Conjugated Polymers: Perspective, Theory, and New Materials [4 ed.] 9781138065697

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Conjugated Polymers

Perspective, Theory, and New Materials

Conjugated Polymers

Perspective, Theory, and New Materials

Edited by

John R. Reynolds, Barry C. Thompson, and Terje A. Skotheim

Cover art by Ellen Skotheim. A collage, based on images from important developments in conducting polymers as represented by the 4th edition of the Handbook.

CRC Press Taylor & Francis Group 6000 Broken Sound Parkway NW, Suite 300 Boca Raton, FL 33487-2742 ©  2019 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 on acid-free paper International Standard Book Number-13: 978-1-138-06569-7 (Hardback) This book contains information obtained from authentic and highly regarded sources. Reasonable efforts have been made to publish reliable data and information, but the author and publisher cannot assume responsibility for the validity of all materials or the consequences of their use. The authors and publishers have attempted to trace the copyright holders of all material reproduced in this publication and apologize to copyright holders if permission to publish in this form has not been obtained. If any copyright material has not been acknowledged please write and let us know so we may rectify in any future reprint. Except as permitted under U.S. Copyright Law, 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. Visit the Taylor & Francis Web site at  http://www.taylorandfrancis.com  and the CRC Press Web site at  http://www.crcpress.com 

Contents Authors..................................................................................................................... vii Contributors.............................................................................................................. ix

1

Early History of Conjugated Polymers: From Their Origins to the Handbook of Conducting Polymers...................................................................1 Seth C. Rasmussen

2

Recent Advances in the Computational Characterization of π-Conjugated Organic Semiconductors.. ......................................................... 37 Jean-Luc Brédas, Xiankai Chen, Thomas Körzdörfer, Hong Li, Chad Risko, Sean M. Ryno, and Tonghui Wang

3

Perspective on the Advancements in Conjugated Polymer Synthesis, Design, and Functionality over the Past Ten Years....................................... 107 Brian Schmatz, Robert M. Pankow, Barry C. Thompson, and John R. Reynolds

4

Advances in Discrete Length and Fused Conjugated Oligomers..................149

5

Direct (Hetero)Arylation Polymerization for the Preparation of Conjugated Polymers...................................................................................... 195

Shanshan Chen, So-Huei Kang, Sang Myeon Lee, Tanya Kumari, and Changduk Yang

J. Terence Blaskovits and Mario Leclerc

6

Living Polymerizations of π-Conjugated Semiconductors........................... 239

7

Controlled Synthesis of Polyfurans, Polyselenophenes, and Polytellurophenes. . ......................................................................................... 263

Jeffrey Buenaflor and Christine Luscombe

Shuyang Ye, Emily L. Kynaston, and Dwight S. Seferos

8

Donor-Acceptor Polymers for Organic Photovoltaics.................................. 283

9

Conjugated Polymers for n- and p-Type Charge Transport......................... 325

10

Conjugated Block Copolymers: Synthesis, Self-Assembly, and Device Applications.. .................................................................................................. 429

Desta Gedefaw and Mats R. Andersson

Zachary S. Parr, Zhijie Guo, and Christian B. Nielsen

Jessica Shaw and Malika Jeffries-EL

v

vi

Contents

11

Metal-Containing Conjugated Polymers. . ..................................................... 447

12

Recent Progress in the Development of Optoelectronic Materials Based on Group 13 Element-Containing Conjugated Polymers. . ............................ 489

Christopher M. Brown and Michael O. Wolf

Shunichiro Ito, Masayuki Gon, Kazuo Tanaka, and Yoshiki Chujo

13

Multifunctional Conjugated Polymers: Helically Assembled Spherulites, Photo-Controllable Illuminants, and Helical Graphites............................... 517 Kazuo Akagi

14

Conjugated Polyelectrolytes Designed for Biological Applications............. 547

15

Oxidative Chemical Vapor Deposition for Conjugated Polymers: Theory and Applications.. .............................................................................. 587

Pradeepkumar Jagadesan, Yun Huang, and Kirk S. Schanze

Karen K. Gleason and Xiaoxue Wang

16

Flow Synthesis: A Better Way to Conjugating Polymers?..............................613 James H. Bannock, Martin J. Heeney, and John C. de Mello

Index. . ...........................................................................................................653

