−1)

2.0

FIGURE 5.2.3 Two-dimensional grazing incidence x-ray diffraction patterns of poly(3-octyl thiophene) thin films cast on SiO2/Si substrates with various solution processing techniques from different solvents: (a) chloroform (CHCl3); (b) toluene; (c) tetrahydrofuran (THF); (d) methylene chloride (CH2Cl2). (From Yang, H., unpublished data, 2006. With permission.)

(010)

(300)

(010)

(200) (100) 0.0 0.5 1.0 1.5 qxy(Å−1)

(500) (400) (300) (200) (100)

(010)

0.0 0.5 1.0 1.5 qxy(Å−1)

FIGURE 5.2.4 Two-dimensional GIXD patterns of drop-cast HT-PAT films with different alkyl side-chain lengths: (a) PHT; (b) poly(3-octylthiophene) (POT); (c) poly(3-dodecyl thiophene) (PDDT). (From Yang, H. unpublished data, 2006. With permission.)


376


C12H25 S

S

* S

* n

S C12H25 PQT-12

(a)

(b)

(c)

(d)

(e)

(f )

10−15 nm m 0n

−8

30

(100)

30−80 nm 10−15 nm (010)

(001)

FIGURE 5.2.5 Top shows the chemical structure of poly(3,3′″-didodecylquaterthiophene)s (PQT-12). (a) Topography and (b) phase images (1 × 1 μm) of annealed PQT-12 thin film on unmodified SiO2 surface; (c) and (d) are, respectively, topography and phase images (1 × 1 μm) of annealed PQT-12 thin film on OTS-8-modified SiO2 surface; (e) blown-up image from (d) showing rodlike structures with domains; (f) schematic depiction of a PQT-12 lamellar π–π stack. (Adapted from Wu, Y.O. et al., Appl. Phys. Lett. 86 (14), 2005.)

π-electron system that allows close intermolecular π–π-stacking and facilitates high charge mobility. Ong and coworkers have reported poly(3,3′″-didodecylquaterthiophene) (PQT-12) (Figure 5.2.5) thin films with highly edge-on orientation of polythiophene molecules, which facilitates in-plane charge transport and leads to high charge carrier mobility [16,83]. PQT-12 shows an inverse comb-like liquidcrystalline phase with two melting transitions at ~120 and 140°C, upon heating, and two crystallization transitions at ~118 and 70°C, upon cooling. Specifically, thermal annealing for spin-cast PQT-12 films near the liquid-crystalline phase transition temperatures induces highly oriented crystalline nanofibrils with respect to



377

the substrate, resulting in high charge mobility (~0.2 cm2V–1s–1) in OTFT devices (Figure 5.2.5). Recently, poly(2,5-bis(3-alkylthiophene-2-yl)thioeno[3,2-b]thiophenes (PBTTT) (Figure 5.2.6) incorporating a linear fused conjugated unit, thieno[3,2b]thiophene with alkylthiophenes, were designed to have reduced electron delocalization along the polymer backbone. This results in a lowering of the polymer highest occupied molecular orbital level. Furthermore, the rotational invariance of the symmetrical thieno[3,2-b]thiophene in the backbone facilitates the adoption of the lowenergy chain conformation, promoting formation of highly ordered crystalline domains. PBTTT can grow into two-dimensional crystallites through postannealing spin-cast films achieving high charge mobilities of 0.2 ~ 0.6 cm2V–1s–1 [84]. Solution-deposited films of HT-PATs show various mesoscale crystalline morphologies from an amorphous texture to nanofibrils, mainly depending on their MW, solvents used, and processing conditions [15,17,23,35,36,44,59]. HT regioregularity of substituted side chains in polythiophene backbones strongly affects the solubility and molecular ordering of the resulting polymers. As the HT regioregularity increases, the solubility of PHT decreases and its crystallinity increases. However, high HT-PHT shows low crystallinity of ~15% [23,86] compared to polythiophenes (~35%) without side chains, as calculated from differential scanning calorimetry and x-ray analyses [86,87]. Sirringhaus and coworkers have found that spin-cast PHT films with low (81%) and high (96%) HT structure adapt face-on and edge-on structure of molecules on SiO2 substrates, respectively, as seen in Figure 5.2.7 [17,18]. As a result, field-effect mobility (μ) measured for the high HT-PHT film was 0.05–0.1 cm2V–1s–1, two orders of magnitude higher than the 2 × 10–4 cm2V–1s–1 of the low HT-PHT film (Figure 5.2.7c). Interestingly, drop-casting of low HT-PHT results in an order of magnitude improvement of the mobility compared to spin-casting. This is mainly related to crystallization behaviors of HT-PHT dependent on solvent evaporation rate, which will be discussed in more detail later. It is well known that solution-deposited HT-PAT films form various mesoscale crystalline morphologies according to the nature of the solvent [23,36,44,88] and solvent evaporation rate [89]. Unlike other π-conjugated polymers with rigid-rod conformation by strong π–π-interaction even in organic solvents [90], for a “flexible coil-like” HT-PAT molecule in a solution [63], self-organization of the polymer into an ordered crystal should be related to MW. This relationship is expected because PAT is a semicrystalline polymer. MW dependence is especially evident in the viscosity of semicrystalline polymer solutions or melts where the polymer MW strongly affects homogeneous nucleation and growth rate of the polymers, resulting in different mesoscale crystalline morphologies [63,91–95]. Accordingly, the influence of MW on the crystallization behaviors of semicrystalline polymers has been studied in various articles. For example, linear crystal growth rates of poly(ethylene oxide) and poly(ethylene succinate) (PES) reach a minimum value at a critical MW. This value is related to the crystallization transition from an extended chain to a folded chain conformation [96,97], suggesting that high MW polymers require sufficient reconformation time to achieve an ordered structure. As evidence of this MW dependence of the semicrystalline polymer on


378


R S *

S

S S

* n

R PBTTT (a)

400 nm (b)

400 nm

FIGURE 5.2.6 Top shows chemical structure of poly(2,5-bis(3-alkylthiophen-2yl)thieno[3,2b]thiophene) (PBTTT), where R = C10H21, C12H25, or C14H29. AFM images of PBTTT with side chain of C12H25 (annealed at 180°C): (a) before and (b) after annealing above the liquidcrystal isotherm. The left images show the topography and the right images show the phase images. The dark spots in the topography of the annealed film are voids where the film partially dewetted during annealing. (Adapted from McCulloch, I. et al., Nat. Mater. 5 (4), 328–333, 2006.)



379

(b) a a

b

Intensity (a.u.)

5 b

(a)

4.5 4 3.5 3

(010) (300)

1

(100) (100)

(010)

μsat (cm2 V−1 s−1)

2 1.5

(200)

10−1

2.5

0.5

(c)

10−2 10−3 10−4

70

80 90 % heat-to-tall

FIGURE 5.2.7 Two-dimensional GIXD patterns of spin-cast PHT films with head-to-tail regioregularity of 96% (a) and 81% (b), showing different molecular orientations with respect to the SiO2 substrate. The film thickness was 70–100 nm. (c) Dependence of the roomtemperature mobility on the regioregularity for spin-cast (downward triangles) and drop-cast (upward triangles) top-contact PHT FETs (channel length L = 75 μm, channel width W = 1.5 mm). (From Sirringhaus, H., Nature, 401(6754), 685, 1999. With permission.)

crystallization transition, Magonov and coworkers reported AFM images of the solution-grown single crystals using C60H122 (M n ~ 850 Da) and C390H782 (M n ~ 5.5 kDa): the low MW C60H122 was extended in the crystals, while the high MW C390H782 showed a folded structure [92,93]. For HT PHT crystal on a highly oriented pyrolytic graphite substrate, scanning tunneling microscopy (STM) images [61,62,98–100] have showed that hexyl side chains tended to be organized in a planar, zigzag, interdigitated fashion, and the πconjugated backbones were oriented perpendicular to these side chains [98,99]. As seen in Figure 5.2.8, the STM images indicate that the conjugated chains tend to adopt both trans-linear conformation of the thiophene units and cis-conformation, which is a prerequisite for a “hair-pin” chain folding. MW and PDI of semiconducting polymers have become a popular topic for many research groups [35,44,59,60,88,101]. Commercially available HT-PHTs synthesized by metal-catalyzed cross-coupling polymerization have a broad PDI (>2.0),


380


(a)

(c)

(b)

A A

a B

a

B

100 A

20 A

FIGURE 5.2.8 STM images of long-range (a) and short-range (b) ordering of HT poly(3dodecyl thiophene) (PDT) on highly oriented pyrolytic graphites (HOPGs). A closer look at the images reveals that the individual strands are linear and chain folding, in which cisconformations of the thiophene units are prerequisite for the fold is also clearly observed. Regular hair-pin folds angle can be seen (A) but also folds of 120° (B). (c) Calculated model of PDDT corresponding to the area enclosed in the white square on the STM-image in (b). (Mena-Osteritz, E. et al., Angew. Chem. Int. Ed., 39(15), 2680, 2000. With permission.)

a small portion of regiorandom PHT, and other impurities. Of course, several HTPHT fractions showing narrow MWs and high regioregularity can be physically separated from the HT-PHTs using simple methods, such as a Soxhlet extraction with various solvents [35,70]. McGehee and coworkers have reported that the mobility of spin-cast HT-PHT films differs by several orders of magnitude depending on MW [35,44,102]. Low MW-PHTs (M n < 4 kDa) were found to form a highly crystalline “rod-like” morphology, while high MW polymers (M n > 30 kDa) formed isotropic nodules with low crystallinity. Recently, they also found that, for low MW PHTs, the crystal orientation at a critical buried interface between the film and the dielectric plays an important role on field-effect mobility, using x-ray diffraction rocking analysis for the low MW-PHT films on different self-assembled monolayer (SAM)-treated substrates. They observed a 1,000-fold increase in mobility of the film with highly oriented crystals [102]. More detailed discussions about their analysis with GIXD are given in chapter 4.1. Generally, the charge transport properties are dependent on the charge carrier density [103]. Charge-carrier transport of p-type semiconducting polymers in solid states (bulk or film) is classified into three different categories: intrachain, interchain, and interfibrillar hopping [49]. In crystals with ordered backbone chains, charge carriers can in theory transport along the polymer up to its contour length [50]. The carriers also transport between π-conjugated planes in the crystals. Finally, the charge carrier can transfer across crystalline grain boundaries (GBs) through interfibrillar hopping. For spin-cast HT-P3HT films, McGehee and coworkers proposed that the high MW molecules form small ordered areas separated by disordered regions [44]. In this case, the long chains can interconnect ordered areas and prevent charge carriers from being trapped by the disordered GBs. As a result, spin-cast high MWHT-P3HT films with fewer ordered crystals and GBs show higher charge mobility than low MW films. This results in highly developed crystalline structures, which require high potential energy to transfer the charge carriers across the GBs. However, it is unclear whether the difference in mobility is due to the sharper grain interface



(a)

381

(b)

100 nm

(c)

100 nm

(d)

500 nm

(e)

500 nm (g)

(f)

500 nm (h)

500 nm

500 nm (i)

500 nm

500 nm

FIGURE 5.2.9 TM-AFM phase images of thin films of HT-PHTs of various weight-average molecular weight ( M w ) in FET devices prepared by drop-casting from toluene. M w in (a) through (i), respectively: 2.4, 4.8, 5.1, 7.0, 11.8, 15.7, 17.3, and 18.4 kDa. (Zhang, R. et al., J. Am. Chem. Soc., 128(11), 3480, 2006. With permission.)

between rods in the low MW films or to the fact that longer chains require less hopping between molecules. Recently, for HT-PHT films drop-cast from toluene, McCullough and coworkers demonstrated a change in nanofibrillar dimensions, specifically, the lateral width of the HT-PHT nanofibrils, dependent on weight-average molecular weight (M n) [104]. Figure 5.2.9 illustrates the TM-AFM phase images of drop-cast HT-PHT films on SiO2 substrates. In all samples, HT-PHTs tend to yield nanofibrils, regardless of MW, but the lateral width of the nanofibrils (as determined by Fourier transform of AFM images and grazing-incidence small-angle x-ray scattering) initially increases linearly with M w and then levels off, as shown in Figure 5.2.10. Also, field-effect mobility of HT-PHT films increases with an increase in the lateral width of the nanofibrils with respect to M w . Specifically, below M w < 10 kDa, a direct relationship between the nanofibrillar width and the contour length (Lw) [105,106] suggests that the polymer backbone can be extended up to the Lws and aligned perpendicular


382


Lw-weight average contour length (nm)

30

10

20

30

40 10−2

Width Mobility

20

10−3 10−4

10

Mobility (cm2/Vs)

WAFM-Nanofibrill width (nm)

0

10−5 10−6 5 0 10 15 20 Mw-weight average molecular weight (kDa)

0

FIGURE 5.2.10 Dependence of nanofibril width (WAFM) and field-effect mobility (μ) on weight-average molecular weight (Mw, bottom axis) and weight-average contour length (Lw, top axis) of HT-PHT. (Zhang, R. et al., J. Am. Chem. Soc., 128(11), 3480, 2006. With permission.)

to the nanofibril axis. The increase in lateral width of nanofibrils results in improvement of field-effect mobility of the PHT-based OTFT devices (Figure 5.2.10). Ordered π–π-conjugated structure of HT-PAT polymers correlates strongly with their excellent electrical OTFT performance. However, most semiconducting polymers still have poor mechanical and processing properties compared to other nonconjugated polymers. One approach to solve this problem is to combine these polymers with other polymer units (i.e., to synthesize a block copolymer or to prepare blends) [53]. Block copolymers [107] can self-assemble into a number of nanoscale morphologies, such as lamellar, spherical, cylindrical, and so on. McCullough and coworkers have tried to control the self-assembled structure of thiophene-based polymers chemically linked with other polymers through changing the weight fractions of the HT-PHT component [9,22]. However, these polymer films show much lower conductivity (30 2.3 2.1 0.4 0.015 Insoluble

Source: Yamamoto, T. et al., J. Am. Chem. Soc. 120, 2047–2058, 1998.

only two to three times higher than that of common flexible polymers such as polystyrene or poly(methyl methacrylate), despite its conjugated backbone [63,64]. The magnitude of the Huggins coefficient (κH) is reasonable for a flexible polymer in a “good” solvent. However, in dilute CHCl3 and THF solutions of HTPHT, the κH of HT PHT increases as a function of MW [63]. Consequently, CHCl3 and THF are not “good” but “moderately good” solvents for HT-PHT, and there exists a molecular limit of HT-PHT at which the molecule is no longer soluble at room temperature. Specifically, light-scattering analysis for soluble HT-PHT in CHCl3 shows that photoinduced polarizability of regioregular and regiorandom PHTs is 0.26 and 0.01, respectively, indicating that HT-PHT even has a stiff structure at dilute concentrations [79]. As a solvent is evaporated in a solution containing soluble HT-PHT, the solution concentration becomes higher, causing a solventinduced aggregation or crystallization of the polymer due to solubility differences


384


520 a

Absorbance

0.8

560 ih

b c

0.6 0.4

Methanol content (vol %) a 0 g b 10 fe c 20 d d 30 e 40 610 f 50 g 60 h 70 i 80

0.2 0 200

1.0 0.8 Absorbance

1.0

e f g h i

0.6

ab c d

0.4

Methanol content (vol %) a 0 b 10 c 20 d 30 e 40 f 50 g 60 h 70 i 80

0.2 300

400 500 600 Wavelength/nm (a)

700

800

0 200

300

400 500 600 Wavelength/nm

700

800

(b)

FIGURE 5.2.12 Changes in UV-vis spectra of CHCl3 solutions of (a) HT-PHT and (b) mixed PHT (HT/HH = 1/2) on addition of CH3OH at 25°C. [PHT] = 15 mgL–1 or 9 × 10–5 M (monomer unit). Volume percent of CH3OH: 0–80% for a–i as shown. (Yamamoto, T. et al., J. Am. Chem. Soc., 120(9), 2047, 1998. With permission.)

between polythiophene backbones and side chains in organic solvents. This phenomenon is similar to addition of a poor solvent into a stable polymer solution, leading to π-stacking of the polymer molecules and formation of a stable colloidal aggregate [65,68,79,110,115]. For example, addition of methanol to the CHCl3 solution of HT-PHT gives a stable colloidal solution over a wide range of CHCl3/CH3OH ratios. The colloidal solution can yield information about the stacking of the polymer in solution. Figure 5.2.12 illustrates UV-vis spectra of CHCl3 solutions of HT-PHT with various loading of CH3OH. For HT-PATs bearing chiral side groups, reconformation of main chain from nonplanar to planar is accompanied by the appearance of an intense induced circular dichroism in the UV-vis region derived from the main chain or supramolecular chirality [22,114,116–119]. HT-PHT gives the π–π*-absorption band at 450 nm in CHCl3, and addition of CH3OH to the CHCl3 solution leads to a bathochromic shift of the π–π*-absorption band. As shown in Figure 5.2.12(a), the new absorption band shows splitting, and the absorption peaks at 520, 560, and 610 nm agree with those of a film of HT-PHT [56,57,71–73,75,81,120–122], indicating that HT-PHT takes a well-stacked structure in the CHCl3/CH3OH mixture, similar to the film [76,80,111,123–125]. However, PHT containing a majority of head-to-head units (HH) does not seem to form the stacked structure in the CHCl3/CH3OH mixtures at high CH3OH concentrations (Figure 5.2.12b). From these results, Yamamoto and coworkers proposed an aggregate structure of HT-PHT in a colloidal solution (Figure 5.2.13). In this model, HT-PHT molecules aggregate along the main polythiophene chain. Specifically, light scattering analysis of the aggregates supports as the most accepted model for HT-PAT agglomeration in organic solvents the theory that the aggregate tends to grow in the direction of the y and z axes in Figure 5.2.13. The other models predict that, at higher polymer concentration or at poorer solvent quality, HT-PATs aggregate in a face-to-face fashion [126–129] into rod-like micelles [130] in order to decrease the unfavorable interaction between the solvent and the aromatic main chain.



385

x 3.8 A

y

z

16.4 A

FIGURE 5.2.13 A model for the aggregation of HT-PHT in the colloidal solution; solvent = mixture of CHCl3 and CH3OH. The aggregate may consist of several subunits with the sideto-side stacked molecules. (Yamamoto, T. et al., J. Am. Chem. Soc., 120(9), 2047, 1998. With permission.)

On the other hand, Kiriy and coworkers observed a different type of aggregate in HT PAT colloidal solutions, as shown in Figures 5.2.14 and 5.2.15. They demonstrated kinetically driven dynamic growth of one-dimensional aggregates of P3AT adsorbed onto a hydrophobic Si substrate, from a CHCl3/hexane solution using AFM [68]. From AFM and UV-vis analyses for aggregates of HT-PATs in CHCl3/hexane, they suggest that PAT molecules adopt a helical conformation with anticonfiguration of thiophene units in which all sulfur atoms are directed inside the cavity and hydrocarbon groups are oriented outside the helix (Figure 5.2.15). However, on the substrates, dried aggregates inherently have effects of solvent evaporation and substrate-induced crystallization. CHCl3 is a commonly used solvent for film deposition of HT-PHT in field-effect transistors. However, its low boiling point (60.5–61.5°C) and rapid evaporation limit the time for crystallization during a spin-casting process. As a result, charge mobilities achieved for the HT-PHT films are typically on the order of 0.01 cm2/Vs. Recently, it was reported that trichlorobenzene with good solubility and a high boiling point (218–219°C) significantly improved the charge mobility of spin-cast HT-PHT films up to 0.12 cm2/Vs with on/off ratios of 106 [36]. AFM and GIXD analyses for HT-PHT films spin-cast from CHCl3 and TCB support that high crystalline morphology and efficient π–π-stacking parallel to the in-plane current flow in the spin-cast HT-PHT films can be developed with high-boiling-point and highsolubility solvents, resulting in high electric performance of HT-PHT n based OTFTs. As a result, the use of appropriate solvents with high boiling points for spin-casting combines the advantages of increased drying time similar to that of drop-casting while retaining high thin film uniformity. This solvent-dependent structure of


386


(a)

(b)

150 nm

50 nm

300 nm

Height, nm

(c)

(d)

6 4 2 0

–20

(e)

200 HM = 2.2 nm HN = 1.5 nm

–15

PDI = 1.5

–10 5 0 0

2

4

Height, nm

6

Frequency (a.u.)

0

400 30 (f )

600

nm

LM = 428 nm LN = 347 nm

20

PDI = 1.23

10 0 0

300

600

900

Length, nm

FIGURE 5.2.14 Solution AFM topographs of PHT (0.01 g/L) adsorbed onto hydrophobic Si substrates from chloroform/hexane mixture (1:5 v/v): (a, b) immediately after addition of hexane into chloroform solution of PHT; (c) 15 min after addition of hexane; (d) cross-section taken along the particle marked by aHTow. (e) Height distribution of spherical particles adsorbed from PHT solution (0.001 g/L); (f) length distribution of one-dimensional aggregates formed at 0.01 g/L of PHT. (Kiriy, N. et al., Nano. Lett., 3(6), 707, 2003. With permission.)



387

R R R

R

S S

R S

S

S

S

R

R

S R

All anti-conformation

S

RS

(a)

S S

S S

S R

R R

(c) R

S S

R

R

R

(b)

1 nm

1 nm

FIGURE 5.2.15 (a) The helical conformation of PHT molecules. Their one-dimensional aggregation into helical nanotubes: side view (b); top view (c). (Kiriy, N. et al., Nano. Lett., 3(6), 707, 2003. With permission.)

HT-PHT has also been reported in drop-cast films from various solvents with different boiling points and solubilities (Figure 5.2.16) [23]. Figure 5.2.17 illustrates AFM phase images for HT-PHT nanocrystals grown on SiO2 substrates from different solvents: CH2Cl2 and CHCl3. With a low-boiling point (40°C) and poor solubility, CH2Cl2 did not sufficiently develop nanofibrillar structure of HT-PHT during solution deposition, specifically due to fast solvent evaporation rate, showing “nanorod” structure. HT-PHT nanofibrils in CHCl3 (boiling point: 61°C), on the other hand, grew up to several microns in length and were interconnected with each other. The charge mobilities (μ) of the corresponding 30-nm-thick HT-PHT films from CH2Cl2 and CHCl3 were ~1 × 10–4 cm2/Vs and ~0.01 cm2/Vs in bottom-contact OTFTs, respectively.

5.2.2.3 PROCESSING CONDITION As mentioned previously, mesoscale crystalline morphologies and molecular orientation of HT-PATs in solution-deposited films can be controlled through different solution-processing methods and subsequent thermal treatments [15,17,23,35,36,44,59,89]. Figures 5.2.18 and 5.2.19 represent typical examples for tunable crystalline morphologies and molecular orientations of HT-PHT in solution-deposited films. First, in terms of mesoscale morphologies, HT-PHT (M n = 18.8 kDa) films drop-cast from CHCl3 show well developed nanofibrils interconnected with each other, while the spin-cast


388


(a)

(b)

(c)

FIGURE 5.2.16 AFM topographs (left) and phase images (right) of HT-P3HT films dropcast from various solvents in closed petri dishes: (a) toluene; (b) THF; (c) CH2Cl2. (All scale bars represent 1 μm.) (Yang, H. et al., Adv. Func. Mater., 15(4), 671, 2005. With permission.)