Authors John R. Reynolds, a native Californian, obtained his B.S. in Chemistry at San Jose State University (1979) followed by his M.S. (1982) and Ph.D. (1984) in Polymer Science and Engineering at the University of Massachusetts. He became interested in the field of conducting and electroactive polymers through a position with the IBM Research Laboratories in the late 1970s. After developing his own research effort at the University of Texas at Arlington (1984–1991), he moved to the University of Florida where he was a Professor of Chemistry and Associate Director of the Center for Macromolecular Science and Engineering until Spring 2012, when his group moved to Georgia Tech and where he is currently a Professor of Chemistry and Biochemistry, and Materials Science and Engineering. He serves as Director of the Georgia Tech Polymer Network (GTPN) and is a member of the Center for Organic Photonics and Electronics (COPE) management team. Barry C. Thompson was born in Milwaukee, Wisconsin, in 1977 and moved at a young age to Gallipolis, Ohio, where he attended elementary and high school. Barry then attended the University of Rio Grande in Rio Grande, Ohio, where he majored in Chemistry and Physics and minored in Mathematics. After completing his undergraduate studies at Rio Grande, Barry moved to the University of Florida to pursue a Ph.D. in Chemistry with Prof. John R. Reynolds as an NSF Graduate Research Fellow. During his Ph.D. studies, Barry focused on the design and synthesis of electroactive conjugated polymers for electrochromic and photovoltaic applications. Upon completion of his Ph.D. in 2005, Barry moved to Prof. Jean Fréchet’s lab at UC Berkeley to further pursue his interests in polymer-based photovoltaics as an ACS-PRF Postdoctoral Fellow. After a three-year stay at Berkeley, Barry moved to the University of Southern California Department of Chemistry and Loker Hydrocarbon Research Institute as an Assistant Professor of Chemistry. Barry was promoted to Associate Professor with Tenure in 2015. Terje A. Skotheim is the founder of Lightsense and has a successful record in developing new technologies and launching new products through several startups in fields as diverse as advanced lithium-sulfur batteries, MEMS devices, photovoltaic cells, and biosensors. His research interests have spanned across several disciplines in materials science, including conducting polymers, semiconductors, ion conductors, and diamond-like carbon. He has held research positions and co-founded companies in Europe and the US, and was head of the conducting polymer group at DOE’s Brookhaven National Laboratory before launching his career as an entrepreneur. He received his B.S. in Physics from the Massachusetts Institute Technology and his Ph.D. in Physics from the University of California at Berkeley.

vii

Contributors Kazuo Akagi Research Organization of Science and Technology Ritsumeikan University Kusatsu, Japan Mats R. Andersson Flinders Institute for NanoScale Science and Technology Flinders University Adelaide, Australia James H. Bannock Department of Chemistry Imperial College London London, United Kingdom J. Terence Blaskovits Département de Chimie Université Laval Québec City, Québec, Canada Jean-Luc Bredas School of Chemistry and Biochemistry and Center for Organic Photonics and Electronics Georgia Institute of Technology Atlanta, Georgia Christopher M. Brown Department of Chemistry University of British Columbia Vancouver, British Columbia, Canada Jeffrey Buenaflor Department of Chemistry University of Washington Seattle, Washington

Shanshan Chen Department of Energy Engineering School of Energy and Chemical Engineering Ulsan National Institute of Science and Technology (UNIST) Ulsan, South Korea Xiankai Chen School of Chemistry and Biochemistry and Center for Organic Photonics and Electronics Georgia Institute of Technology Atlanta, Georgia Yoshiki Chujo Department of Polymer Chemistry Graduate School of Engineering, Kyoto University Kyoto, Japan John C. de Mello Department of Chemistry Norwegian University of Science and Technology (NTNU) Trondheim, Norway Desta Gedefaw Flinders Institute for NanoScale Science and Technology Flinders University Adelaide, Australia and School of Biological and Chemical Sciences Faculty of Science, Technology and Environment The University of South Pacific Suva, Fiji

ix

x

Karen K. Gleason Department of Chemical Engineering Massachusetts Institute of Technology Cambridge, Massachusetts Masayuki Gon Department of Polymer Chemistry Graduate School of Engineering, Kyoto University Kyoto, Japan Zhijie Guo Materials Research Institute and School of Biological and Chemical Sciences Queen Mary University of London London, United Kingdom Martin J. Heeney Department of Chemistry Imperial College London London, United Kingdom Yun Huang Department of Chemistry University of Texas at San Antonio San Antonio, Texas Shunichiro Ito Department of Polymer Chemistry Graduate School of Engineering, Kyoto University Kyoto, Japan Pradeepkumar Jagadesan Department of Chemistry University of Texas at San Antonio San Antonio, Texas Malika Jeffries-EL Department of Chemistry Boston University Boston, Massachusetts So-Huei Kang Department of Energy Engineering School of Energy and Chemical Engineering Ulsan National Institute of Science and Technology (UNIST) Ulsan, South Korea