389

1 μm (a)

1 μm (b)

1 μm (c)

1 μm (d)

FIGURE 5.2.17 TM-AFM phase images of HT-PHT ( M w = 11.4 kDa, PDI = 2.2) nanofibrils grown on SiO2 substrates from 0.05 and 0.1 wt% solution with different solvents: (a) 0.05 and (b) 0.1 wt% in CH2Cl2; (c) 0.05 and (d) 0.1 wt% in CHCl3. (Yang, H. et al., Adv. Func. Mater., 15(4), 671, 2005. With permission.)

films had smooth featureless surfaces, showing similar results to literature. Visible crystalline development of HT-PHT in the spin-cast films is just observed through thermal treatment above the melting temperature (216°C) because morphological development from less ordered to ordered structure requires reconformation of polythiophene backbones (Figure 5.2.18d). On the other hand, spin-cast HT-PHT films with low MW show crystalline enhancement through thermal annealing below melting temperature [44]. As seen in Figure 5.2.19, two-dimensional GIXD can provide detailed information for crystalline structure and molecular orientation of HT-PHT in these films. First, twodimensional GIXD patterns of the drop-cast film showed intense (100) peaks with higher order peaks and (010) peak along the qz- and qxy-axes, respectively (white arrows in Figure 5.2.19a). The result indicates that the chains have an edge-on structure on the substrate, where π-conjugated polythiophene backbones are preferentially oriented parallel to the substrate (the inset in Figure 5.2.19a). Conversely, the spin-cast film had both edge-on and face-on chain orientations, as indicated from the diffraction peaks of (100) crystal planes in both the meridian (qz) and equatorial (qxy) directions. In addition, the film had much lower crystallinity than the drop-cast film, as calculated by overall integration of the diffracted crystal peaks. Specifically, face-on structure of HT-PHT in spin-cast films tends to increase with increasing substrate temperature. Most interestingly, unlike significant


390


(a)

(b)

(c)

(d)

FIGURE 5.2.18 AFM topographs of HT-P3HT films (Mn = 18.8 kDa) fabricated from CHCl3 under different casting conditions: (a) drop-cast; (b) as prepared; (c) high-temperature annealing (at 200°C, 12 h); (d) meltcrystallized (with a cooling rate of 2°C/min). (All scale bars represent 500 nm.) (From Yang, H., unpublished data, 2006. With permission.)

enhancement in annealed film morphologies, the molecular orientation of HT-PHT in the melt-crystallized films is not equivalent to that in the drop-cast film. It is related to the fact that semicrystalline polymers have residual local order even in the melt (termed the “melting memory effect”), suggesting that initial self-assembled structures of HT-P3HT on dielectric substrates play an important role in the electrical properties of the corresponding films. The AFM and two-dimensional GIXD results for high MW spin-cast HT PHT films strongly support the conclusion that low charge mobility (≤10–3 cm2V–1s–1) in most spin-cast films from HT-PHT soluble in CHCl3 is mainly related to the poor crystallinity and unfavorable orientation of π-conjugated backbone planes normal to the substrate (i.e., the in-plane current flow) [89]. During direct solution deposition of HT-PHTs on dielectric substrates, the solvent evaporation rate can be changed in a wide range from a few seconds to a few hours by controlling a combination of the following parameters: solvent boiling point processing mode: spin-casting, drop-casting, and controlling the substrate temperature solvent vapor pressure (νp) in film deposition facility



391

(a)

(b)

(010) Substrate

(200)

(100) 1.0

1.5

2.0

qxy (Å−1) (c)

(010)

0.5

(100)

0

(100) (300)

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(200)

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FIGURE 5.2.19 Two-dimensional GIXD patterns of drop-cast (a) and spin-cast (b–d) HTP3HT thin films obtained from CHCl3 solution: (b) without any thermal treatment; (c) annealed at 200°C for 12 h; (d) melt-crystallized from 240°C with a cooling rate of 2°C/min. (The insets in the upper right corner of (a) and (b) represent orientation of π–π-conjugated planes in HT-P3HT chains with respect to the SiO2 substrates.) (From Yang, H., unpublished data, 2006. With permission.)

The first two parameters were discussed earlier in this chapter. Generally, the spin-casting method cannot produce nanofibrillar morphologies of HT-PHTs compared to the drop-casting method because of fast evaporation rates. However, during spin-casting, the use of appropriate solvents with high boiling points or the control of νp can provide increased drying time similar to that of drop-casting while retaining high thin film uniformity. Recently, Cho and coworkers demonstrated tunable film structures of HT-PHTs through controlling νp under a spin-casting condition (1,000 rpm) [67]. Figure 5.2.20 represents AFM images for 98% HT-PHT ( M w = 54 kDa, PDI = 1.8) films spin-cast from CHCl3 under various solvent vapor pressures. In the spin-cast films, the HT-PHT nanocrystals grew from short nanorod to long nanofibrils by increasing the νp. Specifically, nucleation and growth behavior of HTPHT crystals are clearly displayed as shown in Figure 5.2.21. On SiO2 substrates covered with the HT-PHT solution during drop-casting, low solvent vapor pressure induced by an open petri dish setup cannot provide sufficient time for HT-PHT crystal growth to occur. Instead, aggregates of short nanorods were observed. On the other hand, an increase in the solvent pressure by a closed environment yields several micron-length nanofibrils. They fabricated micro-size single-crystal wires (height: 700 nm–1.3 μm; width: 1–3 μm; length: 30–500 μm) of PHT (weight-average


392


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(b)

(c)

(d)

(e)

(f )

FIGURE 5.2.20 TM-AFM images (2 × 2 μm2) of PHT molecular nanowires deposited on a substrates (SiOx) at the following solvent vapor pressures: atmospheric condition (a); 6.2 kPa (b); 36.5 kPa (c); 48.9 kPa (d); 53.8 kPa (e); 56.6 kPa (f). (Kim, D.H. et al., Macromol. Rapid. Commun., 36(3), 691, 2005. With permission.)

molecular weight = 54 kg mol–1; regioregularity = 98.5%) through recrystallization of a dilute PHT (0.1 mg/L) solution in CHCl3 (Figure 5.2.22) [131]. In solution-deposited films on dielectrics, the crystalline orientation and morphology depend considerably on polymer–substrate interaction, which can be controlled by modification of the dielectric surface with a self-assembled monolayer (SAM). Cho et al. have controlled the intermolecular interaction at the interface between HTPHT and the dielectric substrate by using SAMs functionalized with various groups (–NH2, –OH, and –CH3) [24]. They have found that, depending on the properties of substrate surface, HT-PHT nanocrystals in the spin-cast films adopt two different orientations of π-conjugated planes — parallel and perpendicular to the substrates. These have field-effect mobilities that differ by more than a factor of four and are as high as 0.28 cm2V–1s–1 for π-conjugated planes parallel to the substrates.

5.2.3 CONCLUDING REMARKS This chapter has described different approaches for the deposition of conjugated polymer films from solution. Many of these films have shown excellent device performance, with mobilities as high as 0.6 cm2V–1s–1. Recently, advances in device performance and device stability have been achieved through both molecular design and the optimization of processing conditions. The investigation of polymer ordering at the semiconductor–dielectric interface has led to improved understanding of the charge transport mechanisms in polymer thin film devices. With continuing improvements in device parameters and stability, conjugated polymer thin film transistors remain the leading candidate for all printed organic transistors.



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(a)

(b)

(c)

FIGURE 5.2.21 AFM topographs (left) and phase images (right) of HT-P3HT films dropcast from CH2Cl2 under different solvent evaporation conditions: (a) open petri dish; (b) closed petri dish; (c) closed jar. (All scale bars represent 1 μm.) (From Yang, H., unpublished data, 2006. With permission.)


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S 8.36Å

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ax

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ire

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S

7.80Å

S

b

S

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S

16.60Å

(b) S

S S

S S

3.9Å 3.9Å

S

(002)

FIGURE 5.2.22 Morphological features and structure characterization of single-crystal PHT microwires. (a) POM image of one-dimensional single-crystal PHT microwire bundles. (b) Field-emission scanning electron microscopy images of PHT wires formed on a silicon substrate modified with an ODTS dielectric layer. The inset shows a side-view image showing the rectangular cross-section of PHT microwires with well-defined facets. (c) Transmission electron microscopy image of PHT wires on a silicon nitride window modified with ODTS, showing preferential growth along the [010] direction. (d) Selected-area electron diffraction pattern of PHT microwire showing diffraction equivalent to a repeating period of 7.80 Å along the π–π-stacking direction, and a repeating period of 8.36 Å along the chain direction. (e) Molecular and crystallographic structures (orthorhombic unit cell) of PHT chains show the enhanced π–π-overlap along the b direction. (f) Schematic representation of the hierarchical self-assembly of PHT chains into single-crystal microwires along π–π-stacking direction. The height, width, and length of the wires are approximately 700 nm–1.3 μm, 1–3 μm, and 30–500 μm, respectively. (Kim, D.H. et al., Adv. Mater., 18(6), 719, 2006. With permission.)



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5.3

Solution Deposition of Oligomers

Howard E. Katz and Chad Landis CONTENTS 5.3.1 Introduction................................................................................................403 5.3.2 Conjugated Oligomers ...............................................................................404 5.3.3 Fused Ring Compounds ............................................................................409 5.3.4 Conclusion and Future Prospects ..............................................................414 References..............................................................................................................414

5.3.1 INTRODUCTION Electronic technology based on printable and largely organic conductors, dielectrics, and semiconductors is a major academic research theme and the basis of significant investment by established and emerging companies [1–4]. While the usual set of passive components such as resistors and capacitors would be needed to produce functional circuits, the key device in these technologies is the organic field-effect transistor (OFET). Platforms have been identified where organic semiconductors offer lower cost fabrication compared to silicon or provide functionality that is distinct from what can be obtained from noncovalent and rigid materials. Prototypes of OFET-based display drivers [5–10] have been demonstrated. The most compatible display media for OFET driver circuits are “electronic inks” — materials in which contrasting pigment particles or domains can be electrophoretically shifted into and out of a field of view. While requiring relatively high switching voltages, they draw little current during switching and virtually no current between switching cycles, matching the typical OFET output. These media provide convincing visual evidence of the effectiveness of the organic-based switches. However, liquid crystal and organic emitter-based materials have also been employed as the visual media. The high-throughput processes and mechanical flexibility associated with these prototypes may open the possibility for electronic newspapers and books onto which information can be downloaded, read, and reset with further information as the reader requests. A second proposed use of organic semiconductors is in radio frequency ID (RFID) tags [11,12]. Several hurdles must be overcome in order to make organic RFID practical. The first is the transistor speed, which will only be adequate for

403


404


frequencies > 1-kHz and 1-V operation with mobilities > 0.1 cm2/Vs and channel lengths < 100 μm. At present, these are the limits for easily printed, all-organic semiconductors. Higher frequencies, which are necessary for all but the most proximate detection schemes, will only arise from much shorter and higher mobility devices. In addition, rectification of the AC input to provide a DC logic signal is challenging in all-organic systems because the fastest diodes tend to be inorganic and not particularly amenable to printing. Schemes to accomplish rectification without diodes have been proposed. In the future, it might be possible to utilize the same kinds of organic semiconductors that have been developed for OFETs to make printable diodes. Related work aimed at printed solar cells is already in progress. Nonprinted organic solar cells with breakthrough efficiencies have recently been reported [13–20]. In addition, organic light-emitting diodes are already in commercial products [21], though the role of particularly high-mobility, printed organic semiconductors has not been extensively considered. While numerous manuscripts have appeared describing new or modified conjugated cores for hole-carrying organic semiconductor molecules and several recent publications have focused on new approaches to electron-carrying molecules in films, very few consider solution deposition of molecular solids in detail. Some recent reviews have discussed the subject relatively briefly [22–24]. Molecular solids offer potential advantages over polymers in that they are generally more ordered, leading to higher mobility, and are more readily purified, leading to better stability and reproducibility. However, the printing of a molecular solid from solution is far more challenging. Molecular solutions maintain low viscosity even as much of the solvent is removed, so film localization is problematic. As crystallites form within the deposited solution, they typically remain separated from each other rather than forming a continuous film. Conversely, at the point that a film forms on a substrate, the molecules are held so tightly in a crystal lattice that they cannot anneal into their best ordered morphology, so macroscale mobility suffers. Thus, the solution deposition of molecular solids requires attention to the molecular design, for the most favorable solvent and substrate interactions, and also the deposition conditions, so that the most favorable phases and highest degrees of continuity can be achieved. This review summarizes techniques that have evolved for the deposition of highmobility organic semiconductor films from solutions of properly designed conjugated molecules. Such molecules are distinguished from polymers by having molecular weights below 1,000. In addition, they tend to have high symmetries and a balance between the volumes of the conjugated cores and solubilizing/self-organizing side chains. The review will be presented in two main sections dealing with heterocyclic oligomers and fused ring compounds, respectively.

5.3.2 CONJUGATED OLIGOMERS The first soluble oligomer with significant field-effect mobility was 5,5′′′-dihexylquaterthiophene, described by the Garnier group in 1998 [25]. The chemical structure is shown in Figure 5.3.1. The compound exhibits a mesophase transistion below 100°C. The hexyl chains improve the ordering of the compound relative to unsubstituted quaterthiophene, which in turn is believed to lead to improved mobility as



C6H13

405

S

S S

S

C6H13

n Dihexylquartethiophene (n = 1) Dihexylquinquethiophene (n = 2)

R

S

S S

S S

S

R

Sexithiophenes: R = hexyl, butoxypropyl, diethylphosphonylbutyl

FIGURE 5.3.1 Chemical structure of alkyl-substituted oligothiophenes.

well. For sublimed films, the improvement is at least an order of magnitude. Importantly, a dihexylquarterthiophene film spin-coated from chloroform onto a hot, primed substrate had mobility only 60% lower. Transistors with vertical architecture were made from this solution [26]. It was shown at about the same time that the special properties of this oligomer could also be related to the large-area, single-crystal grains (Figure 5.3.2.) that could be grown at elevated temperature [27]. The ability to form films with such morphology is shared by few organic semiconductors; pentacene is a notably important example. The ability to tune the size of the grains by varying the substrate temperature during deposition was useful in demonstrating that grain boundaries are a site of chemical sensitivity in OFETs [28]. Attempts to employ solution processing on longer thiophene oligomers were limited by the drastically lower solubility of these compounds. While mobilities for compounds such as dihexyl and bis(butoxypropyl) sexithiophene were on the order of 0.01 cm2/Vs from films cast from solution, the solutions were 1,000. Cross-linked methacrylate and siloxane dielectrics were used, and active areas up to 2 cm2 were obtained. The nonchlorinated solvents xylene and toluene were usable. Extensive data were obtained from the related compound 5,5′-bis(4-hexylphenyl)-2,2′-bithiophene, abbreviated 6PTTP6 [36]. Xylene was an ideal solvent for casting films of this compound, keeping the substrate 5–30° below the boiling point.



407

S

C6H13

S

S

C6H13

S

C6H13

S

S

C6H13

Phenylene-thiophene oligomers: bis(hexylbithienyl) benzene and 6PTTP6

FIGURE 5.3.4 Chemical structures of phenylene-thiophene co-oligomers.

Highly ordered films were obtained on primed oxide and glass resin surfaces. The on/off ratios exceeded 10,000, with the off currents reproducibly low even without having purified the semiconductor by vacuum sublimation. Optimization of multiple processing parameters, including surface functionalization, substrate temperature, and overlying atmosphere (saturation with xylene vapor), led to a highly ordered film by virtue of the traversal of lyotropic mesophases (Figure 5.3.5) [37]. The ordering was monitored by polarized optical microscopy and x-ray scattering. The highest solution-deposited oligomer mobility yet obtained, >0.1 cm2/Vs, was derived from an optimal sample (Figure 5.3.6). This semiconductor was used to introduce the concept of solution-deposited nonvolatile organic transistors [36]. (a)

(b)

FIGURE 5.3.5 Images of lyotropic mesophases from a film of solution-deposited 6PTTP6.


408


0.012

–140 –120 –100 –80 –60 –40 –20 0 20

–100 V

0.010

–80 V

ISD1/2 (A1/2)

ISD (μA)

Y=A+B∗X Parameter Value Error ---------------------------------------------A –3.34643E-4 3.30694E-4 B –1.19112E-4 4.98541E-6 ---------------------------------------------R SD N P ---------------------------------------------–0.99738 3.15305E-4 5 1.60685E-4 ----------------------------------------------

–60 V –40 V –20 V 0

0.008 0.006 0.004 0.002 0

–20 –40 –60 –80 –100 Source-Drain voltage (V)

–20 –40

–60 –80 –100

Gate voltage (V)

FIGURE 5.3.6 Current-voltage characteristics of a solution-deposited 6PTTP6.

C8F5OC

S

S

S

S

COC8F5

Bis (pentafluorophenacyl) quarterthiophene

FIGURE 5.3.7 Chemical structure of the first n-channel, solution-deposited, thiophene oligomer.

While all of these soluble oligomers have displayed p-channel character (hole transporting ability), there is no reason why oligomers of high enough purity and low enough energies of the lowest unoccupied molecular orbitals would not act as n-channel semiconductors. The persistent addition of fluoro and oxo substituents to end-substituted thiophene oligomers has recently led to the first n-channel solution-deposited thiophene oligomers. A mobility of 0.2 cm2/Vs was obtained by drop-casting bis(pentafluorophenacyl)quaterthiophene at 120°C (Figure 5.3.7) [38]. The compound was also successfully blended with a polymeric analog for improved printability. Evidence from dihexylquaterthiophene blended with regioregular poly(3-hexylthiophene) suggests that the blending approach may be more generally applicable [39]. The oligomer considered as an additive improved the mobility of the polymer by a factor of 10. Because of the limited solubility conferred by linear side chains and the desirability of more concentrated solutions for various printing methods, branched side chains are attractive. However, branched chains tend to detract from the ordering of molecular solid films and inhibit high mobility. A collaborative group at UC Berkeley is exploring oligothiophenes with highly branched but thermally cleavable ester side chains [40–43]. The carbon attached to the terminal alpha positions of the oligomer carries an aliphatic and a carbonyloxy substituent. The carbonyloxy is further branched. Thus, highly soluble derivatives are obtained, and the stability of the cation centered on the alpha carbon leads to thermal lability. On heating, the



O

S

S S

S

n

O

O

C6H13

C6H13

n=1–4

O

C4H9

409

C4H9

S

S S

S n

n=1–4 Oligothiophene ester thermolysis

FIGURE 5.3.8 Schematic showing the thermolysis of highly soluble oligothiophene derivatives.

carbonyloxy group is volatilized and a short hydrocarbon chain is left at each alpha position (Figure 5.3.8). Films with a wide variety of morphologies have been obtained, depending on the nature of the substituents, substrate surface treatments, and temperature excursions. Mobilities approaching 0.1 cm2/Vs and very high on/off ratios — into the millions — have been obtained. Materials were deposited by spincoating, dip-casting, and inkjet printing. Active films just a few monolayers thick were demonstrated.

5.3.3 FUSED RING COMPOUNDS The most studied semiconducting fused ring molecule is pentacene, due to its high mobility (1 cm2/Vs) in OFETs and good film-forming capabilities [44,45]. Unfortunately, pentacene is essentially insoluble and is therefore not suitable for solution processing, which is necessary in order to realize cheap, large-area electronics. One example of direct solution processing of pure pentacene requires deposition of a heated solution of pentacene in 1,2-dichlorobenzene and trichlorobenzene under a nitrogen atmosphere [46]. The devices prepared by this method showed a mobility of 0.45 cm2/Vs and on/off ratios of 105. However, the most popular response to the insolubility of pentacene has been the preparation of soluble precursors that can be deposited on a substrate and then converted into pentacene. The first example of a pentacene precursor was a tetrachlorocyclohexadiene adduct prepared by the Müllen group (Figure 5.3.9a) [47]. This derivative is soluble in dichloromethane and forms good films by spin-coating. After deposition, the pentacene film is formed by a thermally activated retro Diels–Alder reaction expelling tetrachlorobenzene as the by-product. The hole mobilities of OFETs prepared from these pentacene precursors depended greatly on the annealing temperature


410


Cl Cl

O

Cl O

Cl

(a)

N S

O

R

(b)

O O

O

N N

(c)

FIGURE 5.3.9 Structures of pentacene precursors.

used, ranging from 5 × 10–5 to 3 × 10–3 cm2/Vs [48]. More recent optimized conditions require annealing at 200°C for 5 sec; thin film transistors prepared in this manner have highest hole mobilities of 0.2 cm2/Vs and on/off current ratios of 106 [49]. This derivative has also recently been used to prepare drivers for a 3.5-cm2 flexible display [5]. More numerous examples of pentacene precursors employ sulfinylamide functionality (Figure 5.3.9b). The first of these derivatives was prepared by the Afzali group and utilized a chloroform-soluble N-sulfinylacetamide derivative (Figure 5.3.9b, R = CH3) where pentacene films were formed thermally by a retro Diels–Alder reaction [50]. There was still a strong temperature dependence between the annealing temperature of the film and the resulting mobility of the devices, with OFETs prepared by annealing at 200°C resulting in hole mobilities of 0.42 cm2/Vs, while devices prepared at 130°C demonstrated mobilties of 0.13 cm2/Vs. Other sulfinylamides have been prepared that allow for differing solubilities [51] or even derivatives that undergo photoinduced retro Diels–Alder reactions that allow for high-resolution features to be printed by curing only the desired regions [52,53]. In one example, a photosensitive precursor was used to pattern an OFET with features of 40 μm and with mobilities of 0.25 cm2/Vs [53]. However, a major problem with the pentacene precursor method is that a residual amount of the Diels–Alder adduct remains in the film and limits the carrier mobility of the final pentacene film [54]. To limit this problem, pentacene precursors with smaller or more easily eliminated Diels–Alder adducts have recently appeared, but the field-effect transistor characteristics of these new derivatives are yet to be published (Figure 5.3.9c) [55,56]. The concept of soluble precursors has also recently been extended to other fused aromatic systems. For example, tetrabenzoporphorins have been prepared from thermal annealing of a soluble precursor [57]. The precursor in this case is a tetrabicyclo derivative (Figure 5.3.10a). Upon thermal annealing at 210°C for 5 min, this derivative eliminates four ethylene molecules, resulting in the electroactive tetrabenzoporphyrin (Figure 5.3.10b). Bottom-contact OFETs prepared from this material had hole mobilities of 0.017 cm2/Vs and on/off current ratios of 105, which is similar to OFETs prepared from vacuum evaporated tetrabenzoporphyrin [58]. Another series of devices from this porphyrin derivative utilized a coating of an n-channel material to limit the effects of oxygen and moisture on the performance of the device [59]. These devices were top-contact configuration, with the n-channel material coated after the electrodes are deposited. The OFETs prepared in this manner displayed



411

N NH

N NH

NH

NH

N

N

(a)

(b)

FIGURE 5.3.10 Structure of tetrabenzoporphyrin (a) and its precursor (b).

hole mobilities of 5 × 10–3 cm2/Vs, similar to top-contact OFETs prepared without the n-channel overlayer [60]. There have been few examples of direct functionalization of fused aromatic systems. Work done by the Katz group used anthradithiophenes instead of pentacene because of the greater oxidative stability resulting from the two fused thiophene rings. In particular, a solution-processed OFET was prepared from the dihexylanthradithiophene derivative (Figure 5.3.11a) [61]. This derivative is soluble in hot chlorobenzene and was formed into films on substrates under vacuum held at 100°C. The OFETs prepared from this derivative had hole mobilities ranging from 0.01 to 0.02 cm2/Vs, but were difficult to reproduce. R

S C6H13

C6H13

S

R (a)

R = (μPr)3Si, Me3Si, Et3Si

(b)

SiEt3 O

O

C7F15CH2 N

N

O

O

S S

SiEt3 (c)

FIGURE 5.3.11 Structures of soluble fused ring derivatives.