Contributors

Thomas Körzdörfer Computational Chemistry Institute of Chemistry University of Potsdam Potsdam, Germany Tanya Kumari Department of Energy Engineering School of Energy and Chemical Engineering Ulsan National Institute of Science and Technology (UNIST) Ulsan, South Korea Emily L. Kynaston Department of Chemistry University of Toronto Toronto, Ontario, Canada Mario Leclerc Département de Chimie Université Laval Québec City, Québec, Canada Sang Myeon Lee Department of Energy Engineering School of Energy and Chemical Engineering Ulsan National Institute of Science and Technology (UNIST) Ulsan, South Korea Hong Li School of Chemistry and Biochemistry and Center for Organic Photonics and Electronics Georgia Institute of Technology Atlanta, Georgia Christine Luscombe Department of Materials Science and Engineering University of Washington Seattle, Washington Christian B. Nielsen Materials Research Institute and School of Biological and Chemical Sciences Queen Mary University of London London, United Kingdom

xi

Contributors

Robert M. Pankow University of Southern California Department of Chemistry Los Angeles, California Zachary S. Parr Materials Research Institute and School of Biological and Chemical Sciences Queen Mary University of London London, United Kingdom Seth C. Rasmussen Department of Chemistry and Biochemistry North Dakota State University Fargo, North Dakota Chad Risko Department of Chemistry and Center for Applied Energy Research (CAER) University of Kentucky Lexington, Kentucky Sean M. Ryno School of Chemistry and Biochemistry and Center for Organic Photonics and Electronics Georgia Institute of Technology Atlanta, Georgia and Department of Chemistry and Center for Applied Energy Research (CAER) University of Kentucky Lexington, Kentucky

Dwight S. Seferos Department of Chemistry Department of Chemical Engineering and Applied Chemistry University of Toronto Toronto, Ontario, Canada Jessica Shaw Department of Chemistry Boston University Boston, Massachusetts Kazuo Tanaka Department of Polymer Chemistry Graduate School of Engineering, Kyoto University Kyoto, Japan Tonghui Wang School of Chemistry and Biochemistry and Center for Organic Photonics and Electronics Georgia Institute of Technology Atlanta, Georgia Xiaoxue Wang Department of Chemical Engineering Massachusetts Institute of Technology Cambridge, Massachusetts Michael O. Wolf Department of Chemistry University of British Columbia Vancouver, British Columbia, Canada

Kirk S. Schanze Department of Chemistry University of Texas at San Antonio San Antonio, Texas

Changduk Yang Department of Energy Engineering School of Energy and Chemical Engineering Ulsan National Institute of Science and Technology (UNIST) Ulsan, South Korea

Brian Schmatz University of Southern California Department of Chemistry Los Angeles, California

Shuyang Ye Department of Chemistry University of Toronto Toronto, Ontario, Canada

1 Early History of Conjugated Polymers: From Their Origins to the Handbook of Conducting Polymers 1.1 Introduction........................................................................................... 1 1.2 Basic Synthesis and Doping Processes of Conjugated Polymers........3 1.3 Polyaniline���������������������������������������������������������������������������������������������5 Early Reports of the Oxidation of Aniline • Determination of the Structure of Aniline Oxidation Products • Buvet, Jozefowicz, and Conducting Polyaniline

1.4 Polypyrrole��������������������������������������������������������������������������������������������11 Angeli and Pyrrole Black • Ciusa and ‘Graphite’ from Pyrrole • Weiss and Conducting Polypyrrole • Pyrrole Black at the University of Parma • Diaz and Electropolymerized Polypyrrole Films

1.5 Polyacetylene���������������������������������������������������������������������������������������� 17 Natta and the Polymerization of Acetylene • Tokyo Institute of Technology and Continued Studies of Polyacetylene • Shirakawa and Polyacetylene Films • Smith, Berets, and Doped Polyacetylene • MacDiarmid, Heeger, and Poly(sulfur nitride) • Doped Polyacetylene Films

1.6 Polythiophene��������������������������������������������������������������������������������������23

1.7

Seth C. Rasmussen

Yamamoto and Polythiophene via Catalytic Cross-Coupling • Lin and Related Catalytic Cross-Coupling Methods • Polythiophene via Electropolymerization • Polythiophenes via Chemical Oxidation

The Rise of Synthetic Metals and a Developing Field of Conductive Polymers��������������������������������������������������������������������������26 Synthetic Metals • Dedicated Literature