(d)

CH2C7F15


412


The Anthony group has prepared an extensive library of soluble acene derivatives. The derivatives employ trialkylsilyl solublizing groups held from a pentacene core by ethynyl spacers at the 6 and 13 positions (Figure 5.3.11b) [62,63]. The result of this functionality is a change in the molecular order of these derivatives from the herringbone orientation of unsubstituted pentacene to a one- or two-dimensional πstacking arrangement. These derivatives also demonstrate greatly increased solubility and oxidative stability compared to unsubstituted pentacene. They are soluble in a wide variety of organic solvents and are air stable as a solution for days and as crystals for weeks. One of the best of these derivatives is 6,13-(bistriethylsilylethynyl)anthradithiophene (Figure 5.3.11c) [64,65]. OFETs prepared by toluene solution casting of a mixture of the isomers of this derivative resulted in hole mobilities of up to 1 cm2/Vs and on/off current ratios of 107. Soluble derivatives of fused aromatic systems are not limited to p-channel materials; however, they are far fewer in number. A soluble derivative of NTCDA has recently been developed that can be cast from α,α,α-trifluorotoluene (Figure 5.3.11d) [66]. Devices prepared from this derivative showed electron mobilities of ~0.01 cm2/Vs [67]. This is currently the only nonfullerene solution-cast nchannel material. Directly functionalized discotic molecules have also recently been prepared to serve as liquid crystalline materials, which could be used for electronic devices. One of the impressive points of these materials is their ability to self-assemble into columnar arrays. One particular example involves a derivative of hexabenzocoronene, which is induced into a high mesoscopic order by the introduction of a magnetic field (Figure 5.2.12a) [68]. OFETs prepared from this liquid-crystalline material showed mobility of 10–4 cm2/Vs when the magnetic field is held perpendicular to the charge transport direction and 3 × 10–6 cm2/Vs when it is held parallel to the charge transport direction. A n-channel example of these materials is a tris(pentafluorobenzylester) derivative of hexaazatrinaphthylene (HAT-NA, Figure 5.3.12b) [69]. Films of the isomer shown in the picture formed textured films by rapid cooling from a melt. While OFETs were not prepared from this derivative, the measured effective electron mobility of the film was found to be 0.07 cm2/Vs. A soluble fullerene derivative is a desirable semiconducting material due to the effective electron transporting capability of C60. The most used soluble fullerene derivative is [6,6]-phenyl-C61-butyric acid methyl ester (PCBM, Figure 5.3.13b); the majority of research into this derivative is in its use as the n-channel material in bulk heterojunction solar cells [70,71]. This derivative is soluble in many organic solvents and tends to form amorphous films that result in OFETs with electron mobilities up to 0.1 cm2/Vs [72,73]. Mobilities were improved to 0.2 cm2/Vs by using a PVA dielectric [74,75]. This derivative has also been used to prepare ambipolar OFETs [76]. More recently, another derivative has been prepared in an attempt to induce greater molecular order in spin-cast films of soluble C60 and thereby improve the electron mobility of the films. This derivative, C60MC12 (Figure 5.3.13a), employs a dodecyl alkyl chain for solubility and to induce long-range order by interacalation of the alkyl chains [77]. Spin-cast films of this derivative show greater order and OFETs prepared from it display electron mobilities of 0.067 cm2/Vs and on/off current ratios of 1.6 × 105.



413

C12H25

R C12H25

C12H25 N

N

N R

N N

N R

C12H25

C12H25

R = CO2CH2C6F5

C12H25 (a)

(b)

FIGURE 5.3.12 Chemical structure of two soluble discotic organic semiconductors. C12H22 O

Me

O

N

C60MC12

[6, 6]- PCBM

FIGURE 5.3.13 Structures of two soluble fullerene derivatives.

During the past decade, the limits of carrier mobility have continuously been sought. Currently, the upper limit for organic single crystals is rubrene with an intrinsic mobility of ~15 cm2/Vs (Figure 5.3.14, left) [78]. Evaporated films of rubrene have not been effective and methods of preparing thin film devices of rubrene have not been fruitful. One unique method has been developed to prepare rubrenebased OFETs. This relies on an analog of rubrene (diphenylanthracene; Figure 5.3.14, right) to act as a glass-inducing impurity to inhibit crystal growth of the rubrene until after casting of the film [79]. A 55 wt% solution of rubrene with diphenylanthracene and an additional 5 wt% of ultrahigh-molecular-weight polystyrene (UHMW-PS) was cast from toluene (the UHMW-PS acting as a mechanical support for the film). The films were held at the crystallization temperature for rubrene in this particular mixture (235°C) for 2–5 min and then cooled. Devices prepared from this method showed hole mobilities up to 0.7 cm2/Vs and on/off current rations of 107.


414


FIGURE 5.3.14 Chemical structure of rubrene and diphenylanthracene.

5.3.4 CONCLUSION AND FUTURE PROSPECTS This chapter has described numerous approaches to the deposition of OFET semiconductor films from solutions of organic compounds. Many of these films were demonstrated in devices with mobilities from 0.01–1 cm2/Vs and could in principle be patterned through printing processes. Factors limiting the usability of these compounds in printed logic circuits are similar to the limitations on organic semiconductor use in general: semiconductor stability, uniformity of coverage and electrical properties, and interfacial effects. In addition, physical properties of the solutions, such as surface energy and viscosity, need to be optimized for particular printing approaches. Little systematic work has been done along these lines, leaving a significant opportunity to develop these solutiondeposition methods taking the liquid-phase properties into account. Judging from the promising electronic properties reported for individual devices and the appearance of initial circuit prototypes, such development could prove particularly fruitful.

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5.4

Inkjet Printed Organic Thin Film Transistors

Ana Claudia Arias CONTENTS 5.4.1 Introduction................................................................................................419 5.4.2 Subtractive Methods — Printing Electrodes .............................................420 5.4.3 Additive Methods — Printing Semiconductors and Encapsulation Layers.........................................................................................................422 5.4.4 Display Backplane Fabrication..................................................................427 References..............................................................................................................431

5.4.1 INTRODUCTION Materials development together with the better understanding of organic-based transistors has enabled solution-processed materials that are stable in air, show good performance, are compatible with flexible substrates, and can be used in applications where thin film transistors (TFTs) are the building blocks of electronic circuits. Suitable applications for organic TFTs are large-area devices, such as active matrix displays, where high switching speeds may not be essential. The significance of polymers in particular is that they can be deposited from solution, therefore allowing device patterning by direct marking techniques. Inkjet printing is an attractive process because it is low cost, applicable to large-area processing, and compatible with flexible substrates, and can be adapted to high-throughput manufacturing processes, such as roll-to-roll printing methods. The technique provides a drop-ondemand digital deposition process without the need for physical masks, and materials are only applied where they are needed, decreasing materials cost and environmental impact. A variety of printing approaches has been used to fabricate TFTs. Each has different attributes with advantages and disadvantages regarding registration, process temperature, and device performance. For example, contact printing has been demonstrated for patterning electronics, but tends to have poor registration accuracy for multilayer structures. Patterning techniques using a fixed printing master are ill suited to flexible substrates unless their dimensional stability can be precisely controlled. In this section, different approaches used to fabricate organic-based TFTs using inkjet printing as a patterning method and as a deposition method will be reviewed.

419


420


The first use of inkjet printing on the fabrication of organic-based TFTs was to pattern the electrodes [1,2]. Printed source-drain electrodes were achieved by combining inkjet printing with photolithography, achieving small features in one layer of the device while the gate electrodes were directly printed on top of the gate dielectric layer [1]. A different approach used to pattern electrodes by inkjet printing is the combination of inkjet printing with conventional wet etching processes [3,4]. In order to achieve a patterned device feature, subtractive steps are used in both methods. These techniques are reviewed in Section 5.4.2. Direct writing was used later to deposit polymeric semiconductors. In direct writing, there are no subtractive steps and materials are not wasted. However, device dimensions are limited by ejector sizes and interactions between solution and substrate. The fundamental parameters controlling inkjet printed liquids are the viscosity and surface energy [5]. The pattern formed when an ejected drop of liquid hits the surface depends in large part on the ink–surface interaction. The wetting contact angle determines the spread of a liquid drop on the surface and depends on the relative surface energy of the solid–liquid, solid–vapor, and liquid–vapor interfaces. High-energy surfaces result in a small wetting angle and an extended drop, while a low surface energy results in a smaller footprint. The surface energy and contact angle also relate to the adhesion of the liquid to the surface. Strong adhesion is associated with good wetting and low adhesion with high contact angles (poor adhesion). Unfortunately, most situations need a high contact angle to limit the spread of the printed liquid and good adhesion to the surface to allow further processing [6]. To inkjet print fine features on a flat surface is not an easy task because of the difficulty in controlling the spread of a liquid on a free surface. The relation between printed line-width and contact angle is displayed in Figure 5.4.1. The solid lines are obtained from a simple model of a small volume of liquid with a cylindrical surface. The measurements of printed nanoparticle metals on different surfaces, shown in Figure 5.4.1, follow the expected trend; the printed line width decreases as the contact angle increases [6]. The printed feature size will not only determine adhesion of material to the substrate, but will also play an important role on the design of devices. Important parameters related to design of printed TFT backplanes are discussed in Section 5.4.4.

5.4.2 SUBTRACTIVE METHODS — PRINTING ELECTRODES Inkjet printing can be used to deposit resist materials that can then replace photoresist in the fabrication of conventional TFT amorphous silicon (a-Si) processes [3,7]. In this technique, conventionally deposited materials (metal, semiconductors, and dielectrics) can be patterned by printing following the steps shown in Figure 5.4.2. In this process, a thin metal film — for example, gold — is deposited everywhere. But instead of coating with photoresist and exposing, the mask is inkjet printed directly onto the layer, as shown in Figure 5.4.2(b). The amount of wax reflow that occurs before solidification can be adjusted by controlling the substrate temperature. Drop sizes from 20–40 μm are possible with industrially produced printheads.



(a)

421 (b) Φ = 30°

3

Line width (relative)

2.5

350 μm

2

Φ = 60°

1.5

180 μm

1

Φ = 80°

0.5 0

100 μm 0

20

40

60

80

100

Contact angle (deg)

FIGURE 5.4.1 Calculated printed line width as a function of wetting contact angle, assuming a small cylindrical liquid pattern. Square data points and the optical micrographs show measurements of printed Ag nanoparticles ink on surfaces that show different water contact angle. (From Street, R.A. et al., Mater. Today 9, 32–37, 2006. With permission.)

Digital lithography

Printed wax features

300 μm (b) Etched metal features

300 μm (a)

(c)

FIGURE 5.4.2 (a) Steps of patterning a metal feature using inkjet printing and digital lithography. (b) Optical micrographs of printed wax features on a metal film. Printed wax features are used as an etch mask for conventionally deposited materials. (c) Optical micrographs of metal features patterned by inkjet printing.


422


Because these heads exhibit superior directionality, features produced by gaps may be fabricated down to 5 μm. In Figure 5.4.2(b), an example of a printed wax etch mask is displayed. The line width is 40 μm and the smallest spacing between two printed lines is 5 μm. The printed wax features work as an etch mask protecting the material that needs to remain on the substrate. The exposed material is etched, and the mask is stripped. The etched features are displayed in Figure 5.4.2(c). This process is called digital lithography because it uses subtractive methods for patterning, as in conventional semiconductor processing. But the mask used here is a digital mask (a file that is printed) instead of hard masks used for photolithography. Paul et al. combined wax printing with direct writing of semiconductors to fabricate polymer-based TFTs [4]. Device performance and reproducibility are discussed in Section 5.4.3 [4,8]. Another hybrid technique used to print high-resolution features is to inkjet print the active material into a predefined well. This well is usually obtained by photolithography to create a prepatterned feature that exhibits higher resolution than what is obtained by direct writing. In this technique, the liquid spread and the drying pattern are controlled by the well. The liquid flows over the surface until it reaches the well wall, which prevents further spread. Since the resulting pattern does not depend on exactly where the liquid is inkjetted, the precision requirement for the printing system is reduced. Both polymer light-emitting diodes (PLEDs) and color filters are made using this technique and are expected to be the first applications of inkjet printing in display applications [9,10]. Sirringhaus and Kawase used this approach to confine the spreading of waterbased conducting polymer ink, PEDOT, to define the critical dimensions of a TFT (Figure 5.4.3a) [1,2]. A line of PEDOT droplets is deposited into hydrophilic wells at a distance d from a hydrophobic polyimide line that is sufficiently small for the spreading droplets to reach the repelling line. After deposition of the source-drain material, TFT devices are fabricated in a top-gate configuration (Figure 5.4.3c) by spin-coating a continuous film (150–300 Å) of the active semiconducting polymer, poly(9,9-dioctylfluoreneco-bithiophene) (F8T2), from a xylene solution. A 400- to 500-nm thick film of the gate dielectric polymer, polyvinylphenol (PVP), is spincoated from isopropanol solution. Finally, a PEDOT/PSS gate electrode line overlapping the channel is inkjet printed in air. In these devices, the semiconductor, F8T2, is not patterned. After deposition by spin-coating, the device was annealed to 265°C to bring F8T2 to its liquid crystal phase and achieve higher TFT mobility [11]. The TFT mobility extracted from the transfer characteristics was 0.01 cm2/Vs. The threshold voltage was equal to –10 V and the Ion/Ioff ratio was equal to 105 for gates voltages from –40 to 10 V and W/L of 600.

5.4.3 ADDITIVE METHODS — PRINTING SEMICONDUCTORS AND ENCAPSULATION LAYERS One of the great advantages of polymeric semiconductors is that they can be deposited from solution, therefore allowing device patterning by direct marking



n

(a)

m

423

(c)

–

SO3 HSO3 O O O O S S nS S S O O O O H

S

S

n

O O

PVP G

PI

d

S

D

(d)

(b)

S 200 nm

50 μm

G D

μm PEDOT 10

PEDOT

PI 20

30

40

20 10 μm

Channel

FIGURE 5.4.3 (a) Schematic diagram of high-resolution inkjet printing onto a prepatterned substrate. (b) AFM showing accurate alignment of inkjet-printed PEDOT/PSS source and drain electrodes separated by a repelling polyimide (PI) line with L = 5 μm. (c) Schematic diagram of the top-gate inkjet printed TFT configuration with an F8T2 semiconducting layer (S, source; D, drain; and G, gate). (d) Optical micrograph of an inkjet printed TFT (L = 5 μm). The arrow indicates pronounced roughness of the unconfined PEDOT boundary. (From Sirringhaus, H. et al., Science 290, 2123–2126, 2000. With permission.)

techniques. The ability to pattern discrete areas of semiconductor improves performance of devices by reducing the off current and in an array by reducing deviceto-device leakage [8]. Paul et al. demonstrated printed organic TFTs that were fabricated on a layer of thermal oxide with a common silicon gate electrode [4]. The metal source and drain contacts were patterned by digital lithography using printed wax as an etch resist, as described in Section 5.4.2. The inkjet printing technology used for the polymeric organic semiconductors, acoustic inkjet printing, is described elsewhere [12]. The OTFTs demonstrated by Paul and coworkers used coplanar device geometry. First, gold contact metal (100 nm on a 2-nm adhesion layer of chromium) was deposited by evaporation onto a silicon wafer having a 300-nm layer of thermal oxide. The gold film was then patterned by digital lithography to define the source and drain contacts. Channel lengths varied between 40 and 400 μm, resulting in a range of W/L between 0.3 and 27. The thermal oxide gate dielectric was chemically modified to be hydrophobic using either octadecyltrichlorosilane (OTS) or octyltrichlrosilane (OTS-8) by solution-based deposition [4]. The semiconductor was then deposited by acoustic inkjet printing onto the patterned and treated substrate. Two polymeric organic semiconductors were studied: (1) 0.45% solution of


424


poly(9,9′-dioctyl fluorene-co-bithiophene), F8T2 [11,13] in xylene; and (2) a solution of a regioregular poly(thiophene) (PQT-12) synthesized at Xerox Research Centre of Canada [14]. The morphology of the printed semiconductors is shown in Figure 5.4.4. The ejected drops coalesce within a printed line; however, if the solvent evaporates before the next line is printed, the dried lines result in the striated topography seen in Figure 5.4.4. Figure 5.4.4(a) is an optical micrograph of printed PQT-12 and Figure 5.4.4(b) shows an image of printed F8T2 acquired by tapping mode atomic force microscopy. The topography of the surface shows the increased thickness of the polymer as each line is applied with a drop size of 35 μm and an overlap ratio of 50%. (a)

Print direction

400 μm

(b)

Pri

nt d i

Height (μm)

2 Au contact

rec tion

1

Printed polymer

0 20 40 Substrate

60 μm

Au contact

FIGURE 5.4.4 (a) Optical micrograph of an array of OTFTs with printed polymer PQT-12 on gold source/drain contacts defined using wax printing and etching. (b) Topography of a 0.45% solution of F8T2 printed on gold contacts acquired by atomic force microscopy in tapping mode. A drop size of 35 μm and overlap ratio of 50% were defined in the printing program. The line-to-line overlap of the printed polymer is clearly seen in the corrugation of the surface. (From Paul, K.E. et al., Appl. Phys. Lett. 83, 2070–2072, 2003. With permission.)



425

The TFT data shown by Paul and colleagues for printed semiconductors conform well to conventional transistor models in the linear and saturation regimes, with TFT mobility equal to 0.1 cm2/Vs, Ion/Ioff ratio > 106, threshold voltage of –3 V, and subthreshold slope of 1.7 V/decade. Extracted linear mobility for the printed F8T2 on OTS treated gate dielectrics had an average value of 4 × 10–3 cm2/Vs and a subthreshold slope of ~1 V/decade for coplanar device geometry. All the data are within the range of values that were observed in the equivalent spin-cast devices, indicating that the use of additive inkjet printing for depositing polymeric organic semiconductors does not compromise device performance [1,13,15]. It has been shown that in order to achieve high device performance, the surface of the gate dielectric layer needs to be treated with monolayers that create a hydrophobic coating [15–17]. However, it is difficult to coat polymer semiconductors uniformly on hydrophobic gate dielectrics by spin-coating. Therefore, direct marking also simplifies processing on hydrophobic substrates, allowing the deposition and patterning of polymeric semiconductors onto optimized gate dielectric surfaces and achieving high-performance OTFTs in one processing step. The process used by Paul et al. can also be used to fabricate TFTs with patterned gate electrodes instead of using a doped silicon wafer as a common gate [8]. The processing steps to fabricate a bottom-gate TFT structure using inkjet printing as the only patterning technique are described in Figure 5.4.5. First, a 100-nm chromium film was deposited onto glass by thermal evaporation and patterned by digital lithography using wax ejected from a multi-ejector piezoelectric printhead [14]. The dielectric material was 30 nm of silicon oxide on a 200-nm silicon nitride film, Wax

(1)

Define S/D

(4) Cr

Substrate Define Cr Gate Deposit Semiconductor H25C12

(2) S

S S

(3) Wax

S

n

C12H25

Cr/Au (5)

Si3N4/SiO2

FIGURE 5.4.5 Process flow for the fabrication of polymer TFTs by inkjet printing. Steps 1 to 4 use a subtractive process for patterning of materials. The subtractive process prints an etch mask on a previously deposited film. Step 5 illustrates the additive inkjet printing process. The additive process simultaneously deposits and patterns the semiconductor material PQT12. (From Arias, A.C. et al., Appl. Phys. Lett. 85, 3304–3306, 2004. With permission.)


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deposited using plasma-enhanced chemical vapor deposition (PECVD) at 350°C. The source-drain metal layer, 100-nm gold film with a chromium adhesion layer, was deposited by thermal evaporation and subsequently defined by digital lithography. The surface of the dielectric layer was modified by deposition of a selfassembled monolayer of octyltrichlorosilane (OTS-8) before the semiconductor material was printed on the channel region of the TFTs. An additive inkjet printing process, using a separate print head, was used to simultaneously deposit and pattern the solution-based polymeric semiconductor, completing the bottom-gate TFT device fabrication. Two different print head technologies were used for depositing the semiconductor layer. The first, used for printing small areas, is a nozzleless acoustic inkjet print head that produces features of 30–40 μm of PQT-12 on an OTS-treated surface [12]. The second is a commercially available MicroFab Technologies piezoelectric print head. Feature sizes of 80–100 μm were obtained by using a 60-μm nozzle. The size of printed features obtained with this head was comparable to the channel area of the TFTs and one drop of PQT-12 per transistor was sufficient to give good TFT performance. Figure 5.4.6(a) shows typical transfer and output current-voltage characteristics for a TFT taken from an array of TFTs. The devices exhibit a field-effect mobility between 0.05 and 0.10 cm2/V·s and an off current of ~10–12 A, giving an on/off ratio of ~106 at VSD = –40 V. The onset voltage is close to 0 V and the subthreshold slope is 75 nF.V/decade.cm2. The output characteristics show good saturation and no sign of significant contact resistance. In addition, TFTs made with PQT-12 have shown minimal gate bias stress effect [18]. All the parameters obtained for the inkjet printed TFTs are similar to those obtained for spin-coated materials on thermal oxide-coated silicon substrates [14,15]. The probe test across the diagonal of the arrays shows a high yield of working TFTs and uniformity in both the on- and off-currents. A typical series of measurements, shown in Figure 5.4.6(b), gave an average mobility of 0.06 cm2/V·s, with a standard deviation of 0.02 cm2/V·s. The deviation showed here, using inkjet printing as the only patterning technique, is very promising and comparable to the deviation obtained by the photolithographic process used by Gelinck et al. [19]. As with any solution-based technique, materials compatibility is a challenge requiring device design to prevent the solvent used in successive layers from dissolving previous layers. Using digital lithography to pattern wax resist for etching metal contacts circumvents the issue of dissolution of previous layers, but further materials engineering or device design may be required when inkjet printing organic semiconductor inks on polymeric substrates or dielectric layers. For example, TFTs must be passivated, and this is also a process that can be done by inkjet printing. It has been shown that deposition of the semiconductor and the passivation can be performed in a single-step process by printing a polymer blend [20,21]. Instead of conventionally depositing the semiconducting polymer solution directly over the channel region of the TFT, the semiconductor is blended with an insulator polymer and the blend is deposited from solution over the TFT structure. Viscosity, solubility, and surface energy are manipulated to ensure phase segregation in a direction vertical to the substrate. The final thin film consists of a self-organized double layer, formed in one step deposition, where the semiconductor segregates to the substrate forming the TFT channel and is encapsulated by the insulating polymer.



427

(a)

(b)

VDS = –40 V

μ~0.063cm2/Vs sd 0.021 cm2/Vs

0.12

1E-6 0.10

ID(A)

1E-8 1E-9

ID(A) –6 1.0 × 10

μlinear(cm2/Vs)

1E-7

VG = –40 V

–7 8.0 × 10

1E-10 1E-11 1E-12

–7 6.0 × 10 –7 4.0 × 10 VG = –20 V

–7 2.0 × 10

1E-13 –40

0.0 0

20

30

40

0.06 0.04 0.02

VG = –10 V 10

0.08

Vd(V)

0.00 –30

–20

–10 VG(V)

0

10

0

5

10

15

20

25

30

35

Transistors

FIGURE 5.4.6 (a) Transfer and output (inset) characteristics of a typical printed TFT array with W/L ~ 2 and C = 25 nF/cm2. (b) TFT mobility obtained for different transistors of a printed array, showing an average value of 0.06 cm2/Vs. (From Arias, A.C. et al., Appl. Phys. Lett. 85, 3304–3306, 2004. With permission.)

Figure 5.4.7 demonstrates successful encapsulation of the TFTs via surfaceinduced phase separation. Devices were left in air and consecutive measurements were taken every eight hours. The device characteristics are very stable for a period of two days operating in air. The transfer characteristics, shown in Figure 5.4.7(a), do not show any change during that period of time, illustrating an improved environmental stability of polymer-based devices. Devices made with the other blend compositions and with P3HT instead of PQT show similar trends [20,21]. A small shift in the onset voltage and subthreshold slope is observed when devices are left in air for longer periods of time. The transfer characteristics of a device left in air for 20 days are shown in Figure 5.4.7(b). There was a 7-V shift of the onset voltage and the subthreshold slope increased by 0.7 V/decade. The shift in the onset voltage of devices fabricated from a blend solution in 20 days is 50% less than the shift observed in one day for nonencapsulated devices, and is similar to films separately encapsulated with PMMA [20,21]. This is one example of possible simplification of the processes that can be expected from the inkjet printing technology.

5.4.4 DISPLAY BACKPLANE FABRICATION A typical active matrix thin film transistor (AM-TFT) backplane contains one transistor per pixel in a structure having four to six separate layers composed of conductors, dielectrics, and semiconductors. The TFT works as a switch with the on and off states controlling when each pixel can be addressed. Each pixel is connected to the other through the addressing lines, gate line, and data line. The gate lines turn the TFT on or off while the data (image to be displayed) are transmitted by the data


428


(b) (a) Time exposed to air

1E-6

8h 24h 48h

1E-6

1E-8

ID(A)

ID(A)

1E-8 1E-10

1E-10 1E-12

1E-12

1E-14 –40

1E-14

–20

0 VG(V)

20

–40

as prepared 1.16 V/dec Von~ 0V 20 days in air 1.72V/dec Von~ 7V μ = 0.06 ± 0.02 cm2/V.s –20

0

20

VG(V)

FIGURE 5.4.7 Improved environmental stability of TFTs fabricated with self-encapsulated polymer films. (a) Transfer characteristics of a TFT exposed to air for two days. (b) Transfer characteristics of TFT as prepared and after been exposed to air for 20 days. (From Arias, A.C. et al., Adv. Mater., 18, 2900–2904, 2006. With permission.)

lines. Given the previously demonstrated ability to print polymer TFT arrays, it is important to verify whether inkjet printed features can lead to pixel designs that meet the application needs. Figure 5.4.8 shows a printed 128 × 128 pixel array, along with optical micrographs of the pixels and the equivalent circuit for the pixel. The metal lines were defined by digital lithography and the semiconductor was directly deposited by inkjet printing, as described in Section 5.4.3. The pixel is 340 μm, corresponding to display backplanes with resolution of 75 dots per inch (dpi) [8]. The TFT contacts are formed by the data line and the top plate of the storage capacitor. The metal line width of about 50 μm was determined by the size of the wax drop from the inkjet printer. However, it is shown in Figure 5.4.8(c) that the space between the data line and the storage capacitor plate is smaller than 10 μm. The ability to control high-resolution gaps is important in order to obtain short TFT channel lengths and low parasitic capacitances. The channel length of 30–50 μm was chosen such that the contact resistance that may be observed for polymer-based TFT with short channel lengths is negligible [17,22]. The printed polymer is precisely confined to the region of the gate electrode between the TFT contacts, as seen in Figure 5.4.8 [8]. The extension of the polymer beyond the gate electrode is detrimental to TFT performance because it causes higher leakage current from the ungated semiconductor. The well defined area of inkjet printed PQT-12 also prevents the formation of a continuous layer of polymer between TFTs that would otherwise form a conductive path and result in significant crosstalk between pixels.



(a)

429

(b)

300 μm (c)

(d)

Gate line Data line

CPL Gate line

CP

CST

CST

40 μm

Data line

FIGURE 5.4.8 Optical images of a printed polymer TFT array at increasing magnification, showing the whole 128 × 128 array (a), small regions of the array (b), and a single device (c). Note the printed semiconductor confined to the channel region and figures (b) and (c). (d) Equivalent circuit for the pixel, showing the gate and data address lines, the TFT, and various capacitances. In the display pixel, the storage capacitor, CST, is connected to the next gate line. CP is the parasitic capacitance between the gate line and the pixel and CPL contains all the sources of gate line capacitance, of which the main additional capacitance is to the data line and the TFT channel. (From Arias, A.C. et al., Appl. Phys. Lett. 85, 3304–3306, 2004. With permission.)

The electrical response of a backplane is governed by its RC time constants since the electrical signals sent to the pixels are time dependent [23]. A storage capacitor (CST) is added to the pixel, shown in Figure 5.4.8(c), in order to retain the charge while data are written to the other pixels. The RC time constant for charging a pixel is given by [8]: tRC = RTFT CST = const. CST/[μ (W/L) VGS]

(5.4.1)

The overlap between gate and data levels of a backplane creates a capacitive coupling between the gate line and the pixel. Consequently, the data voltage at the pixel is reduced by an amount equal to the feedthough voltage VFT (Equation 5.4.2), where CP is the parasitic capacitance coupling the gate line to the pixel and VG is the gate voltage relative to the pixel voltage (see Figure 5.4.8c and d): VFT = VGS CP/CST

(5.4.2)

Fast addressing is obtained when tRC and VFT are minimized. These are conflicting requirements since a large value of CST reduces VFT but increases tRC, while the opposite is true for large VGS (see Equations 5.4.1 and 5.4.2). The inverse product of these parameters is a convenient combined figure of merit, F, given by


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F = 1/[tRC VFT] = const. μ [W/L]/CP

(5.4.3)

where larger F corresponds to more optimal performance [8]. An optimal pixel design minimizes the parasitic capacitance CP while maximizing the mobility μ and W/L. The parasitic capacitance is largely due to the overlap between the gate and source-drain metal of the TFT and can be minimized by printing small features or by improving deposition accuracy to ensure careful positioning of the two metals. Low parasitic capacitance and large W/L are easily achieved by conventional photolithography allowing the use of semiconducting polymers with TFT mobilities of 0.02 cm2/Vs to drive electrophoretic displays [19]. Arias et al. combined positioning precision of inkjet printing techniques with the high intrinsic mobility of PQT-12 compared to other polymers to optimize the pixel design of an all-printed patterned backplane. With this printing technology, it was possible to control the placement of metal lines minimizing the parasitic capacitance and obtaining VFT as low as 1.2 V, for operating gate voltage swing of 20 V. Figure 5.4.9 shows a typical pixel response for the printed array. The pixel is charged to the data-line voltage (10 V) during the gate-ON pulse (Vgate-ON = –5 V). After the gate is switched off (Vgate-OFF = +15 V), the charge level drops slowly over several seconds, which indicates good transistor and pixel performance. The gate feedthrough voltage for the chosen voltage levels is around 1–2 V in this pixel design [24]. 16 14

Gate signal Gate OFF

12 10

Pixel signal Gate feedthrough

Voltage (V)

8 6 4 1 Sec 2 Data signal

0 –2 –4 –6

Gate ON Time (500 msec increments)

FIGURE 5.4.9 The pixel signal from a printed backplane shows good switching time, pixel charge storage, and low gate feedthrough. (From Daniel, J.H. et al., SID Symp. Dig. Tech. Papers 36, 1630–1633, 2005. With permission.)



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Another important parameter to consider when choosing materials to fabricate TFT backplanes is the resistance of the address lines. The time constant to address a line is tADR = 0.4 N2 RPL CPL

(5.4.4)

where RPL and CPL are the total resistance and capacitance, respectively, of the metal line for one pixel and N is the number of address lines [23]. A small time constant is easy to achieve in a small format array, but the N2 dependence on the number of lines quickly increases the time constant. Typical values are CPL ~ 1–5 pF and RPL ~ 1–10 Ω for a conventional metal or RPL ~ 10–100 kΩ for conducting polymers. With conventional metals, the time constant remains small enough, around 20 μs, even for arrays where N = 1,000, while conducting polymers could only address arrays where N ~ 30 [8]. Therefore, the development of solution-processed low-resistance conductors is fundamental for an all-additive printing process.

REFERENCES 1. Sirringhaus, H. et al., High-resolution inkjet printing of all-polymer transistor circuits, Science 290, 2123–2126, 2000. 2. Kawase, T. et al., Inkjet printing of polymer thin film transistors, Thin Solid Films 438–439, 279–287, 2003. 3. Wong, W.S. et al., Amorphous silicon thin-film transistors and arrays fabricated by jet printing, App. Phys. Lett. 80, 610–612, 2002. 4. Paul, K.E. et al., Additive jet printing of polymer thin-film transistors, Appl. Phys. Lett. 83, 2070–2072, 2003. 5. De Gennes, P.G., Wetting: Statics and dynamics, Rev. Mod. Phys. 57, 827–863, 1985. 6. Street, R.A. et al., Jet printing flexible displays, Mater. Today 9, 32–37, 2006. 7. Wong, W.S. et al., Hydrogenated amorphous silicon thin-film transistor arrays fabricated by digital lithography, IEEE Electron Device Lett. 24, 577–579, 2003. 8. Arias, A.C. et al., All jet-printed polymer thin-film transistor active-matrix backplanes, Appl. Phys. Lett. 85, 3304–3306, 2004. 9. Shimoda, T. et al., Inkjet printing of light-emitting polymer displays, MRS Bull. 28, 821827, 2003. 10. Gans, B.-J.D., Duineveld, P.C., and Schubert, U.S., Inkjet printing of polymers: State of the art and future developments, Adv. Mater. 16, 203–213, 2004. 11. Sirringhaus, H. et al., Mobility enhancement in conjugated polymer field-effect transistors through chain alignment in a liquid-crystalline phase, App. Phys. Lett. 77, 406–408, 2000. 12. Elrod, S.A. et al., Nozzleless droplet formation with focused acoustic beams, J. Appl. Phys. 65 (9), 3441–3447, 1989. 13. Brennan, D.J. et al., Polyfluorenes as organic semiconductors for polymeric field effect transistors, Proc. SPIE Int. Soc. Optical Eng. 5217 (Organic Field Effect Transistors II), 1–6, 2003. 14. Ong, B.S. et al., High-performance semiconducting polythiophenes for organic thinfilm transistors, J. Am. Chem. Soc. 126, 3378–3379, 2004.


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15. Salleo, A. et al., Thin-film transistors with chemically modified dielectric interfaces, Appl. Phys. Lett. 81, 4383–4385, 2002. 16. Sirringhaus, H. et al., Two-dimensional charge transport in self-organized, highmobility conjugated polymers, Nature 401, 685–688, 1999. 17. Bürgi, L. et al., Close look at charge carrier injection in polymer field-effect transistors, J. Appl. Phys. 94, 6129–6137, 2003. 18. Street, R.A., Salleo, A., and Chabinyc, M.L., Bipolaron mechanism for bias-stress effects in polymer transistors, Physical Rev. B 68, 085316-1–085316-17, 2003. 19. Gelinck, G.H. et al., Flexible active-matrix displays and shift registers based on solution-processed organic transistors, Nat. Mater. 3, 106–110, 2004. 20. Arias, A.C., Vertically segregated polymer blends: Their use in organic electronics, J. Macromolec. Sci. Part C: Polymer Rev. 46, 103–125, 2006. 21. Arias, A.C., Endicott, F., and Street, R.A., Surface induced self-encapsulation of polymer thin film transistors, Adv. Mater., 18, 2900–2904, 2006. 22. Street, R.A. and Salleo, A., Contact effects in polymer transistors, Appl. Phys. Lett. 81, 2887–2889, 2002. 23. Tsukada, T., Active-matrix liquid-crystal displays, in Technology and applications of amorphous silicon, ed. R.A. Street, Springer Verlag, Heidelberg, 2000. 24. Daniel, J.H. et al., Flexible electrophoretic displays with jet-printed active-matrix backplanes, SID Symp. Dig. Tech. Papers 36, 1630–1633, 2005.


5.5

Soft Lithography for Fabricating Organic Thin-Film Transistors

Kimberly C. Dickey, Kwang Seok Lee, and Yueh-Lin Loo CONTENTS 5.5.1 Typical Device Structures and Conventional Fabrication Techniques .....433 5.5.2 Stamps for Soft Lithography .....................................................................435 5.5.3 Microcontact Printing (μCP) .....................................................................438 5.5.3.1 Selective Etching .........................................................................439 5.5.3.2 Selective Electroless Plating .......................................................441 5.5.3.3 Selective Chemical or Electrochemical Polymerization.............443 5.5.3.4 Stamp-and-Spin-Cast...................................................................444 5.5.3.5 Other Microcontact Printing Derivatives ....................................445 5.5.4 Nanotransfer Printing (nTP) ......................................................................447 5.5.5 Soft-Contact Lamination (ScL) .................................................................458 5.5.6 Cold Welding .............................................................................................466 5.5.7 Metal Transfer Printing..............................................................................468 5.5.8 Hot Lift-Off................................................................................................469 5.5.9 Micromolding in Capillaries (MIMIC) .....................................................469 5.5.10 Soft-Contact Optical Lithography .............................................................473 5.5.11 Laser Thermal Transfer Printing ...............................................................475 5.5.12 Imprint Lithography...................................................................................475 References..............................................................................................................483

5.5.1 TYPICAL DEVICE STRUCTURES AND CONVENTIONAL FABRICATION TECHNIQUES Organic thin-film transistors (OTFTs) can be classified as bottom-contact or topcontact devices, depending on where the electrodes make contact to the organic semiconductor layer, as shown in Figure 5.5.1 [1,2]. Bottom-contact devices (Figure 5.5.1a) — where the electrodes are prepatterned prior to the deposition of the organic semiconductor layer — are generally known to have poorer device characteristics compared to those of top-contact devices (Figure 5.5.1b), where the electrodes are 433


434


Organic semiconductor Source Source

Drain Dielectric

Drain Organic semiconductor Dielectric

Gate

Gate

Substrate

Substrate

Bottom-contact device

Top-contact device

(a)

(b)

FIGURE 5.5.1 Side view of (a) bottom-contact and (b) top-contact OTFTs.

directly deposited on top of the organic semiconductor [1,2]. Such disparity in device performance is speculated to result from uniformity differences in the organic semiconductor thin film between the two types of device geometries [1,3]. In the case of bottom-contact devices, the organic semiconductor is deposited on two different surfaces: on the dielectric surface and on top of the metal electrodes. Due to surface energy differences, the growth behavior of the organic semiconductor can change dramatically across the dielectric–electrode interface. The discontinuity in grains across the interface can lead to structural disorder, which can hamper charge transfer and in turn limit device performance [1,3]. In contrast, a uniform organic semiconductor film is first deposited on the gate dielectric in top-contact devices. Electrodes are then defined directly on top of the uniform organic semiconductor layer. As a result, the grains are continuous across the charge transport region. The top-contact device structure is therefore desirable to maximize device performance [4–6]. Top-contact devices are frequently fabricated by sequential deposition of functional materials. Specifically, the organic semiconductor and metal electrodes are independently and sequentially deposited through shadow masks to define the active channel regions. While this technique allows highperformance OTFTs to be fabricated, these devices are typically large, with feature sizes limited by the resolution of the shadow masks (25–30 μm) [7]. Smaller features are easily patterned using photolithography, a well-developed technique for fabricating micro- or submicron inorganic semiconductor devices. According to the 2005 International Technology Roadmap for Semiconductors (ITRS 2005), features as small as 70 nm can be patterned using photolithography [8]. Combining photolithography with lift-off provides an effective method for fabricating the electrodes for OTFTs. Photolithography, however, is rarely used to define the organic semiconductor layer because the chemicals used during the subsequent processing steps (etchants, developers, etc.) can cause degradation [9] or delamination [10] of the organic semiconductor. As a consequence, photolithography is typically limited to defining electrodes in bottom-contact devices in which the organic semiconductor is deposited postpatterning. Halik and coworkers subsequently extended this process to pattern conducting polymer electrodes, such as poly(3,4-ethylenedioxythiophene) doped with polystyrene sulfonic acid (PEDOT-PSS) [11]. In this process, PEDOT-PSS is uniformly coated on a substrate that has prepatterned photoresist. The PEDOT-PSS-coated substrate is then immersed in acetone to swell the photoresist underneath the



435

PEDOT-PSS overlayer. The swollen photoresist lifts off along with the PEDOT-PSS overlayer, leaving behind patterns of PEDOT-PSS in the regions that did not originally contain any photoresist. PEDOT-PSS source, drain, and gate electrodes as small as 2 μm were fabricated in this fashion. Typical of polymer conductors, these PEDOT-PSS electrodes exhibit an electrical conductivity of around 0.1 S/cm [11]. A uniform, continuous pentacene (p-type organic semiconductor) layer is then deposited on the PEDOT-PSS source and drain electrodes to complete the bottom-contact OTFTs. Since pentacene is deposited in the final step, it is not exposed to any of the solvents (e.g., acetone) required to prepattern the PEDOT-PSS electrodes. These bottom-contact pentacene TFTs with PEDOT-PSS electrodes exhibit charge-carrier mobilities comparable to charge-carrier mobilities of typical bottom-contact pentacene TFTs with gold electrodes (0.2 cm2/Vs) [1] and on/off current ratios as high as 106. Further, these device characteristics’ are on par with requirements for driving most display applications [12,13]. As mentioned previously, patterning the active semiconductor layer photolithographically is complex since many organic semiconductors tend to degrade during this process. Despite such technical difficulties, Sheraw and coworkers [14] recently demonstrated a derivative photolithographic patterning proven to define the active pentacene layer. In this process, water-soluble, photopatternable polyvinyl alcohol (PVA) is used as the etch resist for patterning pentacene. Device fabrication begins with depositing and patterning nickel gate electrodes, SiO2 gate dielectric, and palladium source and drain electrodes on a polyethylene naphthalate (PEN) substrate using ion sputtering, and photolithography followed by lift-off. Pentacene is subsequently deposited by thermal evaporation. PVA is then spin-coated directly on the pentacene layer and is photolithographically patterned to create features that define the channel regions. Subsequent oxygen plasma etching removes pentacene outside the channel (regions that are no longer protected by PVA). Bottom-contact pentacene TFTs with a channel length of 10 μm and a channel width of 200 μm exhibit charge-carrier mobilities as high as 1.2 cm2/Vs, and on/off current ratios of 108 [14]. While this process successfully patterns pentacene active regions, it has not been successfully extended for patterning other organic semiconductors, such as poly-3-hexylthiophene (P3HT), due to mechanical cracking of the PVA etch barrier during development. While photolithography has successfully been demonstrated for OTFT fabrication, it is potentially an expensive process [15]. This process is thus not well-suited for fabricating low-cost organic electronic devices. Inexpensive, noninvasive processing technologies that are physically and chemically compatible with electricallyactive organic materials need to be developed to lower fabrication costs and maximize device performance. Among a number of promising techniques that have been demonstrated to date, soft lithography and its derivative techniques will be reviewed in this chapter.

5.5.2 STAMPS FOR SOFT LITHOGRAPHY The key component of soft lithography is the stamp that is used to transfer patterns. To achieve high-quality pattern transfer, conformal contact must be established


436


between the stamp and the substrate, preferably without external pressure. Such conformal contact is often achieved with the use of an elastomeric stamp, such as one made of crosslinked poly(dimethylsiloxane) (PDMS). Dow Corning offers a commercially available PDMS formulation known as Sylgard 184 [16]. Sylgard 184 is a two-component system that consists of a PDMS prepolymer and a cross-linker. To create PDMS stamps, the prepolymer and cross-linker are mixed in a 10:1 (w:w) ratio and degassed. The degassed mixture is cast and cured against a prepatterned master as shown in Figure 5.5.2. The raised and recessed features of the master are typically defined by photolithography. After curing the prepolymer mixture overnight at room temperature, or at 60°C for 2 h, the PDMS stamp is peeled from the master. To ensure that the PDMS stamp releases effectively, the master is typically pretreated with tridecafluoro(1,1,2,2,tetra-hydrooctyl)-1-trichlorosilane (FSAM) [17]. The final stamp has the negative image of the master.

Fabricate and silanize master

SiO2, Si3, N4, metals, photoresists, or wax Si

Pour PDMS prepolymer over master

PDMS

Si

Cure, peel off PDMS

PDMS h l

d

FIGURE 5.5.2 Schematic of the PDMS stamp fabrication procedure. A master is fabricated by conventional photolithography and lift-off. The master is then treated with FSAM for ~30 min to ensure that the cured PDMS releases easily from the master. The mixture of PDMS prepolymer and cross-linker (Sylgard 184, Dow Corning) is poured over the master and cured. After peeling from the master, the PDMS stamp has the negative image of the master. (From Y. Xia and G. M. Whitesides, Angew. Chem., Int. Ed., 37, 550, 1998.)



437

Although photolithography is necessary initially to create the master from which the stamp is cast, once the master is prepared, many stamps can be fabricated from a single master. A stamp with the positive image can also be fabricated by using the initially generated PDMS stamp (with the negative image) as the master. To ensure that the second-generation PDMS stamp releases effectively, the PDMS master can also be treated with FSAM. While the mechanical properties of PDMS are attractive for establishing conformal contact between the stamp and the substrate, the low elastic modulus and high compressibility of the Sylgard 184 PDMS formulation limit the minimum feature size achievable to several microns [18,19]. To realize nanoscale features, a high-resolution hard stamp is required. This hard stamp can be fabricated from rigid materials, such as fused silica, silicon, or GaAs [20,21] or from composite PDMS (h-PDMS) materials [22,23]. When a rigid stamp is used, the nanometer-sized features on that stamp are defined lithographically with techniques, such as electronbeam lithography or holographic lithography [24]. The surface of the rigid stamp is made nonstick with FSAM treatment [21]. To transfer patterns at ambient pressures with a rigid stamp, the substrate being patterned must be a conformal material, such as PDMS, or a plastic substrate coated with PDMS. To eliminate pressure concerns, elastomeric, composite PDMS stamps can be fabricated. These composite stamps were initially developed by Schmid and coworkers at IBM-Zurich [23]. They formulated a siloxane polymer composite, referred to as h-PDMS (“hard” PDMS), that employs a trilayer stack (shown in Figure 5.5.3) to create conformal yet high-resolution elastomeric stamps. The composite stamp consists of a flexible glass or foil backplane to prevent long-range pattern distortion, a soft cushion of Sylgard 184 PDMS to establish conformal contact with the substrate, and a hard silane polymer optimized by Schmid and coworkers to enable nanometer feature replication (as low as 80 nm) at aspect ratios (Figure 5.5.2) ranging from 5 to 0.02. The flexible backplane, as compared to a rigid backplane, allows easier release of the stamp from the master during preparation and the stamp from the substrate after patterning. Further, as shown in the bottom half of Figure 5.5.3, the flexible backplane permits conformal contact with substrates with moderate surface topography. Odom and coworkers [22] subsequently modified the design of the h-PDMS stamp to enable 50-nm feature replication in a two-layer stamp. Specifically, the backplane is completely eliminated in favor of a thicker (~3 mm) Sylgard 184 PDMS layer that supports a thin layer of h-PDMS (30–40 μm). Eliminating the glass backplane makes the composite stamp easier to handle and to release from the master. In addition to the mechanical limitations discussed, PDMS is highly susceptible to solvent swelling, which can affect feature resolution and fidelity during patterning. Recently, DeSimone and coworkers [25–27] developed photocurable perfluoropolyethers (PFPEs) that can be used as new stamp materials for soft lithography. The mechanical properties of PFPEs are similar to PDMS (i.e., conformal contact can be established at ambient conditions). PFPEs, however, are highly resistant to solvent swelling. Further, the low surface energy of PFPEs allows replication of nanoscale features without resorting to composite or hard stamps [28]. PFPE stamps are


438


(a)

Hard backplane Soft cushion

Stiff pattern

(b)

FIGURE 5.5.3 (a) Schematic of a trilayer h-PDMS stamp, which consists of a glass backplane, a PDMS cushion layer, and a patterned h-PDMS layer. (b) Trilayer h-PDMS stamp in contact with an uneven substrate. (From T. W. Odom et al., Langmuir, 18, 5314, 2002.)

fabricated by the same procedure as PDMS stamps, with the exception that the material cures by UV irradiation (365 nm), rather than by heat treatment.