References��������������������������������������������������������������������������������������������������������29

1.1 Introduction Modern society is largely a plastic-based culture in which plastics developed from organic polymers have become more ubiquitous than other common materials such as metals, glass, or ceramics. This has led some to postulate that there is sufficient justification to refer to the period beginning with the 20th century as the Age of Plastics [1]. Although common organic plastics comprised of polymers such as 1

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Conjugated Polymers: Perspective, Theory, and New Materials

polyethylene or polystyrene are electrically insulating materials, it was discovered in the 1960s that certain types of organic polymers could be made to exhibit semiconducting properties, with even metallic conductivity demonstrated by the late 1970s. The most common and successful examples of such conductive organic polymers are developed from conjugated organic polymers (Figure 1.1), a class of organic semiconducting materials that exhibit enhanced electronic conductivity in their oxidized or reduced states [2, 3]. These organic materials thus combine the electronic properties of classical inorganic materials with many of the desirable properties of organic plastics, including mechanical flexibility and low production costs. This combination of properties has led to substantial fundamental and technological interest, resulting in the current field of organic electronics and the development of a variety of modern technological applications. Such applications commonly include sensors, electrochromic devices, organic photovoltaics (OPVs), organic light-emitting diodes (OLEDs), and field effect transistors (FETs) [2–7]. In addition, the flexible, plastic nature of the organic electronic materials used as the active layers in such devices has led to the realistic promise of flexible electronics in the near future [4–7]. The ability to imbue organic polymers with electrical conductivity is typically viewed as a somewhat recent advancement, and discussions of the history of these materials generally begin in the mid- to late-1970s with the collaborative work of Hideki Shirakawa, Alan G. MacDiarmid, and Alan J. Heeger on conducting polyacetylene [8–12]. The common view that the field essentially began with this polyacetylene work was further reinforced with the awarding of the 2000 Nobel Prize in Chemistry to these investigators “ for the discovery and development of electrically conductive polymers” [13]. Such a view, however, overlooks the fact that reports of electrically conductive conjugated polymers date back to the early 1960s [14–16], with the study of conjugated polymers in general dating nearly back to the very beginning of the 19th century [17]. In fact, it has been recently argued that polyaniline is the oldest known fully synthetic organic polymer [17], with a nearly continuous string of publications on this material dating back over the last 180+ years. Recently, the current author has worked to educate the conjugated materials community with a series of publications detailing the early history of conjugated polymers and the discovery of their conductive nature when treated with appropriate oxidizing or reducing agents [17–24]. Along with these contributions, two additional historical accounts have been published during this time frame that have also tried to highlight some of the contributions from research that predate the polyacetylene work of the 1970s [25, 26]. In continuing these collective efforts, the present chapter will provide an overview of the known

FIGURE 1.1  Common parent conjugated polymers and the years of their first reports in the literature.

Early History of Conjugated Polymers

3

history of four of the primary parent conjugated polymers—polyaniline, polypyrrole, polyacetylene, and polythiophene—from the origins of polyaniline in the early 19th century up through the development of polythiophenes in the 1980s. While early work in the field may not have been as dramatic or as fully realized as the later polyacetylene studies, many aspects and relationships attributed to the work recognized by the 2000 Nobel Prize can be seen in these earlier contributions.

1.2 Basic Synthesis and Doping Processes of Conjugated Polymers Much of the following history predates the ability to determine the molecular structure of the polymeric materials in question, and even predates the macromolecular model of polymers as introduced by Hermann Staudinger (1881–1965) in the 1920s [27–31]. As such, it is worthwhile to briefly review our modern understanding of the polymerization methods under discussion, as well as the basic redox processes involved in the doping of conjugated polymers, in order to provide context to what will be presented in the following sections. The majority of modern conjugated materials are produced via various transition metal-catalyzed cross-coupling methods, including Kumada, Stille, and Suzuki crosscoupling [32–35], as well as recent efforts in direct arylation polymerization [36–39]. In contrast, however, the early period of conjugated polymers was dominated by oxidative polymerization as the primary synthetic method for the generation of these materials. Electron-rich monomers polymerize anodically via either chemical or electrochemical oxidation of the π-system to form the corresponding radical cation intermediate [40–46], for which multiple resonance forms exist. For the 5-membered heterocycles (thiophene, pyrrole, furan, etc.), spin density studies support the localization of the unpaired electron at the α-position (Figure 1.2) [40, 42, 43]. As such, coupling of the radical cations occurs predominately through the α-positions, followed by deprotonation to give the neutral α,α′-dimer [40–43]. Chain propagation then continues through a step-growth mechanism involving sequential oxidation, coupling, and deprotonation steps. Aniline undergoes oxidative polymerization in a similar manner, but with some important differences [43–46]. As with the heterocycles above, oxidation results in the formation of the corresponding radical cation, which again can exist in multiple resonance forms (Figure 1.3). Spin density studies indicate nearly equal distribution of the unpaired electron between the nitrogen and the para-carbon of the benzene ring [46], thus providing the opportunity for three possible couplings: nitrogen-nitrogen (head-to-head, HH); nitrogen-arene (head-to-tail, HT); and arene-arene (tail-to-tail, TT) [43–46]. Diarylhydrazines formed via HH coupling are not stable and quickly undergo conversion via either disproportionation or the benzidine rearrangement, particularly under acidic conditions [43], and thus

FIGURE 1.2  Oxidative polymerization mechanism of 5-membered heterocycles.

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Conjugated Polymers: Perspective, Theory, and New Materials

FIGURE 1.3  Oxidative polymerization mechanism of aniline.

HH units do not contribute to polymer formation [44]. Of the other two possible regio couplings, the high radical cation concentrations typical of most polymerization conditions favor TT over HT coupling [45]. As with the heterocycles above, initial coupling is followed by deprotonation to give the neutral dimer, after which chain propagation continues through a step-growth mechanism involving sequential oxidation, coupling, and deprotonation steps. The oxidation of conjugated organic polymers generates positive charge carriers (i.e. holes, Figure 1.4) and an increase of p-type character [23, 47, 48]. Thus, conjugated polymers in their oxidized form are referred to as p-doped in analogy to p-doped inorganic semiconductors such as gallium-doped silicon. As the product of this p-doping process is a polycationic material, anionic species must be incorporated into the material to maintain charge neutrality. If p-doping is accomplished via an oxidizing agent, the anions generated by the redox process then become the counterions incorporated into the polymer. In contrast, materials p-doped via electrochemical oxidation incorporate anions from the supporting electrolyte utilized during the electrochemical process [23]. As the polymer products undergo oxidation at lower potentials than the initial monomers, the materials generated via oxidative polymerization are initially produced in their oxidized state and require reduction in order to isolate the neutral form of the polymer. Although less common, some conjugated polymers are also able to undergo reduction, or n-doping, resulting in the addition of negative charge carriers (i.e. electrons, Figure 1.4) and an increase of n-type character [23, 47, 48]. This, too, can be accomplished via either electrochemical reduction or the use of a suitable reducing agent, but in either case cationic species must be incorporated in order to maintain charge neutrality. The counterions integrated into these polymers during p-doping or n-doping are commonly referred to as “dopants,” which can be somewhat misleading as the counterion itself does not cause the enhanced conductivity. However, such counterions are necessary in order to generate the oxidized or reduced forms that do provide the resulting conductive materials [23].

Early History of Conjugated Polymers

5

FIGURE 1.4  Doping of conjugated polymers.

1.3 Polyaniline The earliest conjugated polymer was polyaniline, which has been known by a number of names over the years, including emeraldine and aniline black [17, 23]. First reported in the first half of the 19th century, the discovery of polyaniline was dependent primarily on the isolation of aniline itself, which was independently “discovered” by various researchers. The first of these was the German chemist Otto Unverdorben (1806–1873) who reported the isolation of an oil from the dry distillation of indigo in 1826, which he named krystallin. Eight years later, another German, F. Ferdinand Runge (1794–1867) (Figure 1.5), isolated a volatile oil from the distillates of coal tar that formed colorless salts when treated with acid. Both the oil and its salts became aquamarine when treated with chlorine of lime, which led him to name this oil kyanol [combination of kyanós (Greek for “blue”) and oleum (Latin for “oil”)] [49].

FIGURE 1.5  Friedlieb Ferdinand Runge (1794–1867). (Edgar Fahs Smith Memorial Collection. Kislak Center for Special Collections, Rare Books and Manuscripts. University of Pennsylvania.)

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Conjugated Polymers: Perspective, Theory, and New Materials

Both krystallin and kyanol are now known as aniline, and it was Runge that was the first to oxidize this species in 1834 [17, 23].