5.5.3 MICROCONTACT PRINTING (μCP) Microcontact printing (μCP) [17,29,30] is part of a set of nonphotolithographic fabrication techniques known as soft lithography. In μCP, features are patterned with inexpensive, elastomeric PDMS stamps (see Section 5.5.2 for details on stamp fabrication). Consequently, μCP can be less expensive in terms of capital and operation costs compared to photolithography for patterning large-area micron- or submicron-sized, features [31,32]. The PDMS stamp is used to transfer inks onto rigid inorganic substrates or flexible polymer substrates. The transferred ink can serve as a molecular resist for subsequent selective etching, or as a hydrophobic/hydrophilic molecular template that allows the selective deposition of functional materials in subsequent steps. To transfer patterns, the PDMS stamp is inked with a solution typically containing molecules or colloids. The inked stamp is then brought into contact with the substrate onto which the ink is transferred. Through van der Waals forces, the raised regions of the stamp and the substrate form conformal contact. In the regions of contact, the ink binds chemically or physically to the substrate. Similar to photolithography, subsequent processing steps are required to define features. For this reason, μCP is also typically limited to the fabrication of electrodes in bottom-contact OTFTs. The subsequent processing steps can be classified into three categories: selective etching [29,33–35], selective electroless plating [36–39], and selective chemical and electrochemical polymerization [40–43]. Depending on



439

PDMS (a)

1−2 μm

Photoresist pattern

Si Exposure of PDMS to thiol solution (b)

PDMS

(c)

PDMS

Gold (50–200 nm) supported on 10 nm Ti adhesion layer (d)

Alkanethiol Stamping onto The thickness of gold substrate the alkanethiolate layer is 12Å Si Etching in CN−/O2

(e)

Si

FIGURE 5.5.4 Defining gold patterns by μCP followed by selective etching. (a) A PDMS stamp is cast from a master. (b) The PDMS stamp is removed from the master and inked with an alkylthiol solution. (c) The inked-PDMS stamp is brought into contact with a goldcoated Si substrate. (d) In the regions of contact, the ink transfers to the gold. (e) The ink on the gold effectively acts as an etch resist. Subsequent etching in an aqueous basic solution removes gold only in the regions that are not protected by the ink. (From A. Kumar and G. M. Whitesides, Appl. Phys. Lett., 63, 2002, 1993.)

the processing details, the final features are created subtractively or additively. Each of these techniques is reviewed in following sections.

5.5.3.1 SELECTIVE ETCHING Kumar and Whitesides [29] first patterned gold features with μCP followed by selective etching. In this example, illustrated in Figure 5.5.4, microcontact-printed alkylthiols act as an etch resist for the subtractive patterning of gold. The procedure begins with inking the PDMS stamp with an alkylthiol solution, and contacting the inked stamp against a gold-coated substrate. In the regions of contact, the thiol end groups form covalent gold–sulfur bonds with the substrate. The covalently bound alkylthiol self-assembled monolayer (SAM) acts as an etch resist, effectively protecting the underlying gold during subsequent etching. Etching the exposed gold regions with an aqueous basic solution results in gold features only in the SAMcovered regions. Gold features, as small as 1 μm, can be produced by μCP followed by selective etching. In addition to patterning gold, this technique has been extended


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to patterning silver [33], copper [34], palladium [35], and indium tin oxide (ITO) [44] features. Rogers and coworkers [44] used μCP and selective etching to fabricate gold source and drain electrodes, as well as interconnects, for active matrix backplane circuits to drive flexible electronic paper displays. Since μCP requires subsequent etching to define features, source and drain electrodes were first defined prior to the deposition of chemically and mechanically fragile organic semiconductor. The resulting OTFTs were therefore built in the bottom-contact geometry. The circuits were fabricated on a flexible, prepatterned indium tin oxide (ITO)-coated poly(ethylene terephthalate) (PET) substrate. Specifically, ITO served as the gate electrode and organosilsesquioxane spin-on glass (SOG) served as the gate dielectric. Thin layers of Ti (1.5 nm) and Au (20 nm) were deposited on the dielectric layer and patterned by μCP and selective etching to define the source and drain electrodes and interconnects. The typical channel width and channel length in these μCP-patterned circuits were ~200 and ~20 μm, respectively. By minimizing the mechanical distortion of the PDMS stamp during printing, Rogers and coworkers were able to limit the registration error between the source and drain level and the gate level to as little as 50 μm over a 16 cm × 16 cm printed footprint [44]. Subsequent deposition of the organic semiconductor layer through a shadow mask completed the bottom-contact OTFTs, which are shown in Figure 5.5.5. A variety of organic semiconductors were used as active materials, and their device performance was comparable to those of typical OTFTs fabricated with bottom-contact metal electrodes defined by conventional photolithography. Another attractive feature of μCP is its ability to pattern features on curved surfaces [45,46]. Jackman and coworkers demonstrated the versatility of μCP by patterning gold features on a cylindrical surface by μCP and selective etching [45].

100 μm

FIGURE 5.5.5 Photograph of a plastic active matrix backplane circuit. Gold source and drain electrodes and interconnects were fabricated by μCP and selective etching. An optical micrograph of a typical transistor is shown in the inset. (From J. A. Rogers et al., Proc. Natl. Acad. Sci. U.S.A., 98, 4835, 2001.)



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In their process, a glass or silicon dioxide cylinder is coated with titanium and gold. The cylinder is then rolled across an alkylthiol-inked PDMS stamp. Since the PDMS stamp makes conformal contact with the curved surface of the cylinder, the alkylthiols are transferred to the metal-coated cylinder with minimal pattern distortion. The final features are generated after an etching step. Alternatively, a cylindrical PDMS stamp can be used to transfer patterns onto a flat gold/titanium-coated substrate [47]. By mounting the PDMS stamp on a glass cylinder, high-speed, reel-to-reel fabrication methods are feasible. Rogers and coworkers demonstrated such a reel-to-reel process for fabricating gold electrodes for bottom-contact regioregular poly(3-hexylthiophene) (P3HT) TFTs using a cylindrical PDMS stamp [48]. These bottom-contact P3HT TFTs exhibited an average charge-carrier mobility of 0.02 cm2/Vs, which is comparable to those of typical bottom-contact P3HT TFTs (from 0.01 to 0.1 cm2/Vs) [1]. An on/off current ratio of 10, however, was extracted from these OTFTs. This low on/off current ratio likely stems from a high off current due to leakage since the organic semiconductor layer was not patterned. Since μCP relies on conformal contact between the stamp and the substrate and the transfer of molecular inks from the stamp to the substrate, the mechanical distortion of the PDMS stamps [49] and diffusion of molecular inks [50,51] can be problematic when printing fine features with high resolution. Controlling these factors is therefore crucial for defining submicron channels. Leufgen and coworkers [52] fabricated bottom-contact OTFTs with channel lengths as small as 100 nm using μCP followed by selective etching. To achieve submicron features, they carefully controlled the contact pressure, the contact time, and the concentration of the molecular inks during μCP. These parameters directly affect how much ink (etch resist) is transferred to the substrate during printing. In Leufgen’s process, the PDMS stamp was supported on a flat, rigid, silicon backplane, so pressure had to be applied uniformly across the stamp. Increasing contact pressures and/or times broadened the stamped regions. This ultimately resulted in narrower channels after etching. Changes in the ink concentration, however, did not have a significant effect on pattern size for a given contact time (2 sec). Figure 5.5.6 shows how the channel length (etched region) can be reduced from 700 to 200 nm simply by increasing the printing pressure from 2 to 4 bar while holding all other conditions constant. While this method provides a useful tool to study short-channel OTFTs, it is doubtful that it can be extended to reproducibly fabricate more complex features.

5.5.3.2 SELECTIVE ELECTROLESS PLATING While μCP and selective etching can potentially replace costly photolithographic tools in OTFT fabrication, expensive deposition tools are still required to deposit the metal layers. Recently, a lower cost metal electroless plating deposition technique has been explored as a potential replacement for the more expensive sputtering and evaporation processes. During electroless plating, metal is deposited onto the desired surface from solution through an autocatalytic redox process [36]. Specifically, a reductant in the


442


Au

SiO2/Si Low p Stamp

2 bar 10s Substrate

1.3 μm 1.8 μm

0.7 μm

0.2 μm High p Stamp

4 bar 10s

Substrate

FIGURE 5.5.6 Submicron channel devices can be fabricated by μCP and selective etching by controlling the pressure during contact. After etching, the channel length is reduced from 700 to 200 nm by increasing the stamping pressure from 2 to 4 bar. (From M. Leufgen et al., Appl. Phys. Lett., 84, 1582, 2004.)

electroless plating solution reduces the cation of the metal, thereby depositing elemental metal on the surface. The surface, however, needs to be activated by a catalyst, such as palladium, to initiate metal deposition. The activated surface can be created by microcontact printing the catalyst on the substrate surface. The combination of μCP and electroless plating has successfully generated copper [36,37] and nickel [38,39] features. For example, Hidber and coworkers [36] fabricated copper electrodes by μCP and electroless plating. Glass, silicon/silicon dioxide, and polymer substrates were patterned with palladium colloids by μCP and submerged in an electroless plating solution. During electroless plating, copper growth only occurs in the regions containing palladium colloids. Using a derivative technique, Zschieschang and coworkers [39] microcontact printed FSAMs [30] on a hydrophilic flexible polyethylene naphthalate (PEN) substrate. The patterned substrate was then activated in a palladium bath; palladium catalyst only adsorbed in the hydrophilic, unstamped regions. Subsequent electroless plating resulted in the selective deposition of nickel only in the hydrophilic regions and not in the regions, stamped with FSAM. The electroless-plated nickel electrodes served as the gate electrode in flexible OTFTs. Solution-processable polyvinylphenol (PVP) was then deposited on top of the nickel electrodes as the gate dielectric. Gold source-drain electrodes in these devices were subsequently defined by photolithography. Finally, pentacene was thermally evaporated to complete the bottom-contact OTFTs. These OTFTs exhibited a charge-carrier mobility



443

of 0.06 cm2/Vs, a subthreshold swing of 1.1 V/decade, and an on/off current ratio of 106. These characteristics are typical of bottom-contact pentacene TFTs with source and drain electrodes defined by ink-jet printing or other soft lithography techniques [39]. Electroless plating, however, is limited to bottom-contact device applications since the electroless plating solutions are generally not compatible with organic semiconductors.

5.5.3.3 SELECTIVE CHEMICAL POLYMERIZATION

OR

ELECTROCHEMICAL

An alternative candidate for source and drain electrodes is conducting polymers. To define conductive polymer electrodes, μCP is used to create hydrophilic and hydrophobic patterns (e.g., a molecular template), which can facilitate the selective deposition of conductive polymer [40–43]. Huang and coworkers [40] demonstrated the selective deposition of polypyrrole and polyaniline (PANI) conducting polymers on silicon/silicon dioxide and glass substrates by chemical polymerization. In this technique, regions of hydrophilic glass and silicon substrates are rendered hydrophobic by microcontact-printing alkylsilanes on the substrates [30,51]. Immersing the patterned substrates in a monomer/oxidant solution (pyrrole or aniline) results in the polymerization of the conducting polymer. Polymerization and deposition rates of the conducting polymers were found to be much higher in the hydrophobic regions than in the hydrophilic regions. As a consequence, a thicker polymer film is deposited in the hydrophobic regions of the substrate. But the conducting polymer adheres more strongly in the hydrophilic regions. Huang and coworkers were thus able to exploit the difference in adhesion to preferentially remove the polymer from the hydrophobic regions of the substrate. Since the polymer adheres more strongly to the hydrophilic surface, the thicker polymer layer can be easily peeled from the hydrophobic regions of the substrate with Scotch™ tape. Using this technique, conducting polymer features with lateral dimensions as small as 2 μm can be produced. Similar results can be obtained by electropolymerization [41–43], in which long alkylthiols (>10 carbons in length) are microcontact-printed onto a gold-coated substrate. These long alkylthiols are thought to inhibit interfacial charge transfer, thereby significantly reducing electropolymerization rates [41]. As such, polymer growth only takes place significantly in the gold regions that had not been stamped with alkylthiols. Using this technique, Gorman and coworkers [41] grew polypyrrole features as small as 2 μm with conductivities ranging from 1 to 5 S/cm. With a similar electropolymerization procedure, Parashkov and coworkers [53] selectively grew PEDOT-PSS source and drain electrodes on a gold substrate. Again, an alkylthiol template was microcontact-printed onto a gold-coated glass substrate for the selective growth of PEDOT-PSS. The PEDOT-PSS patterns exhibited conductivities of 1–5 S/cm. Since the underlying substrate (gold) is laterally conductive, the individual PEDOT-PSS patterns must be transferred to an insulating substrate for application (otherwise, all the conducting polymer patterns are electrically shorted by the underlying gold substrate). To transfer PEDOT-PSS patterns from the


444


gold surface, polyimide (PI) was cast and cured directly on the PEDOT-PSS patterned substrate. Peeling the PI substrate removed the PEDOT-PSS features from the gold surface. Bottom-contact OTFTs were fabricated on the PI platform in an “upside-down” fashion. A thin layer of pentacene was evaporated onto the PEDOTPSS source and drain electrodes. Poly(vinyl alcohol) (PVA) was subsequently spincoated on top of the pentacene layer to serve as the gate dielectric. PEDOT-PSS gate electrodes were then screen-printed on the PVA layer to complete the OTFTs. These devices thus adopt a bottom-contact, top-gate device geometry. The OTFTs fabricated by this technique had channel widths from 200 μm to 3.6 mm and channel lengths ranging from 5 to 50 μm. OTFTs with a channel width of 1 mm and a channel length of 10 μm typically exhibited a charge-carrier mobility of 0.02 cm2/Vs, a threshold voltage of –2.9 V, and an on/off current ratio of 13. These values are lower than those expected for pentacene bottom-contact transistors, but may result from leakage through the PVA polymer gate dielectric. Similar to the various μCP techniques described previously, the solutions required for chemical or electrochemical polymerization are generally not compatible with the organic semiconductors of OTFTs. As a consequence, the application of these procedures is also limited to the fabrication of bottom-contact devices.

5.5.3.4 STAMP-AND-SPIN-CAST Expanding the application of conducting polymers as source and drain electrodes in OTFTs, Lee and coworkers recently demonstrated a direct patterning technique for generating conducting polymer electrodes directly on insulating substrates [54]. Consequently, conductive polymer features can be directly patterned on the gate dielectric of an OTFT. Unlike the previously mentioned techniques for patterning conducting polymers, this technique, coined “stamp-and-spin-cast,” does not require any postpolymerization steps. Such direct patterning of a conducting polymer is enabled by the use of a water-dispersible PANI [54]. Stamp-and-spin-cast begins with microcontact printing hydrophobic alkylsilanes on a hydrophilic substrate (–OH terminated), as shown in Figure 5.5.7(a). To fabricate PANI source and drain electrodes, Lee and coworkers microcontact-printed octadecyltrichlorosilane (OTS) on highly-doped silicon with a thermally-grown silicon dioxide overlayer. Subsequent spin-coating of an aqueous PANI dispersion on the patterned substrate creates conductive PANI patterns in the hydrophilic regions. PANI features as small as 5 μm were achieved by Lee and coworkers using this technique [54]. Pentacene was thermally evaporated though a shadow mask on top on the PANI source and drain electrodes to complete bottom-contact devices on a silicon/silicon dioxide platform. Bottom-contact pentacene TFTs with PANI electrodes patterned by stamp-and-spin-cast with channel widths of 800 or 1,000 μm and channel lengths ranging from 30 to 300 μm exhibited an average charge-carrier mobility of 0.016 ± 0.008 cm2/Vs and on/off current ratio as high as 104. The reduced charge-carrier mobility (compared to typical pentacene bottomcontact TFTs) stemmed from the dielectric surface roughness (OTS was stamped on the dielectric surface). Using solution deposition methods to create the molecular templates yielded much smoother dielectric surfaces, and the performance of these



445

UV/Ozone treatment to produce a hydrophilic surface O H O H O H O H O H O HO H

(a)

(b)

Deposit molecules R

Transfer OTS

R

R

R

R

R

R

OTS OTS OTS OH OH OH OH OH OH OH UV/Ozone exposure

Shadow mask

R

R

OH

R OH

R OH

Spin-cast PANI-PAAMPSA aqueous solution PANI-PAAMPSA

R OH

R

R OH

R OH

FIGURE 5.5.7 Schematic of stamp-and-spin-cast. Hydrophobic patterns are created by (a) μCP or (b) subtractively removing hydrophobic molecules by UV-ozone irradiation through a shadow mask. An aqueous PANI dispersion is spin-coated on the patterned substrate. PANI features selectively deposit in the hydrophilic regions immediately after the spin-coating. (From K. S. Lee et al., Appl. Phys. Lett., 86, 074102, 2005.)

devices improved accordingly [54]. These OTFTs with PANI electrodes exhibited linear output characteristics in the low source-drain voltage regime [54]. This observation implies lower contact resistance at the PANI–pentacene interface compared to that at the gold–pentacene interface, and has recently been verified by scanning surface potential measurement experiments [55].

5.5.3.5 OTHER MICROCONTACT PRINTING DERIVATIVES In addition to using μCP to fabricate electrodes for OTFTs as discussed previously, μCP can also be used to pattern the growth of organic–inorganic hybrid semicon-


446


(a)

(b)

FIGURE 5.5.8 Examples of the selective growth of oligoacene crystals on SAM-templated substrates. The dark regions and the bright regions indicate oligoacene and bare gold, respectively. (a) Oligoacene crystals grow on terphenylthiol–SAM-covered squares in the dodecylthiol–SAM-covered background. (b) Anthracene crystals are deposited on dodecylthiol–SAMcovered squares terphenylthiol SAM in the background. Scale bar = 300 μm. (From A. L. Briseno et al., J. Am. Chem. Soc., 127, 12164, 2005.)

ductors [56], organic semiconductor single crystals [57], and inorganic semiconductor single crystals [58]. Kagan and coworkers [56] demonstrated the selective deposition of a soluble organic–inorganic hybrid semiconductor, (C6H5C2H4NH3)2SnI4, on microcontact-printed molecular templates. In their method, hydrophobic inks, such as alkylsilanes, fluorinated alkylsilanes, alkylphosphonic acids, or alkylhydroxamic acids, were microcontact-printed on SiO 2 or ZrO 2 . A solution of (C6H5C2H4NH3)2SnI4 was then spin-coated on the patterned substrate. The semiconductor only adsorbed in the hydrophilic bare oxide regions due to the favorable wetting properties. The ability to pattern the semiconductor layer selectively is useful for building addressable arrays of transistors for display applications. Recently, Briseno and coworkers [57] demonstrated the selective growth of organic semiconductor single crystals on microcontact-printed molecular templates. Specifically, they investigated the growth of oligoacene crystals (anthracene and 5chlorotetracene) from a THF solution on substrates prepatterned with various SAMs. Their results showed that the crystal growth of oligoacenes varied dramatically depending on the chemical functionality of the stamped SAM. In particular, terphenylthiol SAMs induced the growth of large oligoacene single crystals from THF solution, while dodecylthiol-treated surfaces suppressed crystal growth. The use of a template containing terphenylthiol and dodecylthiol patterned regions therefore yielded patterned growth of oligoacene single crystals on the terphenylthiol-printed regions as shown in Figure 5.5.8. Functional 5-chlorotetracene TFTs were demonstrated by patterning dodecylthiol and terphenylthiol on gold source and drain electrodes. These patterned OTFTs exhibited an average charge-carrier mobility of 5.7 × 10–4cm2/Vs, which is comparable to that of single-crystal 5-chlorotetracene devices (1.4 × 10–4cm2/V) [59]. Finally, the selective growth of inorganic crystals was demonstrated on a SAMtemplated silver surface. Hsu and coworkers [58] preferentially grew zinc oxide (ZnO) from an aqueous solution on silver surfaces. The silver surface was patterned by μCP with a carboxyl acid (–COOH)-terminated SAM to inhibit ZnO growth. Figure 5.5.9(b) and (c) contain micrographs that demonstrate ZnO growth only



(a)

447

(b) COOH terminated SAM micropatterns

Ag Si ZnO growth from aqueous solution

2 μm (c) 200 μm

ZnO

Ag Si

FIGURE 5.5.9 (a) COOH-terminated SAM is microcontact-printed on silver. ZnO preferentially grows on the bare silver surface rather than on COOH-terminated surface. (b, c) SEM images reveal that ZnO nanorods (white) only grow in the bare silver regions. The surrounding regions are covered with HSC10H20COOH. (From J. W. P. Hsu et al., Nano. Lett., 5, 83, 2005.)

occurs in the bare silver regions of the substrate. Since ZnO has a wide band gap, potential applications of patterned ZnO microstructures include photovoltaics, active sensor platforms, and microlasers [58].

5.5.4 NANOTRANSFER PRINTING (nTP) While μCP is a highly effective patterning technique, it is still subtractive in nature, requiring etchback to define functional features. Often, the chemicals used for etchback are incompatible with organic semiconductors, thereby limiting the application of μCP to bottom-contact devices. To eliminate the materials incompatibility issues and enable top-contact device fabrication, Loo and coworkers developed an additive patterning technique called nanotransfer printing (nTP) [20,21,60–62]. Nanotransfer printing is a solventless patterning technique for directly printing functional solid materials — these can be metal features or stacks of functional materials — from the raised regions of a stamp onto a substrate. Moreover, nTP occurs at ambient conditions and allows features as small as 100 nm [20] to be patterned over large areas on both rigid and flexible substrates. Using nTP, patterns have been transferred from rigid stamps [21,63] (e.g., GaAs or fused silica) and elastomeric PDMS stamps [20,21,60–62,64,65]. Pattern transfer results from an interfacial chemistry between the substrate and the material being transferred. As a consequence, the contact between the stamp and the substrate is


448


critical for achieving high-quality printed patterns. Using a PDMS stamp or substrate during nTP ensures conformal contact between the stamp and substrate via van der Waals interactions. The intimate contact results in high-quality printing at ambient pressures. If a rigid stamp is used to transfer patterns to a rigid substrate, external pressure (>150 MPa [66]) is required to obtain intimate contact between the stamp and the substrate. This technique has been demonstrated with cold welding of metal features [66–68], and is discussed in Section 5.5.6 of this chapter. This section will focus on systems with at least one conformal surface so that printing can be carried out at ambient pressures. The success of nTP relies on interfacial chemistries between the solid material being transferred and the substrate. For the nTP procedure shown in Figure 5.5.10, the reactivity of titania is exploited for pattern transfer [20]. Thin films of gold (~40 nm) and titanium (~5 nm) are deposited on the raised and recessed regions of a patterned stamp (either PDMS or GaAs) via sputtering or electron beam evaporation. Collimation of the metal fluxes during deposition prevents the sidewalls of the stamp from becoming coated with metal. Exposing the titanium to ambient conditions oxidizes its surface to form titania (TiOx). The gold and titania stack can then be transferred from the raised regions of the stamp onto a plastic substrate on contact. To prepare the plastic substrate for pattern transfer, a thin film of PDMS is spin coated onto a PET substrate to create a conformal, flexible substrate. Before contacting the stamp with the substrate, titania and PDMS are exposed to oxygen plasma to create reactive –OH groups on their respective surfaces [69–71]. Contacting the stamp with the substrate results in conformal contact between the raised regions of the stamp and the substrate. In the regions of physical contact, condensation reactions between the hydroxyl groups of the titania and those on the PDMS surfaces result in permanent covalent bonds. Removing the PDMS stamp (to which gold has poor adhesion) from the plastic substrate transfers the gold/titania metal from the raised regions of the stamp to the substrate. The entire procedure occurs at ambient conditions with contact times less than 15 sec. Scotch tape adhesion tests further confirm that the transferred patterns are strongly bonded to the substrate. Since aluminum also readily forms surface oxides, the same interfacial chemistry can be used to generate aluminum patterns [20]. Figure 5.5.11 illustrates the printing quality of nTP. Scanning electron micrographs of a PDMS stamp with gold deposited on the raised and recessed regions before and after nTP are shown in Figures 5.5.11(a) and (b), respectively. These micrographs clearly illustrate that the metal stack transfers completely from the raised regions of the stamp. Further, Figure 5.5.11(c) demonstrates that the fidelity of the PDMS pattern is replicated on the PDMS/PET substrate. Combining nTP with soft contact lamination (ScL), which will be discussed in Section 5.5.5, allows functional top-contact devices to be built on plastic substrates [20,63]. Loo and coworkers fabricated organic transistors and complementary inverter circuits on PDMS/PET substrates in which the gold source/drain electrodes and interconnects were patterned by nTP [20]. Figure 5.5.12(a) shows the currentvoltage characteristics of a representative pentacene [72] (p-type) TFT fabricated by a combination of nTP and soft contact lamination. The electrical properties of these top-contact devices (charge-carrier mobility ~ 0.1 cm2/V-s, on/off current ratio ~ 104)



449

Stamp

(a)

Deposit Au/Ti on stamp (b) 0.2 − 10 μm 0.1 − 10 cm 0.05 − 100 μm

(c)

(d)

Plasma oxidizes surfaces of stamp, substrate; print

Substrate

Si Ti

O

TiOx chemically bonds to substrate

Remove substrate; Au/Ti printing complete (e)

FIGURE 5.5.10 Schematic of nanotransfer printing (nTP): (a) Stamp with relief features ranging from 0.2 to 10 μm and the lateral dimensions ranging from 0.05 to 100 μm. (b) 20nm Au and 5-nm Ti are sequentially evaporated on the raised and recessed regions of the stamp. (c) The stamp and substrate are exposed to oxygen plasma to create surface (–OH) groups and brought into contact with one another. (d) An interfacial condensation reaction occurs in the regions of contact binding the metal to the substrate. (e) Peeling away the stamp transfers the pattern to the substrate via permanent covalent bonds. (From Y.-L. Loo et al., Appl. Phys. Lett., 81, 562, 2002.)