1.3.1 Early Reports of the Oxidation of Aniline Shortly after the isolation of his kyanol, Runge treated its nitrate salt with copper oxide in hydrochloric acid to produce a dark green-black color [50]. He then showed that the treatment of either the nitrate or hydrochloride salts with a variety of copper salts resulted in the same reaction and noted that if enough of the aniline salts could be prepared, the colored products resulting from their treatment with metal species could provide a practical use [50]. In an effort to illustrate this, he treated cotton with lead chromate, after which he printed the fabric with aniline hydrochloride to give green patterns within 12 hours. If more concentrated aniline hydrochloride solutions were used, black patterns developed rather than green, with both patterns proving to be color-fast upon rinsing in water [50]. Later in Russia, Carl Julius Fritzsche (1808–1871) (Figure 1.6) isolated a colorless oil from indigo in 1840, which he called anilin after the Spanish name of indigo, añil [51]. A few years later, the German organic chemist August Hoffmann (1818–1892) showed that krystallin, kyanol, and anilin were all the same compound [52], for which Fritzsche’s name was retained as the modern aniline [17]. Fritzsche found that, under the right circumstances, treatment of aniline with nitric acid caused the production

FIGURE 1.6  Carl Julius Fritzsche (1808–1871).

Early History of Conjugated Polymers

7

of a blue or green material, although he was not able to generate enough of the product to adequately characterize. He then found that the addition of aniline salts to chromic acid produced a dark green precipitate, which ultimately became black-blue. Combustion analysis revealed that the precipitate contained significant amounts of chromium [51]. He later returned to the treatment of aniline with oxidants in 1843, this time combining alcohol solutions of aniline salt with potassium chlorate in hydrochloric acid to give a blue precipitate. This blue solid turned green upon washing with alcohol, and ultimately dark green upon drying. Analysis of the product revealed an empirical formula of C24H20N4Cl 2O [53], in near perfect agreement with the chloride salt of the most common form of oxidized polyaniline [17]. True efforts to commercialize these colored products then began in England when Frederick CraceCalvert (1819–1873), Samuel Clift, and their assistant, Charles Lowe, developed green and blue dyes from the oxidation of aniline for the coloring of cotton in 1860 [54–56]. They filed a joint patent for these dyes on June 11, 1860, using methods very similar to those of Runge and Fritzsche that involved the application of an aniline salt (either the hydrochloride or tartate) and potassium chlorate to give a green color after 12 hours. This green color was given the name emeraldine [54–58], which ultimately became the name adopted for the most common form of polyaniline. If the green-dyed fabric was then boiled in either an alkaline or soap solution, the color became a blue that was given the name azurine [54–57]. The printers Wood and Wright commercialized these dyes in late 1860, while also introducing improvements resulting in darker shade that could be considered black [57, 59]. Additional aniline-based black dyes were also introduced by both the English colorist John Lightfoot, Jr. and the German industrial chemist Heinrich Caro (1834–1910) in 1860 [54, 55, 59], with Caro’s black dye commercialized by Roberts, Dale & Co. in 1862 [60] and Lightfoot’s dye commercialized by Jakob J. Muller-Pack in 1863 [59]. These various black dyes from aniline became collectively known as aniline black by 1871 [17, 55], which eventually became the first real general term for polyaniline. Following the introduction of aniline black dyes, the next major innovation came from the English physician and chemist Henry Letheby (1816–1876) (Figure 1.7) [61], who is often incorrectly credited with the first production of polyaniline. Letheby’s study of aniline began with the treatment of acidic aniline solutions with various oxidizing agents to produce blue-to-purple colors in 1862 [61]. Continuing these efforts, Letheby then electrochemically oxidized a sulfuric acid solution of aniline via a Pt electrode at the positive pole of a small Grove cell (an early, high current battery) to generate a deep blue to bluishgreen pigment [61]. Using greater quantities of aniline and two larger Grove cells connected together for intensity, he was then able to prepare the material on a larger scale as a dirty bluish-green pigment that coated the large platinum sheet acting as the positive electrode. This pigment was then removed from the electrode, washed with water, and dried to give a bluish-black powder that was only soluble in sulfuric acid. Dilution of acid solutions with water caused the precipitation of a dirty emerald green powder that could be made blue with concentrated ammonia or turned from blue to purple with concentrated sulfuric acid [61]. This was the first example of the electrochemical oxidation of aniline and the earliest report of this method for the production of conjugated materials.

1.3.2 Determination of the Structure of Aniline Oxidation Products It is important to note that the identity and structure of the oxidation products discussed above were completely unknown, and it was not until the early 1900s that significant efforts to determine the structure of these products were reported. These efforts began in 1896, when Heinrich Caro treated aniline with limited oxidant under cold, alkaline conditions, resulting in the isolation of azobenene and a second yellow product (Figure 1.8). Reduction of the yellow product generated p-amidodiphenylamine, which was found to undergo oxidation to give a yellow species. This oxidation product was proposed to be phenylquinonediimide and believed to be the same as the initial yellow product [62].

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Conjugated Polymers: Perspective, Theory, and New Materials

FIGURE 1.7  Henry Letheby (1816–1876). (By W. & D. Downey albumen carte-de-visite, 1860s © National Portrait Gallery, London.)

FIGURE 1.8  Caro's study of the initial oxidation products of aniline.

Further efforts were then continued by the German organic chemist, Richard Willstätter (1872–1942, Figure 1.9) [63–65], beginning in 1906 with the direct polymerization of Caro’s p-amidodiphenylamine to emeraldine under acidic conditions [63]. A blue species was then separated from this mixture that was viewed as a form of the emeraldine base or azurin, which was crystallized from hexane. Analysis led to a formula of C24H20N4 and the conclusion that this product was the combination of two molecules

Early History of Conjugated Polymers

9

FIGURE 1.9  Richard Willstätter (1872–1942). (Edgar Fahs Smith Memorial Collection. Kislak Center for Special Collections, Rare Books and Manuscripts. University of Pennsylvania.)

FIGURE 1.10  Willstätter's proposed generation of aniline black.

of phenylquinonediimide. Oxidation of the blue base gave a red product with two less hydrogens (C24H18N4), both of which could be reduced to give the leucobase (C24H22N4), which could not be further reduced. After proposing all of the possible combinations of two units of the phenylquinonediimide, it was viewed that the experimental observations best supported the two linear isomers given in Figure 1.10. Willstätter viewed emeraldine as an intermediate in the formation of aniline black and showed that his red oxidation product could undergo polymerization into insoluble black products under a variety of conditions. As such, he believed the production of aniline black proceeded as outlined in Figure 1.10. It

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Conjugated Polymers: Perspective, Theory, and New Materials

was finally concluded that aniline black appears as (C6H4.5N)x, where x equaled a minimum of 8, with the simplest possible formula being C48H36N8 [63]. Willstätter then continued the study of aniline black in 1909, focusing on its oxidation and hydrolysis [64]. In the process, he performed a more detailed analysis of the chemical composition of aniline black, with multiple samples resulting in a mean atomic ratio of C5.97H4.55N. It was concluded that this molecular formula described a material consisting of the structural unit shown in Figure 1.10. Lastly, it was concluded that this constitutional formula provided the prospect of various additional oxidation states and oxygen-containing derivatives formed via hydrolysis [64]. Competing studies starting in 1910 were then reported by the English industrial chemists Arthur G. Green (1864–1941) and Arthur E. Woodhead [66–69]. Although Green and Woodhead agreed that the aniline units were linked by single connections, they felt that these could exist as either linear chains or ring-like structures, and actually favored ring structures over the linear structures of Willstätter. In addition, they found issues with some of Willstätter’s conclusions and found the models to be oversimplified. As such, they reinterpreted the previously reported results and provided additional data to present a more detailed model of the primary oxidation products of aniline. Although they believed that the question of linear chains vs. ring structures was still undecided, they presented these models as linear octameric species for simplicity. As illustrated in Figure 1.11, the primary oxidation products were given the previous names emeraldine and nigraniline. In addition, they acknowledged that there were additional stages of oxidations below and above these states, which were named protoemeraldine and pernigraniline. As previously recognized by Willstätter, all of these states were quinonoid derivatives of the same parent substance, which they named leucoemeraldine. Unlike Willstätter, however, Green was firm that these species were only intermediates and did not represent “true aniline black”, which he maintained to be an azine material. It should be pointed out that Green and Woodhead are typically given the credit for determining the structures of the polyanilines, while the primary structure determination was really accomplished by Willstätter. The contributions of Green and Woodhead were

FIGURE 1.11  Green and Woodhead's names and structures for the series of aniline oxidation products.

Early History of Conjugated Polymers

11

primarily the recognition of the full series of oxidative forms and the associated names, both of which are still applied for the descriptions of modern polyanilines.