450


(a)

100 μm (b)

100 μm (c)

100 μm

FIGURE 5.5.11 Scanning electron micrographs of a PDMS stamp with 20-nm Au and 5nm Ti on the raised and recessed regions before (a) and after (b) nTP. (c) The resulting printed metal features on a plastic substrate. (From Y.-L. Loo et al., Appl. Phys. Lett., 81, 562, 2002.)



Drain-source current (μA)

Drain-source current (μA)

2 10 102

0 10 100

-30

Vg = −100 V

−2\-2 1010 0 −50 -50 −100 -100 -20 0 Gate Voltage Gate Voltage −10

0 0

451

40

100 μm m

30

Vg = 100 V

20 10

−20 −40 −60 −80 −100

0 0

20

40

60

80

Drain-source voltage (V)

Drain-source voltage (V)

(a)

(b)

100

Vout (volts)

40 30 20 10 0 0

10

20

30

40

Vin (volts) (c)

FIGURE 5.5.12 (a) Current-voltage characteristics of pentacene thin-film transistors whose gold electrodes were printed by nTP. Inset: saturated source-drain current as a function of gate voltage (W/L = 8). (b) Switching characteristics of a complementary inverter circuit whose gold electrodes were printed by nTP. (From Y.-L. Loo et al., Appl. Phys. Lett., 81, 562, 2002.)

are comparable to top-contact transistors in which the gold electrodes are evaporated through a shadow mask directly on top of pentacene [20]. Figure 5.5.12(b) shows the transfer characteristics of a complementary organic inverter circuit with pentacene (p-type) and hexadecafluoro copper phthalacyanine [73] (n-type) organic semiconductors. Again, the performance of this circuit compares well with top-contact devices in which the electrodes and wiring are fabricated by evaporation through a shadow mask [20]. The nTP process described in the preceding paragraphs relies on interfacial condensation chemistry to transfer metal patterns from a PDMS stamp to a plastic substrate. Realizing that the interfacial chemistry can be modified according to the material to be transferred and the desired substrate, nTP can easily be extended for patterning a variety of single- and multilayer conductors, semiconductors, and dielectrics. In one such variation, SAMs serve as a covalent “glue” for transferring gold from PDMS stamps onto GaAs substrates. Figure 5.5.13 illustrates this concept. Difunctional 1,8-octanedithol molecules are deposited on freshly etched GaAs. The adsorption of the molecules on the substrate occurs such that only one of the thiol endgroups bonds to the GaAs substrate, leaving the second thiol functionality free for subsequent reaction [60].


452


(a) Etch oxide; deposit molecules 20 nm Au PDMS stamp

SH

SH

SH

SH

(CH2)X (CH2)X (CH2)X (CH2)X S

S

S

S

Au

Au

S

S

GaAs

(b) Bring stamp into contact with substrate

(c) Remove stamp; complete nTP

SH

SH

(CH2)X (CH2)X (CH2)X (CH2)X S

S

S

S

GaAs

FIGURE 5.5.13 Schematic of nTP on GaAs. (a) 1,8-octanedithiol molecules are deposited on freshly etched GaAs. Inset: idealized orientation of 1,8-octanedithiol molecules on GaAs. (b) A PDMS stamp with 20 nm of gold evaporated on the raised and recessed regions of the stamp is contacted against the treated GaAs surface. (c) Peeling away the PDMS stamp effectively transfers gold onto the GaAs substrate in the regions of contact. Inset: idealized orientation of 1,8-octanedithiol molecules as “glue” between GaAs and gold. (From Y.-L. Loo, et al., J. Vac. Sci. Technol., B, 20, 2853, 2002.)

An idealized schematic of the orientation of the deposited 1,8-octanedithiol layer is shown in the inset of Figure 5.5.13. Separately, a thin layer (~20 nm) of gold is deposited onto the raised and recessed regions of a PDMS stamp. The stamp is then brought into contact with the thiol-functionalized GaAs substrate. In the regions of conformal physical contact between the stamp and substrate, an instantaneous chemical reaction occurs between the unreacted thiol end groups and gold on the raised regions of the stamp, resulting in permanent gold–sulfur covalent bonds. Removing the stamp leaves behind a gold pattern, as shown in Figure 5.5.14. Since gold is permanently adhered to the substrate through covalent bonds, the printed patterns always pass Scotch tape adhesion tests. Again, this chemistry works at room temperature with contact times less than 30 sec. Further, the interfacial thiol chemistry is expected to work with all coinage metals. In fact, a similar nTP process has been demonstrated for patterning copper [62].



1 μm (a)

100 μm (b)

50 μm (c)

453

100 μm (d)

FIGURE 5.5.14 Optical micrographs of gold patterns printed by nTP. (a) 100-nm holes printed with a GaAs hard stamp on a plastic substrate coated with PDMS; gold patterns printed with a PDMS stamp on (b) a PDMS substrate; (c) a silicon/silicon dioxide substrate; and (d) an ITO-coated plastic substrate. (From Y.-L. Loo, et al., J. Vac. Sci. Technol., B, 20, 2853, 2002.)

There is, however, a significant difference between printed gold and copper patterns. Unlike printed gold patterns, printing with as-cast PDMS stamps results in nonconductive copper patterns. An additional step of leaching the PDMS stamps prior to metal deposition must therefore be instituted to print conductive copper patterns. X-ray photoelectron spectroscopy carried out during depth profiling experiments revealed that uncross-linked PDMS oligomers can penetrate between copper grains, thereby disrupting the lateral conductive pathway. Leaching removes uncrosslinked PDMS oligomers from the stamps. This phenomenon — that PDMS oligomers penetrate between copper grains — appears to be related to whether the metal oxidizes at ambient conditions. Specifically, gold, which does not oxidize at ambient conditions, is always conductive whether it is printed from as-cast or leached PDMS stamps. Copper [62] and silver [74], both of which oxidize at ambient conditions, are only electrically conductive when printed from leached PDMS stamps. In addition to patterning on GaAs, SAMs can also be used to bind gold patterns to silicon substrates. Similar to the procedure for patterning on GaAs, a molecule with thiol functionality, 3-mercaptopropyltrimethoxysilane (MPTMS), is employed. The silicon wafer is first treated (with a 6:1:1 mix of water:H2O2:HCl, for 10 min at 75°C) [21] to create reactive hydroxyl surface functionality (–OH). The methoxy groups of MPTMS react with the –OH functionality to generate a


454

(a)

200 μm


(b)

200 μm

FIGURE 5.5.15 Optical micrographs of (a) silver dots and (b) copper lines printed on silicon by nTP with as-cast PDMS stamps.

thiol-terminated (–SH) silicon surface [21,75]. Contacting a PDMS stamp (with gold deposited on the raised and recessed regions of the stamp) with the thiolfunctionalized silicon surface transfers gold patterns in the regions of contact. Again, the gold patterns are covalently bound to the silicon substrate through permanent gold–sulfur linkages. Similar results, shown in Figure 5.5.15, can be achieved for patterning silver and copper on silicon [62,74]. As noted previously, the PDMS stamps must be leached prior to printing conductive silver and copper patterns. When leached stamps are used, the conductivities of the printed copper and silver features are comparable to evaporated thin films of copper and silver of similar thicknesses. Additionally, the same interfacial silane chemistry can be employed with any substrates that have surface hydroxyl groups or can be functionalized with surface hydroxyl groups. For example, gold patterns can be printed on PET substrates coated with organosilsesquioxane (glass resin) [21]. As discussed in Section 5.5.2, the mechanical properties of commercially-available Sylgard 184 PDMS limit the minimum feature size achievable to several microns [18,19]. Specifically, shallow relief features tend to deform, buckle, or collapse [18,19] when brought into contact with the substrate to be patterned. To generate nanoscale features with nTP, hard stamps must be used. Loo and coworkers generated submicron patterns with GaAs hard stamps, as illustrated in Figure 5.5.16 [20]. Here, gold patterns are transferred at ambient pressures onto a conformal plastic substrate coated with PDMS. The printed patterns are uniform over large areas. Figures 5.5.16(b) and (c) show the edge roughness of the printed features to be less than 15 nm, suggesting that when a stamp with smooth edges is used, the edge resolution of printed features is limited by the grain size of the evaporated metal. While rigid hard stamps are successfully used to print submicron features, elastomeric stamps remain attractive because they can be easily fabricated and they readily conform to rigid substrates at ambient pressures. The low elastic modulus and high compressibility of elastomeric PDMS stamps can be overcome with composite PDMS stamps (h-PDMS) [22,23]. As discussed in Section 5.5.2, h-PDMS stamps consist of a Sylgard 184 PDMS layer that provides conformal contact with



(a)

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(b)

5 μm

1 μm

(c)

100 nm

FIGURE 5.5.16 Scanning electron micrographs of Au/Ti patterns printed onto PDMS/PET substrates. (a) 500-nm lines printed with a fused silica stamp; (b) intersecting trenches printed with a GaAs stamp; and (c) 130-nm hole array printed with a GaAs stamp. The bright regions in the micrographs represent transferred Au/Ti. (From Y.-L. Loo et al., Appl. Phys. Lett., 81, 562, 2002.)

the substrate and a rigid h-PDMS layer that provides nanoscale feature resolution. These h-PDMS stamps have successfully been used by Zaumseil and coworkers to print three-dimensional nanostructures with features as small as 50–100 nm [64]. The ability to print complex multilayer stacks and the ability to separate the deposition and patterning steps effectively make nTP a powerful tool for building devices on plastic. For example, Loo and coworkers demonstrated the versatility of nTP by fabricating metal–insulator–metal capacitors on plastic at ambient conditions [21]. In these devices, SiNx functions as the dielectric. Direct plasma deposition of SiNx on plastic is generally limited by the low working temperature of the substrate. The nTP process, however, can be exploited to transfer SiNx onto plastic at ambient conditions. To fabricate the capacitor structure, layers of Au (50 nm), SiNx (100 nm), Ti (5 nm), and Au (50 nm) are sequentially deposited onto a lithographically patterned silicon stamp treated with FSAM. The multilayer stack is then transferred to a gold-coated plastic substrate via cold welding [66–68] (see Section 5.5.6) when the stamp is brought into contact with the substrate. The performance of the printed capacitors was comparable to that of capacitors of similar dimensions fabricated on silicon wafers by photolithography, lift-off, and etchback [21]. To this point, SAMs have been discussed in the context of glues to anchor metal patterns or metal stacks to a substrate. If we extend our thinking to consider SAMs as the electrically-active components [65,76,77], we see how nTP can be a useful


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tool for building nanoscale electronic devices. Specifically, nTP can provide a nondestructive method for making electrical contact to a molecular layer. In contrast, direct evaporation of metal contacts on top of a molecular layer generally leads to electrical shorts in these devices since hot metal atoms can penetrate and/or damage the molecules during evaporation [78]. Using nTP, Loo and coworkers fabricated GaAs/1,8-octanedithiol/gold molecular junctions by nTP. In these devices, the 1,8octanedithiol is not only the glue that anchors gold on GaAs, but also the electricallyactive layer between two electrodes (gold and GaAs) [65]. The electrical properties of 1,8-octanedithiol, however, are not very interesting. Recent work in this area focuses on the assembly of conjugated glues [76,79–83]. These conjugated, electrically-active molecules are attractive because they can potentially carry charges along their molecular backbone. Combining the patternability of nTP with electrically-interesting molecules moves us one step closer to understanding the fundamentals of nanoscale electronics and perhaps even the realization of molecular electronics. In addition to the nTP procedures described in the preceding paragraphs, a noncovalent transfer process was recently demonstrated by Hur and coworkers for patterning gold and gold/titanium multilayers [61]. As the name implies, this procedure does not rely on the formation of specific covalent bonds for pattern transfer. Instead, noncovalent surface forces are exploited. The significant advantage of this procedure, however, is the ability to print electrodes directly on top of the organic semiconductor thin films without prior modification or surface treatment. Figure 5.5.17 illustrates this procedure. Similar to previously described nTP procedures, the metal to be transferred is deposited on the raised and recessed regions of a PDMS stamp. The stamp is contacted against the substrate and heated mildly (50–80°C). The heating time required for pattern transfer is a function of the surface energies, and varies depending on the metal being transferred. A high surface energy substrate, like PET or polythiophene, tends to be “stickier” and therefore requires less heating time for pattern transfer (surface energy of PDMS < surface energy of PET) than a lowenergy surface, like pentacene or polypropylene (surface energy of PDMS ≈ surface energy of pentacene). For successful patterning, the nonspecific adhesion at the metal–stamp interface must be less than the nonspecific adhesion between the metal–substrate interface (i.e., the metal must prefer to adhere to the substrate over adhering to the stamp). Since PDMS has an extremely low surface energy, this is typically not a problem with most materials. Removing the stamp from the substrate transfers the metal pattern in the regions of contact. Figures 5.5.17(b) through (e) demonstrate some of the patterns generated with this technique [61]. Deposition of metals onto PDMS stamps by electron-beam evaporation, however, can modify the PDMS surface, thereby increasing the surface energy of the PDMS above its intrinsic value. If the surface energy of the PDMS becomes greater than the surface energy of the substrate, the metal will not transfer from the PDMS stamp. This is typically not a concern with conventional nTP since it relies on specific interfacial chemistries for pattern transfer. Hur and coworkers found that the surface energy of the PDMS stamp can be restored to close to its intrinsic value by instituting a heating step. They hypothesized that this heating step facilitates reorientation and



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(a) Evaporated metal films PDMS Stamp Contact surface

Apply heat Substrate Remove stamp

(b)

(c)

500 nm 1 cm (d)

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1 cm

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FIGURE 5.5.17 (a) Schematic of noncovalent nTP: A PDMS stamp, with metal deposited on the raised and recessed regions, is contacted against a substrate. After moderate heating, the stamp is removed, leaving behind metal in the regions of contact. (b) Optical micrograph of 30-nm thick gold dots printed on a plastic substrate. (c) Scanning electron micrograph of a small region of the same dot array. (d) Optical micrograph of Ti (2 nm)/Au (30 nm) patterns printed on a silicon/silicon dioxide substrate. (e) Optical and scanning electron micrographs of a smaller region of the same printed pattern. (From S.-H. Hur et al., Appl. Phys. Lett., 85, 5730, 2004.)

segmental motion of the polymer chains and/or diffusion of the low molecular weight components of the PDMS to the stamp surface [61]. In addition to restoring surface energy, the heating step may also improve contact between the stamp and the substrate. Heating times required for pattern transfer can also be affected by the surface roughness of the substrate. Overall, smoother surfaces lead to better contact between the stamp and the substrate and therefore higher quality pattern transfer. Hur and coworkers successfully fabricated top-contact P3HT TFTs on plastic substrates with printed multilayer gold/titanium and gold/titanium/gold electrodes


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by noncovalent nTP [61]. Fabrication starts on a polyimide substrate. Gold, gold/titanium, and gold/titanium/gold electrodes are printed directly on top of P3HT to complete the circuit. The authors observed that the multilayer electrodes have fewer cracks and are therefore more electrically conductive. They attribute this observation to the improved structural integrity of the multilayer stack. Further, the performance of all of the printed devices (charge-carrier mobility ~ 3.5 × 10–3 cm2/Vs, for a device with a channel width of 200 μm, and a channel length of 50 μm) is comparable to the performance of devices in which the electrodes were directly evaporated through a shadow mask (charge-carrier mobility ~ 3.0 × 10–3 cm2/Vs, for devices with a channel width of 200 μm, and a channel length of 50 μm) [61]. The most significant advantage of nTP over traditional patterning techniques is that nTP is purely additive. This allows the functional material to be transferred directly onto the desired substrate with exact placement, eliminating the use of photoresists, etchants, and developers that are required for photolithography and microcontact printing (μCP) techniques. Also, since solid functional materials can be directly transferred, ink diffusion concerns associated with μCP are eliminated. Finally, the stamp is reusable (but does require cleaning over time), and the process allows for direct patterning of three-dimensional and multilayer structures without etching or additional processing steps [21,84].

5.5.5 SOFT-CONTACT LAMINATION (ScL) Combining μCP with lamination techniques [85,86], Loo and coworkers fabricated top-contact organic transistor arrays on plastic substrates. This procedure was subsequently called soft-contact lamination (ScL) [87,88]. ScL exploits the conformal properties of PDMS, so individual components of the device can be made separately and then laminated together in the final step. This feature of ScL is particularly useful for assembling OTFTs considering the chemical and mechanical fragility of the organic semiconductor. In SeL, the organic semiconductor can be deposited independently on one conformal substrate (which also contains the gate and gate dielectric), while the source and drain electrodes are patterned on a separate conformal substrate. The two individual substrates are then laminated together to complete the circuit. Consequently, devices fabricated by ScL adopt the top-contact geometry. Since many organic materials are incompatible with the traditional lithographic and μCP processes used to fabricate and pattern the electrical contacts, separating the deposition and patterning steps greatly simplifies the fabrication of plastic electronics and avoids exposing the organic semiconductor to harsh chemical environments. Further, ScL provides a facile method for fabricating embedded transistors that are mechanically flexible. Figure 5.5.18 illustrates the ScL procedure used by Loo and coworkers to fabricate laminated pentacene TFTs. The top substrate, containing the μCP-printed source and drain electrodes, was constructed on a PDMS-coated plastic substrate [89]. The bottom substrate was constructed on an ITO-coated plastic substrate. The ITO, patterned photo lithographically to define the gate level, was coated with an organosilesquioxane spin-on glass, which comprises the dielectric. Pentacene [44]



(a)

(b)

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PDMS PET

Au OH PDMS

Si

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Semiconductor, dielectric gate Contacts PET

(d)

PET

PDMS Dielectric

Interfacial bonding; complete circuit

Si

Si

FIGURE 5.5.18 Schematic of soft-contact lamination. (a) Uniform layers of Ti (~1 nm) and Au (15–20 nm) are deposited onto a PDMS-coated plastic substrate. (b) Gold source and drain electrodes and interconnects are defined by microcontact printing. (c) This sheet is laminated against a bottom plastic substrate that contains the gate, dielectric and, semiconductor levels. (d) The completed laminated circuit. Insets: magnified view of the laminated circuit and the side profile of a laminated electrode. (From Y.-L. Loo et al., Proc. Natl. Acad. Sci. U.S.A., 99, 10252, 2002.)


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was evaporated on the dielectric layer through a metal shadow mask and annealed at 100°C in a nitrogen environment for 6 h to create the organic semiconductor layer. Aligning and contacting the top substrate with the bottom substrate completed the circuit [89]. The PDMS layer on the top substrate allowed conformal contact with the bottom substrate at ambient pressures through van der Waals interactions, yielding a complete circuit embedded between two sheets of plastic without the use of adhesives [71]. The conformal contact provided by the PDMS layer also provided efficient electrical contact between the source/drain electrodes and the organic semiconductor layer. Zaumseil and coworkers demonstrated that contacts formed by lamination actually exhibit lower contact resistance between the electrodes and the organic semiconductor than contacts formed by direct evaporation of gold on top of the organic semiconductor [87]. Lower contact resistance translates to lower power consumption per unit current output during device operation. Additionally, the encapsulated, laminated transistors are mechanically robust. By placing the active elements of the circuits at the zero-stress plane [90] through lamination, the circuits do not fail due to mechanical fracture of the ITO layer (the most brittle material in this system), nor do they fail due to plastic deformation of the PET substrate [89]. The encapsulation that results from lamination also has practical importance since many organic semiconductors are highly sensitive to their environment. While PDMS is not designed to form a hermetic seal, it did provide several hours of protection when the laminated circuit was immersed in soapy water [89]. ScL was also used to study charge transport in organic semiconductor single crystals [91]. Prior to ScL, organic semiconductor single-crystal transistors were fabricated by placing a thin organic semiconductor single crystal (~1 μm) against a silicon wafer with predefined electrodes [92–94]. While this technique is successful, it requires extremely thin, bendable (i.e., fragile) organic semiconductor single crystals to achieve successful lamination between the crystal and the electrodes, since this process relies purely on electrostatics between the crystal and the electrodes. ScL provides two advantages over this technique. Since the metal electrodes supported on PDMS substrates are conformal to moderate surface topography, thicker (up to a few millimeters) and more rigid organic semiconductor single crystals can be studied. Further, the ScL technique is nondestructive and reversible, so the laminated contacts can be removed from the organic semiconductor single-crystal surface and re-established many times without affecting the transistor characteristics [91]. To contact the organic semiconductor single crystal, the source, drain, and gate electrode levels are all fabricated independently on a PDMS substrate [91,95]. The PDMS substrate containing the patterned gate, source, and drain electrodes is then laminated against the organic semiconductor single crystal to complete the circuit (Figure 5.5.19). This capability allows many transistor structures — each with different electrode configurations and dimensions — to be assembled and characterized on the same region of the crystal. Figure 5.5.20(b) shows data collected from a set of transistors with different channel lengths assembled sequentially on a single region of a rubrene single crystal. As expected, the saturation current scales linearly with the width to length ratio (W/L) of the transistor channel [91]. Additionally, the ScL technique has allowed Sundar and coworkers to study how molecular anisotropy



Drain

Source

461

Top view

Dielectric Gate 500 μm Substrate

(a) Organic crystal

100 μm

Initiate contact

Complete lamination

(b)

FIGURE 5.5.19 (a) Schematic of sequential deposition of gate, dielectric, and source and drain electrodes onto a PDMS substrate. Inset: top view of the PDMS transistor platform. (b) Schematic of procedure for laminating an organic semiconductor single crystal against the PDMS transistor platform. Initial contact (top frame) between the crystal and the PDMS results in a “wetting” front (middle frame) that proceeds across the entire crystal/PDMS interface (bottom frame). Insets: optical micrographs of the progress of the wetting front. (From V. C. Sundar et al., Science, 303, 1644, 2004.)



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FIGURE 5.5.20 (a) Current-voltage characteristics of a laminated rubrene single crystal transistor (L = 75 μm, W = 980 μm; Ci = 0.67 nF/cm2). Inset: linear regime current-voltage characteristics. (b) Transfer characteristics measured at a source-drain voltage of –100 V in the same region of a rubrene single crystal by laminating transistor stamp structures of various channel lengths (L = 220, 170, 120, 100, and 75 μm; W = 980 μm). Inset: linear scaling of the saturation currents with the W/L ratio. (From V. C. Sundar et al., Science, 303, 1644, 2004.)



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1 b direction a direction

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20

FIGURE 5.5.21 Four-probe conductivity measurements of a rubrene single crystal as a function of gate voltage along the b and a axes. Intrinsic mobilities of 15.4 and 4.4 cm2/V-s were measured along the b and a axes, respectively. Inset: contact resistance along the a and b directions of the rubrene single crystal. (From V. C. Sundar et al., Science, 303, 1644, 2004.)

in the organic semiconductor single crystal affects charge-carrier mobility across the transistor channel [91]. The rubrene single crystal (p-type), for example, shows maximum hole mobility along the b axis of its unit cell (Figure 5.5.21), consistent with the expectation of stronger π-orbital overlap along the b axis due to molecular packing [91]. In addition to the transistors described earlier, Zaumseil and coworkers [88] fabricated nanoscale transistors by soft-contact lamination, as depicted in Figure 5.5.22. In this procedure, the stamp is incorporated as an element of the device. Specifically, gold evaporated on the raised regions of the PDMS stamp defines the source and drain electrodes while the recessed regions of the stamp define the transistor channel. To access a nanoscale channel, only the recessed regions of the PDMS stamp need to be on the nanometer length scale. An elastomeric nanoscale transistor platform can thus be generated by depositing titanium and gold on the raised regions of an h-PDMS stamp. The PDMS platform is laminated against a silicon substrate (gate) consisting of SiNx (dielectric), pentacene (organic semiconductor), and metal contact pads to connect to the electrodes on the PDMS substrate. Figure 5.5.23 illustrates the electrical properties of a representative nanoscale laminated transistor (channel width of 20 μm and channel length of 150 nm). The transistor exhibits a lower charge-carrier mobility and on/off current ratio compared to laminated micron-size transistors (channel width of 20 μm and channel lengths of 2.5 and 100 μm). Zaumseil and coworkers attributed the differences to contact resistances and short-channel effects [88]. Methods to lower contact resistance, which may include the clever use of monolayer chemistry and conductive polymers as electrodes, are being explored [87,88].


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PDMS stamp with source and drain electrode relief; Ti/Au deposited on plasma oxidized surface

Au Au Bring stamp in close contact with semiconductor and probe pads

Semiconductor Dielectric Gate electrode Probe pads Ti/Au/Ti

FIGURE 5.5.22 Schematic of soft-contact lamination for building thin-film transistors. The electrodes are defined by the raised regions of the PDMS stamp. (From J. Zaumseil et al., Appl. Phys. Lett., 82, 793, 2003.)


–0.25 150 nm

–0.20 –0.15

1 μm

–0.10 –0.05 0.00 0.05 0.0

–0.5 –1.0 –1.5 Source-drain voltage (V)

–2.0

FIGURE 5.5.23 Current-voltage characteristics of a typical laminated nanoscale pentacene transistor (W ~ 20 μm and L ~ 150 nm). Inset: scanning electron micrograph of the channel region defined by the recessed region of the PDMS stamp. (From J. Zaumseil et al., Appl. Phys. Lett., 82, 793, 2003.)

Lee and coworkers [96] also used ScL to construct organic light-emitting diodes (OLEDs). Similar to the transistors described in preceding paragraphs, the laminated OLEDs consist of two parts: conformal gold electrodes supported on PDMS and a transparent, ITO-coated substrate on which the electroluminescent organic material is deposited. Bringing the two substrates together completes the device, as shown in Figure 5.5.24. For comparison, reference OLEDs were fabricated by evaporating gold



465

PDMS Ti/Au

EL layer ITO Substrate (a) PDMS EL layer ITO

(b)

(c)

FIGURE 5.5.24 Schematic of ScL for fabricating laminated OLEDs. (a) Ti (1 nm)/Au (20–60 nm) electrodes are deposited on a PDMS stamp while the electroluminescent layer (EL) is deposited on an ITO-coated transparent substrate. (b) The PDMS substrate containing the electrodes is laminated against the EL layer at ambient conditions to complete the circuit. (c) Optical micrograph of the laminated OLED. (Scale bar, 50 μm). (From T.-W. Lee et al., Proc. Natl. Acad. Sci. U.S.A., 101, 429, 2004.)


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electrodes directly on top of the electroluminescent organic material. Both sets of devices exhibited uniform spatial emission. The laminated devices, however, exhibited higher quantum efficiencies. Lee and coworkers attributed the difference in quantum efficiencies to the disruption of the π–π-conjugation when hot metal atoms are directly evaporated on top of the electroluminescent material [96]. These results further demonstrate the advantages of separating the organic semiconductor deposition process from the source and drain patterning steps during the fabrication of organic devices.

5.5.6 COLD WELDING Cold welding is another contact printing technique that can be used to generate electrodes for top-contact devices [66,97,98]. Cold welding differs from nTP in that it relies on the formation of metallic bonds between two metal surfaces of similar composition for pattern transfer. Figure 5.5.25 illustrates the cold welding process. A silicon hard stamp is coated with an adhesion reduction layer (typically an organic material, such as pentacene) followed by gold. A thin sacrificial layer (strike layer) of gold is deposited on the substrate. Since a rigid stamp is used in this process, external pressures must be applied to achieve conformal contact between the two gold layers. When the stamp and the substrate are brought into contact at high pressures (~150 MPa), cold welding occurs between gold on the raised regions of the stamp and the gold strike layer on the substrate at room temperature. After transferring the gold patterns from the stamp to the substrate, the exposed gold strike layer is removed from the substrate by sputter etching in an argon environment. Kim and coworkers [66] fabricated top-contact pentacene TFTs with gold electrodes patterned by cold welding. To achieve top-contact devices, a gold strike layer was deposited directly on top of pentacene. Cold welding defines the source and drain electrodes; channels with widths of 97 μm and lengths as small as 1 μm were fabricated [66]. Figure 5.5.26 shows the output characteristics of a top-contact pentacene TFT with the gold electrodes formed by cold welding. Although the OTFT is a top-contact device, the output characteristics exhibited significant nonlinear behavior in the linear regime. Kim and coworkers attributed such nonlinear output characteristics to large contact resistance at the gold-pentacene interface [66]. Further, the electrical characteristics of this top-contact pentacene TFT were lower than expected. There are two possible explanations for this observation. First, the etching process to remove the gold strike layer from the channel may damage pentacene, thereby degrading its performance. Second, the high contact pressures required for cold welding (~150 MPa) may also damage pentacene. The bottom inset of Figure 5.5.26 reveals evidence of wrinkles in pentacene after cold welding. The same authors recently demonstrated cold welding at reduced pressures (~180 kPa) using flexible PDMS stamps supported on a glass backplane, rather than rigid stamps [98]. In this example, Kim and coworkers patterned the cathodes for OLEDs by cold welding. This process begins with the deposition of an organic light-emitting material followed by an electron-transporting layer on an ITO-coated glass substrate (anode). The cathode layers and a gold strike layer (i.e., LiF/Al/Au) are then deposited on the electron-transporting layer.



Step 1

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Stamp

Achesion-reduction layer Strike layer

Metal Substrate

Step 2 Cold welding Metal transfer

Step 3 Strike layer removal

Step 4 Pattern replication in a substrate

FIGURE 5.5.25 Schematic of cold welding. A rigid silicon stamp is coated with an adhesion reduction layer, such as pentacene, followed by a metal layer. A thin metal strike layer is also deposited onto the substrate. The stamp is then brought into contact with the substrate. In the regions of contact, the metal is transferred to the strike layer via cold welding. After separating the stamp from the substrate, the exposed strike layer is removed by subsequent etching. (From C. Kim et al., Appl. Phys. Lett., 80, 4051, 2002.)

Separately, a PDMS stamp is prepared on a glass support and is coated with an adhesion reduction layer and a gold layer. The gold layer on the PDMS stamp is then cold welded to the cathodes with the organic light-emitting underlayer. Here, the conformal contact of the PDMS stamp allows the gold layer to be transferred to the strike layer at reduced pressures, thereby preventing damage to the organic light-emitting layer. The quantum efficiency of the OLEDs fabricated by cold welding is comparable with that of OLEDs fabricated by conventional pattern definition by evaporation through shadow masks, thus implying that stamping does not degrade device performance [98]. Subsequently, Loo and workers demonstrated cold welding at ambient pressures [21]. Specifically, they fabricated a metal–insulator–metal capacitor (Au–SiNx–Au) by cold welding. An Au/SiNx/Au multilayer stack was deposited on a silicon stamp and directly printed on an Au-coated PDMS/PET substrate. The details of this device were discussed previously in Section 5.5.4.


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ID(μA)

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D

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0.0

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FIGURE 5.5.26 Output characteristics of a top-contact pentacene TFT with gold electrodes formed by cold welding. This TFT has a channel length of 1 μm as shown in the top inset. Although the TFT adopts the top-contact device geometry, the output characteristics exhibit significant nonlinear behavior at small source-drain voltages. The bottom inset reveals wrinkles in the pentacene layer after etching to remove the underlying Au strike layer. (From C. Kim et al., Appl. Phys. Lett., 80, 4051, 2002.)

5.5.7 METAL TRANSFER PRINTING Similar to nTP and cold welding, metal transfer printing [99] allows top-contact OTFTs to be fabricated. In metal transfer printing, the metal features are directly printed on semiconducting and insulating polymers. Wang and coworkers [99] demonstrated the patterning of gold and aluminum on a variety of polymers, including poly(methyl methacrylate) (PMMA), polystyrene (PS), and P3HT. The polymer, heated above its glass transition temperature, acts as an adhesion promoter, eliminating the need for the metal strike layer in cold welding. Specifically, a thin layer of metal is deposited on the raised and recessed regions of a PDMS stamp. The PDMS stamp is contacted against a heated polymer semiconductor (>120°C), and slight pressure (~2.9 kPa) is applied. After cooling the polymer semiconductor below its glass transition temperature, removing the PDMS stamp transfers the metal patterns onto the polymer. Using metal transfer printing, Wang and coworkers [99] fabricated top-contact P3HT TFTs. P3HT was spin-coated on a substrate containing the gate electrodes coated with a Ta2O5 film, which served as the gate dielectric. Gold electrodes were directly patterned on top of P3HT by metal transfer printing. These OTFTs had a channel width of ~400 μm and a channel length of ~20 μm and exhibited a chargecarrier mobility of 0.017 cm2/Vs, which was comparable to that of P3HT TFTs with photolithographically defined electrodes [99]. While this technique is useful for fabricating top-contact OTFTs with polymer semiconductors, it will not work for small-molecule organic semiconductors since they do not exhibit a softening temperature prior to degradation.



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5.5.8 HOT LIFT-OFF Current techniques for patterning organic semiconductors mainly rely on thermal evaporation through a shadow mask. Consequently, the feature size of organic semiconductors is limited by the resolution of the shadow mask (25–30 μm). Further reduction in the feature size can be achieved by “hot lift-off” (illustrated in Figure 5.5.27a), a technique developed by Wang and coworkers [100]. Using hot lift-off, they fabricated an array of 7- by 7-μm organic semiconductor squares. In this technique, the organic semiconductor is deposited on a substrate by conventional evaporation methods. A partially cured epoxy stamp is then contacted against the organic semiconductor with an applied pressure of 981 kPa to induce local fracture of the organic semiconductor along the patterned edges of the stamp. The applied pressure is then reduced to 196 kPa while heating (80–120°C, 20 min) to maintain conformal contact between the stamp and organic semiconductor while the stamp is further cured. After cooling to room temperature, the epoxy stamp is peeled away from the substrate. In the regions of contact, the organic semiconductor peels with the epoxy stamp, leaving behind patterned organic semiconductor in the noncontact regions on the substrate (Figure 5.5.27). Such lift-off occurs because the work of adhesion between the epoxy and the organic semiconductor is greater than the work of adhesion between the substrate and the organic semiconductor. Although this technique is subtractive in nature, it avoids exposing the remaining organic semiconductor to harsh chemical environments, such as the etching solutions used in photolithography and μCP. A variety of organic semiconductors, including copper phthalocyanine (CuPc), metal-free phthalocyanine (H2-Pc), N,N′-di(naphthalene-1-yl)-N,N′-diphenylbenzidien (NPB), and tri(8-quinolinolato)aluminum (AlQ3) were patterned by hot lift-off. Additionally, Wang and coworkers [100] fabricated functional top-contact CuPc TFTs by hot lift-off and evaporation of the gold electrodes through a shadow mask. These OTFTs exhibited a charge-carrier mobility of 0.02 cm2/Vs, a threshold voltage of –5 V, and an on/off current ratio of 104–105. Although this OTFT adopts the topcontact geometry, the device performance is not better than that of bottom-contact CuPC TFTs in which the CuPC is unpatterned [101].

5.5.9 MICROMOLDING IN CAPILLARIES (MIMIC) Micromolding in capillaries (MIMIC) [102,103] differs from the techniques discussed previously because it is not a contact printing technique. Rather, MIMIC is a molding technique that relies on the spontaneous filling of the microchannels of the mold with a fluid, which can be a solution, a suspension, or precursors of the material to be patterned. This technique is therefore useful for patterning thick polymer, ceramic, and metal features. In MIMIC, a PDMS mold is brought into contact with a rigid support, forming microchannels between the substrate and the recessed regions of the PDMS mold as shown in Figure 5.5.28. The fluid, placed at the end of mold, fills the channels of the mold by capillary action. The material is then cross-linked, crystallized, or


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FIGURE 5.5.27 (a) Schematic of hot lift-off. A partially cured epoxy stamp is placed on an organic semiconductor film. Heating the stamp while applying pressure creates conformal contact between the stamp and the semiconductor. After cooling, peeling the stamp lifts the organic semiconductor underlayer. A variety of organic semiconductors are patterned using hot lift-off. (b) CuPC on Si. (c) Alq3/NPB on ITO. (d) H2PC on Si. (e, f) The cross-sectional profiles of the patterns obtained by AFM reveal sharp edges of the patterns. (From Z. Wang et al., J. Am. Chem. Soc., 125, 15278, 2003.)



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Channels fill by capillary action

Cure; remove PDMS Polymer

FIGURE 5.5.28 Schematic of MIMIC. At least one end of the PDMS stamp is cut off to create entrances to the microchannels. The PDMS mold is then placed on a support to form the microchannels. When a precursor or a prepolymer solution is placed at one end, capillary action spontaneously fills the channels with the solution. After curing, the PDMS stamp is removed, leaving behind solid polymer features on the substrate. (From E. Kim et al., J. Am. Chem. Soc., 118, 5722, 1996.)


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cured before the PDMS mold is removed. The patterned material is therefore a negative image of the PDMS mold. Rogers and coworkers [104] demonstrated the utility of MIMIC for fabricating electrodes for top-contact OTFTs. Specifically, conductive carbon or conducting polymer (PANI doped with m-cresol) electrodes, as wide as 100 μm and separated by 25 μm, were patterned on top of P3HT organic semiconductor by MIMIC. In these devices, the P3HT and gate dielectric (polyimide) layers were prepatterned using screen printing [103]. The source and drain electrodes were defined by the recessed regions (microchannels) of a PDMS mold. To access the microchannels, vertical holes were created at both ends of the PDMS mold. After placing the mold on top of the P3HT layer, a conductive ink (conductive carbon or a conductive polymer, such as PANI) was injected into the access holes of the mold. The ink filled the microchannels of the PDMS mold by capillary forces. After solvent evaporation, the PDMS mold was removed, leaving behind solid conductive carbon or polymer source and drain electrodes. Top-contact P3HT TFTs with conductive carbon source and drain electrodes exhibited a charge-carrier mobility of 0.01–0.05 cm2/Vs, which is comparable with the charge-carrier mobility of bottom-contact P3HT OTFTs with photolithographically-defined gold electrodes [104]. A variation of MIMIC for patterning polymer semiconductor TFT arrays was recently reported by Salleo and coworkers [105]. This technique exploits capillary forces to pattern a solution-processable polymer semiconductor. In this method, the polymer semiconductor solution is spin-coated onto a substrate. A chemically treated PDMS stamp is then placed directly on top of the substrate coated with the polymer semiconductor solution. In the regions of contact, the PDMS stamp absorbs the solvent, leaving behind a solid polymer semiconductor film between the stamp and the substrate. In the recessed regions of the PDMS stamp, the polymer semiconductor solution wicks into the stamp due to capillary forces, effectively leaving behind a clean surface in the noncontact regions. Using this stamping process, patterned arrays of poly[(9,9′-dioctylfluorine)-cobithiophene] (F8T2) and poly[5,5′-bis(3-alkyl-2-thienyl)-2,2′-bithiophene] (PQT12) with features as small as 2 μm were fabricated. Figure 5.5.29(a) illustrates the stamp fabrication and patterning process. Optical micrographs of patterned polymer semiconducting films and a used PDMS stamp after patterning are shown in Figures 5.5.29(b) and (c), respectively. Both top- and bottom-contact TFTs were fabricated with organic semiconductor arrays patterned by this technique. To fabricate topcontact TFTs, gold electrodes are evaporated through a shadow mask onto an F8T2 organic semiconductor array patterned on thermally-grown silicon dioxide/silicon substrate. These top-contact F8T2 TFTs exhibited charge-carrier mobilities of 0.5–1 × 10–3cm2/Vs, and a threshold voltage of ~1 V, similar to those obtained from analogous top-contact TFTs fabricated by spin-coating the F8T2 layer [106]. For bottom-contact TFTs, gate electrodes (Cr/Au), gate dielectric (SiNx/SiO2 bilayer), and source and drain electrodes (Cr/Au) were defined by digital lithography [107]. F8T2 or PQT-12 organic semiconductor arrays were then patterned. While the bottom-contact F8T2 TFTs showed a charge-carrier mobility of 3 × 10–4cm2/Vs and a threshold voltage of –10 V, the PQT-12 TFTs had better device characteristics with a charge-carrier mobility of 3 × 10–3cm2/Vs and a threshold voltage of 10 V.



473

F8T2 Substrate Jet-print wax pattern Mold PDMS stamp

300 μm 120 μm

Release PDMS stamp O2 plasma

600 μm

Activate stamp surface

1 mm

Benzyl-TS Treat stamp surface Recess

(b)

Raised feature Coat substrate with solution place stamp

Solvent absorption

Raised portion

Capillary action

Stamp recess

Allow solution to wick in recesses and solvent to permeate stamp

Patterned polymer

Remove stamp

(c)

(a)

FIGURE 5.5.29 (a) Schematic of the stamp fabrication and patterning process of F8T2 and PQT-12. (b) Optical micrographs of patterned F8T2 arrays on thermally grown SiO2 on a Si substrate. (c) Optical micrograph of a used stamp reveals the polymer residue in the recess regions of the stamp. (From A. Salleo et al., Adv. Funct. Mater., 15, 1105, 2005.)

These charge-carrier mobilities, however, are two orders of magnitude lower than those achieved by spin-coating or jet-printing the organic semiconductor onto a gate dielectric treated with a hydrophobic molecular layer (0.02 cm2/Vs for F8T2 and 0.1 cm2/Vs for PQT-12). Salleo and coworkers attributed the reduction in chargecarrier mobilities to the fact that these TFTs were fabricated on untreated dielectric surfaces. The observed charge-carrier mobilities are comparable to those obtained by other processing methods on untreated dielectric surfaces.

5.5.10 SOFT-CONTACT OPTICAL LITHOGRAPHY As demonstrated throughout this chapter, the unique mechanical properties of PDMS have enabled many soft lithographic techniques. Soft-contact optical lithography developed by Lee and coworkers [108] is a variation of photolithography for patterning both rigid and flexible substrates at a low cost by replacing expensive photomasks with PDMS “masks.” Since the PDMS mask is flexible, it can make conformal contact against flexible or curved surfaces. Soft-contact lithography therefore allows metals to be patterned on flexible or nonflat substrates via traditional photolithography. A schematic of soft-contact lithography is shown in Figure 5.5.30.


474


PDMS Oxygen plasma Ti/Au Deposition Ti/Au

PDMS

Spin cast photoresist Photoresist Ti/Au

PDMS

No prebaking hv PDMS

PDMS I: Lower dose II: Higher dose Developing PDMS

PDMS

Au etching (Kl sulution)

PDMS

PDMS

FIGURE 5.5.30 Schematic of soft-contact optical lithography. Ti (1 nm)/Au (20 nm) is uniformly evaporated on a flat PDMS substrate and coated with photoresist. A patterned PDMS mask, which is transparent to UV light, is laminated against the photoresist layer. The backside of the PDMS mask is illuminated with UV light. The exposed PR film is developed. The resulting pattern is transferred to the Ti/Au layer by etching. Controlling the exposure dose results in two distinct line-width patterns. Lower exposure doses replicate the PDMS relief pattern; higher exposure doses generate line widths from 50 to 150 nm by phase-shifts in the optical near field. (From T.-W. Lee et al., Adv. Funct. Mater., 15, 1435, 2005.)

Specifically, photoresist is spin-coated on a gold-coated PDMS substrate. After establishing conformal contact between the PDMS mask and the photoresist/gold/PDMS substrate, the assembly is exposed to 330 nm UV light. By controlling the exposure dose and developing conditions, a single PDMS mask is capable of generating two different patterns in the resist. At low-exposure doses, the regions of photoresist in contact with the PDMS mask become soluble and can be removed during developing, leaving behind a pattern that replicates the recessed regions of the PDMS mask. Higher exposure doses or longer developing times generate narrower lines (as small as 150 nm) positioned at the contact edges of the features on the mask. At these higher exposure doses, the PDMS mask acts like a near-field phase-shift mask,



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allowing nanoscale features such as periodic and curved lines, open rings, and posts on planar and curved surfaces to be patterned [22,108–110]. Combining soft-contact optical lithography with ScL enables the fabrication of laminated OLEDs with nanometer-scale electrodes [96,108].

5.5.11 LASER THERMAL TRANSFER PRINTING By definition, laser thermal transfer printing is not a soft lithography technique. It is, however, a useful technique for patterning organic components of OTFTs. Laser thermal transfer printing is a solventless thermal imaging technique [111]. One of the most attractive features of this technique is that it circumvents solvent incompatibility issues frequently found in inkjet printing or screen printing. This technique also provides rapid printing of micron-sized organic material features over large areas. Blanchet and coworkers [111] used laser thermal transfer printing for patterning conducting polymer electrodes on flexible substrates for pentacene TFTs. Figure 5.5.31 illustrates this laser thermal transfer printing process. This method uses two flexible substrates. The multilayer donor substrate consists of uniformly coated thin layers of conducting polymer and metal. The other substrate is the receiver substrate on which patterned features of organic material will be transferred. These two flexible substrates are loaded into the near infrared laser printer and held together by vacuum. Focusing the laser beam on the thin absorbing metal layer in the donor substrate converts light into heat in this layer. The generated heat decomposes the conducting polymer at the metal-conducting polymer interface into gaseous by-products. The expansion of gaseous products transfers the top layer of the conducting polymer onto the receiver layer. The transfer of organic material occurs pixelwise. Gate, as well as source and drain, electrodes are patterned from a PANI/single wall carbon nanotubes (SWNT)-coated donor substrate using this technique. Laser thermal transfer printing allows both bottom- and top-contact devices to be fabricated. Bottom-contact pentacene TFTs with printed PANI/SWNT electrodes on glass resin/ITO/Mylar substrates had channel widths of 750 μm and channel lengths ranging from 10 to 250 μm. The charge-carrier mobility of these OTFTs was 0.3 cm2/Vs, which is twice as high as that of reference OTFTs with gold electrodes (0.15 cm2/Vs), at similar channel dimensions. Interestingly, pentacene TFTs with top-contact printed PANI/SWNT electrodes exhibited nonlinear output characteristics. Analogous to the top-contact devices fabricated by cold welding (where the pentacene layer is wrinkled in Figure 5.5.26), it is speculated that pentacene may have been thermally damaged or degraded during the PANI/SWNT transfer process.

5.5.12 IMPRINT LITHOGRAPHY While imprint lithography does not involve the use of an elastomeric stamp and hence is not a soft lithography technique, it is a next-generation patterning technique that has gained tremendous attention for generating nanoscale features, which is


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FIGURE 5.5.31 Illustration of laser thermal transfer printing. Two flexible substrates are held together by vacuum. One film is the donor substrate, which consists of thin layers of conducting polymer and metal uniformly coated on a substrate. The other film is the receiver onto which patterns of conducting polymer are transferred. Focusing the laser beam onto the thin absorbing metal layer generates heat at the polymer conductor interface. The generated heat decomposes the conducting polymer into gaseous by-products, which then push the conducting polymer onto the receiver layer. (From G. B. Blanchet et al., Appl. Phys. Lett., 82, 463, 2003.)

why we feel that it is important that this patterning technique be discussed. Imprint lithography provides a higher resolution and lower cost alternative for fabricating nanoscale electrical circuitry components as compared to photolithography [112]. The focus of this section is thus to provide an introduction to the technology addressing the advancements and concerns associated with this next-generation patterning technology. For a more detailed discussion on the state of imprint lithography and its prognosis for future application in integrated circuit fabrication, please refer to a recent topical review written by Guo [113]. The two leading imprint lithography techniques are nanoimprint lithography (NIL) [112,114] and step-and-flash imprint lithography (S-FIL) [115,116]. Both techniques rely on a two-step process of imprinting and pattern transfer. A schematic of the NIL process is shown in Figure 5.5.32 [114]. In the NIL process, the resist is a thermoplastic [112,117] or thermal curable polymer [118]. During NIL, the thermoplastic is heated above its glass transition temperature, allowing it to flow, before imprinting with the mold. This mold is typically fabricated on a silicon or



477

Mold

1. Press mold

Resist Substrate

2. Remove mold

3. RIE

FIGURE 5.5.32 Schematic of nanoimprint lithography (NIL). (1) A rigid mold is pressed into a thermoplastic resist, heated above its glass transition temperature, to create an imprint. (2) The mold is removed, leaving behind the patterned resist. (3) An anisotropic etch removes the residual resist and transfers the pattern into the substrate. (From S. Y. Chou et al., J. Vac. Sci. Technol., B, 14, 4129, 1996.)

fused silica template using e-beam lithography. Upon mold contact, the resist is cooled below its glass transition temperature, effectively “locking in” the shape of the mold, which is subsequently removed. It is important to note that NIL is a physical deformation process, not an ink-stamping process. Consequently, the diffusion concerns associated with soft lithography techniques, such as μCP, do not exist. Further, the imprint mold is rigid, so feature collapse and deformation during imprinting are eliminated. In combination, these attributes of NIL have enabled the patterning of extremely small (below 10 nm) yet high-resolution features, as illustrated in Figure 5.5.33 [119–121]. 6 nm Half-pitch

8.5 nm Half-pitch

17 nm Half-pitch

FIGURE 5.5.33 Scanning electron micrograph of polymer resist with 6-, 8.5-, and 17-nm halfpitch gratings fabricated by NIL. (From M. D. Austin et al., Nanotechnology, 16, 1058, 2005.)


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While NIL is a promising next-generation lithography technique with excellent feature resolution, several concerns still need to be addressed before the technique will have industrial implications. In particular, commercially-available molds, resists, and processes still need to be developed [113]. Since the mold plays a crucial role in imprint lithography, procedures for preventing and removing mold contamination must be addressed. Also, resist materials with suitable physical properties [113], such as low viscosity and low glass transition temperatures, must be exploited to minimize the pressure [122] and the time for heating and cooling cycles required for polymer molding — potential bottlenecks in the manufacturing process. Additionally, thermal expansion effects of the resist and the mold must be considered when heating and cooling the thermopolymer to minimize feature distortion and alignment difficulties. This being said, NIL remains a powerful patterning tool. Austin and coworkers demonstrated the usefulness of NIL for fabricating bottom-contact OTFTs with channel lengths below 100 nm [123]. Figure 5.5.34 illustrates their procedure. A negative imprint of the transistor channel region is defined on a silicon platform by NIL. An oxygen plasma etch removes residual polymer from source and drain regions. Subsequently, gold is uniformly deposited across the substrate, creating source and drain electrodes, and the polymer imprint is removed by lift-off. Deposition of P3HT polymer semiconductor completes the transistor. A protective silicon dioxide cap is evaporated onto the P3HT to protect it from moisture and oxygen. Typical current-voltage characteristics are shown in Figure 5.5.35 for transistors with channel widths of 2.5 μm and channel lengths ranging from 70 to 1,000 nm. Short-channel effects are clearly visible as the channel length is decreased. The second imprint lithography technique discussed here is S-FIL. Similar to NIL, S-FIL can be a potentially high-throughput, low-cost technique for fabricating nanometer-scale features. A schematic of the S-FIL process is shown in Figure 5.5.36 [124]. A photocurable liquid (i.e., resist) is dispensed into the gap between the transparent (to UV) template (analogous to the mold used in NIL) and the substrate. Typically, the template is fabricated on a fused silica plate using conventional phaseshift reticle processing [125,126]. Willson and coworkers, however, have developed other mask fabrication techniques that incorporate a transparent indium tin oxide layer to improve imaging of the mask during defect analysis [127–130]. When the liquid resist is dispensed between the template and the substrate, capillary forces cause the liquid to disperse into the gaps and pull the template tight against the substrate. Illumination of the mold and resist with UV radiation initiates polymerization, creating a replica of the mold. Removal of the template completes the imprint process. The pattern is then transferred into the substrate through a series of etching steps. The key difference between NIL and S-FIL is the type of the resist used for pattern transfer. In S-FIL, a low-viscosity, photocurable organosilane liquid is used with a transparent template. Due to the low viscosity of the resist, external pressure is not required for imprinting. Further, the resist photocures at room temperature, thereby eliminating the heating and cooling cycles required in NIL. The rigid, transparent S-FIL template permits flood exposure of the photocurable resist and layer-to-layer alignment through classical optical techniques [131]. These attributes



NIL Mold

479

(a) Nanolmprint: Imprint polymer 5 nm Gate oxide n+ Silicon Substrate (back-side gate)

NIL Mold (b) Mold Separation

(c) Oxygen Plasma Etch Residue Imprint Polymer 70 nm Channel Length (d) Gold Source / Drain Evaporation and Lift-off

(e) Add: Gold Pads Protective SiOx (f ) Add Semiconducting Polymer P3HT: R = C6H13 S Source

m

Drain (g) Protect and Isolate the Device with SiOx cap

Gate

FIGURE 5.5.34 NIL for fabricating organic thin-film transistors. (a) A negative imprint of the transistor channel is created by pressing and (b) removing the mold. (c) Residual polymer is removed with an oxygen plasma etch. (d) Gold source and drain electrodes are evaporated. Lift-off is performed to remove the polymer imprint. (e) Additional gold pads are added for making output characteristic measurements. A layer of SiOx is also added to protect the gate oxide from probe damage during output characteristic measurements. (f) The polymer semiconducting layer, P3HT, is deposited to complete the transistor. (g) A SiOx layer is deposited over the P3HT layer to protect it from moisture and oxygen. (From M. D. Austin and S. Y. Chou, Appl. Phys. Lett., 81, 4431, 2002.)


480


Drain current density [nA/um]

–2


(a) 1000 nm Device –3

101

Vd = –1.0V

Vg = –3V

Vd = –1.5V

10–5 –4 –2

0

2

Gate voltage Vg = –2V

–1 Vg = –1V 0

0

–1

–2

–3

Drain voltage


–4


(b) 200 nm Device –6

10

1

Vd = –1.0V Vd = –0.5V

Vg = –3V 10–5 –4

–2

0

2

Gate voltage

Vg = –2V

–2 Vg = –1V

0

0

–1

–2

–3

Drain voltage (b) 70 nm Device Drain current density [nA/um]


–30

–20

Vd = –1.0V

0

10–5 –4

–2V –2

0

2

–1

–1V

–2 Drain voltage

FIGURE 5.5.35

Vg = –3V

Vd = –0.5V

Gate voltage

–10

0

101

–3



481

Release layer Etch barrier Base layer (a)

(b)

UV

(c) Base layer

(d)

(e)

FIGURE 5.5.36 Schematic of step-and-flash imprint lithography (S-FIL). (a) A photocurable liquid (etch barrier) is dispersed between a transparent template and substrate. (b) The gap between the template and the substrate is closed to create a thin resist layer. (c) The backside of the template is illuminated with UV light to cure the resist. (d) The template is removed, leaving behind the imprinted features. (e) An anisotropic etch removes the base layer and transfers the pattern into the substrate. (From T. Bailey et al., J. Vac. Sci. Technol., B, 18, 3572, 2000.)

FIGURE 5.5.35 (See figure, facing page.) Current-voltage characteristics of P3HT TFTs with a channel width of 2.5 μm and channel lengths of (a) 1,000 nm; (b) 200 nm; and (c) 70 nm. Insets: transfer characteristics of the respective devices. (From M. D. Austin and S. Y. Chou, Appl. Phys. Lett., 81, 4431, 2002.)


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(a)

(b)

40 μm

FIGURE 5.5.37 Optical micrographs of S-FIL template (a) before and (b) after two imprints. Template contamination is removed during imprinting. (From T. Bailey et al., J. Vac. Sci. Technol., B, 19, 2806, 2001.)

allow an entire wafer to be patterned in a “step-and-flash” method similar to conventional photolithography. The resolution of S-FIL is limited by the resolution of the features on the template. To date, features as small as 30 nm [132] and with aspect ratios of 14:1 [133,134] have been demonstrated with S-FIL. Additionally, S-FIL is capable of pattering surfaces with topography [133]. Similar to NIL, in order for S-FIL to become commercially successful, commercially available templates and processes still need to be developed. Again, the quality of the template is critical for successful imprinting. To ensure that the photocured polymer releases cleanly from the template, the surface of the template is treated with an FSAM to create a nonstick surface [124]. Defect susceptibility studies to quantify the effectiveness of the FSAM treatment indicate that S-FIL is actually self-cleaning [124,135], as depicted in Figure 5.5.37. The contamination from the template is entrained in the polymer resist during photopolymerization, resulting in a template that is visually clean after imprinting. More recently, a new strategy was developed by Kim and coworkers [136] to promote polymer release from the template. Specifically, a fluorinated acetate is added to the resist formulation. During polymerization, the additive migrates to the template/resist interface, reducing the separation force by ~0.9 lbf as compared to the same formulation without the fluorinated acetate. While the photocurable polymers used for S-FIL avoid the viscosity and thermal concerns associated with the thermoplastics used in NIL, they are not without their problems. During polymerization, the photocurable polymers undergo a volumetric shrinkage or densification due to chemical bond formation. Consequently, the feature size, shape, and position could be affected. Studies have shown, however, that by controlling the resist composition, the shrinkage is limited to the z-direction, resulting in a reduced aspect ratio [131,137]. While NIL and S-FIL have been shown to be effective tools for creating nanoscale patterns, it is important to keep in mind that the patterned polymer layers utilized in both techniques are sacrificial structures. Additional etchback steps are required to transfer the patterns into the substrate. Further, metal contacts and other functional materials have to be deposited separately to create functional devices.



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6.1

Radio Frequency Identification Tags

Vivek Subramanian CONTENTS 6.1.1 Introduction................................................................................................489 6.1.2 An Overview of RFID Standards and Classifications...............................490 6.1.2.1 135 kHz RFID.............................................................................490 6.1.2.2 13.56 MHz RFID ........................................................................491 6.1.2.3 900 MHz and 2.4 GHz RFID .....................................................491 6.1.3 Radio Frequency Identification Using Silicon: A Review ........................491 6.1.4 All-Printed RFID Tags: Topology and Architecture Framework..............492 6.1.4.1 Antenna Stage .............................................................................493 6.1.4.2 Rectifier/Power Supply and Clamp.............................................495 6.1.4.3 Digital Section and Modulation Stage........................................497 6.1.5 An Archetypal First Organic RFID Tag ....................................................500 6.1.6 Implications of Tag Architecture on Device Considerations ....................501 6.1.6.1 Transistor Performance and Structural Implications ..................501 6.1.6.2 Circuit Issues ...............................................................................503 6.1.7 Conclusions................................................................................................504 References..............................................................................................................504

6.1.1 INTRODUCTION Radio frequency identification (RFID) tags have received substantial attention in recent years as a potential application for printed organic transistors. The primary driver for consideration of this application is cost; it is expected that the cost of an RFID tag fabricated using printed transistors will be substantially lower than that achievable using conventional technologies. In this section, the status and requirements of RFID are reviewed, with particular emphasis on the process and performance requirements of organic transistors. The section begins with a review of the state of the art of RFID standards and technologies using conventional technologies. The shortcomings in these technologies are identified and used to motivate the interest in printed RFID tags. Next, the possible topologies of RFID tags are reviewed, with emphasis on the performance constraints imposed by printing. Finally, the state of the art of organic transistors for RFID applications is reviewed;

489


490


this analysis is used to drive the identification of future needs or requirements for organic transistors, thus establishing a roadmap of activities and innovation that will drive the realization of printed RFID.

6.1.2 AN OVERVIEW OF RFID STANDARDS AND CLASSIFICATIONS In general, RFID tags may be classified based on two partitions: • •

the method by which they obtain power to operate the specific frequencies at which they communicate with the reader

In the first level of classification, tags are categorized based on whether they have an on-tag battery, or depend on the reader to provide them with power. The former are called “active tags,” while the latter are called “passive tags.” Active tags, due to their higher costs and extended ranges (several meters or more) are currently using in inventory management and high-value asset tracking applications. Passive tags do not contain a battery. Instead, power is supplied to them by the reader through electromagnetic coupling. The reader broadcasts large amounts of power, a small percentage of which is captured or “harvested” by the antenna on the tag. Tags that are close to the reader are thus able to collect enough power to become energized. Since this power transfer is extremely inefficient, the range of passive tags is usually limited; common ranges are a few centimeters to a few meters. Organic transistorbased tags are generally only considered for use in passive tags; since the cost points of active tags are higher, the costs provided by silicon-based circuitry are adequate for most active tag applications. Therefore, in this section, only passive tag architectures will be considered. In the second level of classification, (passive) RFID tags are classified based on the manner in which they communicate with the reader. This is performed based on the frequency at which the reader broadcasts information and power to the tag. In general, the frequency bands already used for RFID around the world are: (1) below 135 kHz; (2) 13.56 MHz; (3) 900 MHz; and (4) 2.4 GHz.

6.1.2.1 135 KHZ RFID Tags operating at 135 kHz operate in the near-field of the reader (i.e., they interact primarily with the magnetic component of the electromagnetic signal broadcast by the reader). In this regime, both the tag and reader antennae are inductors. When the tag is within a usable operating range of the reader, the two inductors are coupled, resulting in a mutual inductance and magnetic coupling between them. Since the inductors required to produce resonant circuits at such low frequencies are large, these inductors are typically wound coil inductors rather than planar spiral inductors [1]. Hence, the fabrication costs associated with low-frequency RFID are generally high, and their use is primarily limited to applications where low frequency is specifically advantageous.



491

Low frequencies generally do not significantly interact with water and other fluids and also work fairly well in the presence of metals. Therefore, these tags work well in environments containing liquids and/or metals. Indeed, one of the most important uses of these tags is for livestock inventory control; tags are inserted into livestock and are readable from the innards of the animal with an external reader. However, given the high costs associated with wound coils, the cost of silicon is only a small fraction of the cost of the tag and, as a result, the use of conventional silicon-based circuitry (with its associated performance and reliability) is acceptable for most low-frequency RFID applications.

6.1.2.2 13.56 MHZ RFID Arguably, the most important RFID operating frequency used today is 13.56 MHz RFID. At this frequency, power is still inductively coupled. However, given the higher frequencies used, it is possible to use planar spiral inductors. These may be fabricated at substantially lower costs than the wound coil inductors used in LF RFID; as a result, 13.56 MHz RFID has received substantial attention as a candidate frequency for low-cost RFID applications [2]. In the presence of liquids, 13.56 MHz RFID works well, but is moderately susceptible to interference from metals nearby. However, compared to other candidate frequencies, 13.56 MHz RFID is generally considered to be the most promising frequency for use in RF barcodes.

6.1.2.3 900 MHZ

AND

2.4 GHZ RFID

RFID tags operating at 900 MHz and 2.4 GHz are called ultrahigh frequency (UHF) and microwave tags. Since these tags typically operate in the far-field region of the reader electromagnetic field, they primarily interact with the electric component of the same. Backscatter systems are used for communication, and the antenna is typically a dipole configuration [3]. This is conveniently formed in a printed planar configuration; as a result, printed antennae have already been adopted in these frequency ranges. These systems are able to achieve relatively long-range, high datarate operation and are therefore considered to be extremely attractive for pallet-level and case-level tracking applications, for example. Unfortunately, such tags are extremely sensitive to liquids and metals and also show various interference phenomena due to reflections, which has resulted in reduced interest in their use for item-level tracking.

6.1.3 RADIO FREQUENCY IDENTIFICATION USING SILICON: A REVIEW RFID technology to this point has had mixed success. While RFID has been tremendously successful in high-value asset tracking applications, it has only had moderate success in the low-cost applications that offer the largest potential markets. This has primarily been limited by the cost of current RFID, which is still a little too high for the target markets [4]. This, in fact, is the primary motivation for the use of organic transistors in low-cost RFID. Therefore, to facilitate a better


492

Individual chips are fabricated on water, but not separated


Chips are separated for handling

Chips are attached to a strap Si

Straps are attached to antennae

FIGURE 6.1.1 Conceptual process for silicon-based RFID tag fabrication.

understanding of the economics therewith, an overview of RFID based on conventional silicon is provided. The cost of silicon chips rises dramatically with increase in chip size, due to a reduction in process yield with increasing chip size and also due to a high cost per unit area associated with silicon. As a result, size reduction is a dominant goal in silicon-based RFID. This has important consequences for circuit partitioning. The size of the antenna is independent of the process technology used and, in fact, can be very large, depending on the operating frequency and desired range. Therefore, in silicon-based RFID tags, the antenna is moved off-chip. In other words, everything but the antenna is fabricated on the silicon chip, which is then attached to a separate antenna, typically fabricated on a plastic inlay or strap, as shown in Figure 6.1.1. While this partitioning strategy reduces the cost of the silicon chip, it adds additional components into the cost equations associated with the overall tag. The overall cost equation for an RFID tag therefore includes three main components: (1) the cost of silicon; (2) the cost of attachment; and (3) the cost of the antenna. Coupling all of these factors together, costs for silicon-based RFID will likely fall into the range of US$0.05 per tag in the foreseeable future, but will struggle to get into the range of US$0.01–0.02 required for many RF barcode applications. This has driven the interest in printed electronics, since, in an ideal world, the cost of a fully printed RFID tag would not be significantly larger than the cost of today’s printed antennae, resulting in a dramatic reduction in overall tag cost to less than US$0.02 per tag.

6.1.4 ALL-PRINTED RFID TAGS: TOPOLOGY AND ARCHITECTURE FRAMEWORK As has been discussed in previous sections, while the economic imperatives for considering printed electronics in general and organic electronics in particular are attractive, there are clear performance trade-offs. First, the performance of printed semiconductors is much worse than the performance of silicon. Second, the line widths and layer-to-layer registration achievable by printing are substantially worse than achievable in silicon technology as well. Based on these constraints, it is



493

Modulation circuit

Clock S SET Q R CLR Q

Power supply

Digital (implements finite state machine, memory, etc.)

Antenna stage

Protection clamp

FIGURE 6.1.2 Block diagram of archetypal HF passive RFID tag.

possible to examine possible circuit architectures for printed RFID tags and then discuss the implications of printed device performance on the viability of the same. As a first constraint, recall that all passive tags require that power be derived from the incident carrier frequency broadcast by the reader. For low-cost applications, this will likely be in the high frequency (HF) or UHF bands, as discussed previously. Given the performance limitations of organic materials, it is highly unlikely that harvesting of UHF frequencies will be possible. As a result, the circuit analysis here will focus exclusively on HF tags. As discussed previously, HF tags operate at 13.56 MHz. Power is inductively coupled from the reader to the tag. A block diagram architecture for a typical HF RFID tag is shown in Figure 6.1.2. For the purposes of simplifying analysis, the tag has been broken down into several blocks, which will be individually discussed.

6.1.4.1 ANTENNA STAGE As discussed earlier, the antenna in an HF tag operates through inductive power coupling. Conceptually, the reader antenna coil acts as a primary coil in a transformer, with the coil in the tag acting as a secondary. Coupling occurs through air; if the tag is within the near-field region of the reader antenna, magnetic coupling between the two coils occurs, and some voltage is “harvested” by the tag coil. To increase the voltage generated by the harvesting process, the antenna stage on the tag is typically a resonant circuit with the antenna inductor connected in parallel


494


L1 (Inductor on reader)

V1

Rs (Resistance of L2)

C2 (Tuning capacitor)

L2 (Inductor on tag)

RL (Load resistance of RFID circuit/ regulator circuit)

FIGURE 6.1.3 Conceptual equivalent circuit for an inductively coupled RFID tag.

with a tuning capacitor. The specific values of the inductance and capacitance are chosen so as to cause resonance at 13.56 MHz. The consequence of this is that the voltage seen at the terminals of the tuned circuit is Q-boosted, where Q is the loaded quality factor of the antenna stage [1]. Typical unloaded Q-values in conventional RFID circuits range from one to two, depending on the desired range. To achieve high Q, low series resistance of the inductor metallization is desired. As a result, substantial effort in recent years has been devoted to the development of low-resistance printed metals [5,6]. This antenna series resistance is extremely important, since the series resistance of the antenna acts as a loss mechanism, falling in series in the LC circuit (Figure 6.1.3) and thus reducing the Q of the circuit. This in turn reduces the power harvested and made available to the RFID circuit. In Figure 6.1.3, L1 is the inductor on the reader, L2 is the inductor in the tag, C is the tuning capacitor on the tag, Rs is the series resistance of the inductor, and Rl is the equivalent load resistance of the RFID circuitry. As is apparent from Figure 6.1.3, the two inductors in fact form a transformer with: v1 = L 1

di1 di2 +M dt dt

(6.1.1)

di1 di2 v2 = M + L2 dt dt where M is the mutual inductance and is related to the coupling factor by: k=

M

(6.1.2)

L 1L 2

Note that the coupling factor between the two inductors L1 and L2 is typically extremely low (

Organic Field-Effect Transistors

Transistors

Thin-Film Transistors