1.3.3 Buvet, Jozefowicz, and Conducting Polyaniline Although the basic structures of the aniline oxidation products were known by 1915, it was not until the mid-1960s that the electronic properties of polyaniline were first characterized in any detail. This work was carried out in the laboratory of Rene Buvet (1930–1992) at the Ecole Supérieure de Physique et de Chimie Industrielles de la ville de Paris (ESPCI ParisTech), primarily by Marcel Jozefowicz (b. 1934) [18]. This work initially focused on optimizing methods for the reproducible preparation of polyaniline samples via the oxidative polymerization of aniline using persulfate in sulfuric acid solutions to afford emeraldine sulfate [70]. Efforts to control the level of protonation were then investigated, as well as effects of the counterion employed in the polymer salts. These latter efforts were carried out via neutralization of the initial emeraldine sulfate to generate the emeraldine base, followed by production of emeraldine salts consisting of either chloride or formate counterions. This was then followed with a 1965 report of the redox properties of the various polyaniline materials generated [71]. Optimized methods for the production of emeraldine sulfates of controlled compositions had been developed by 1966, allowing studies of the resulting conductive properties [72–75]. Pressed pellets of the polyaniline materials were indeed found to be quite conductive, with Jozefowicz stating [72]: The conductivity of the polyanilines is very high and classifies these compounds among the best known organic conductors. This conductivity is, without possible dispute, electronic. They went on to show that the conductivity was dependent on both the extent of protonation and the water content of the oxidized polyaniline, with early conductivities ranging from 10−5 to 10 Ω−1cm−1 (or S cm−1 in more modern units) [72–75]. The pH dependence of the material was found to provide the most significant effects, with a linear increase in the log of the conductivity with decreasing pH. In contrast, the effect of water was more limited, but did exhibit an increase in conductivity with increasing water content. By 1969, polyanilines with conductivities as high as 100 S cm−1 had been achieved [75]. Buvet reviewed the electronic characteristics of polyaniline during a lecture presented at the eighteenth meeting of CITCE (Comité International de Thermodynamique et Cinétique Electrochimiques) in April of 1967, and concluded that its conductivity was electronic in nature and not due to any ion transport [25, 74]. However, it has been pointed out by Inzelt that this did not give rise to great excitement at the time [25].

1.4 Polypyrrole Although not known for as long as polyaniline, polypyrrole also has a very long history dating back to 1915 [20]. Polypyrrole is also notable for two significant firsts in the field: it was the first organic polymer reported to exhibit significant conductivity [18, 20] and was the first of the conjugated polymers to be prepared as a plastic film [20]. All of this started with the work of Angelo Angeli (1864–1931, Figure 1.12) at the University of Florence, who was the first to study the oxidation of pyrrole.

1.4.1 Angeli and Pyrrole Black Angeli produced a black precipitate that he named nero di pirrolo (e.g. pyrrole black) in 1915 via the treatment of pyrrole with mixtures of hydrogen peroxide and acetic acid [76, 77]. Typically, this was accomplished by adding 50% H2O2 to a solution of pyrrole in acetic acid, which resulted in the formation of a greenish-brown color that ultimately turned black-brown over the period of a couple of days. A thin black powder could then be isolated via either spontaneous precipitation, dilution of the final solution with water, or addition of aqueous sodium sulfate. The product was only soluble in basic solutions and

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FIGURE 1.12  Angelo Angeli (1864–1931). (Courtesy of the “Ugo Schiff” Chemistry Department, University of Florence, Italy.)

purification was therefore accomplished by dissolving the powder in base, after which either acetic acid or dilute sulfuric acid was added to induce precipitation. The purified solid was then isolated via filtration and dried at 120°C to give a fine, dark brown-to-black powder. In addition to peroxide, Angeli found that pyrrole blacks could also be obtained via the application of nitrous acid [78], potassium dichromate [79] or chromic acid [80], lead oxide [79], potassium permanganate [81], and various quinones [80]. Even oxygen could be used as the oxidant when used in combination with either light or ethylmagnesium iodide [78]. Comparison of the materials produced via these various oxidants ultimately led Angeli to conclude [79]: These facts are of special interest because it shows that the formation of pyrrole blacks is most likely preceded by a process of polymerization of the pyrrole molecule, which takes place more or less rapidly depending on the reagents that are used. The insoluble nature of the pyrrole blacks made attempts to probe the structure of these materials difficult. Analysis by oxidative degradation, however, revealed cleavage products consistent with pyrrole and indole derivatives, leading Angeli to conclude that the pyrrole ring was retained within the composition of pyrrole black. These studies were then extended to functionalized pyrroles, which revealed that the oxidation of various functionalized pyrroles generated colored products but did not result in the typical solid pyrrole black [78, 79]. These collective studies ultimately led Angeli to propose that pyrrole black consisted of units comprised of direct carbon-carbon bonds between pyrroles as shown in Figure 1.13a [79]. It should be noted that Angeli’s proposed structure is remarkably similar to the currently accepted structure for oxidized portions of the polypyrrole backbone (Figure 1.13b). Angeli moved onto other research topics in the early 1920s, although he did return to pyrrole black with a final paper in 1930. As Angeli moved away from pyrrole black, however, another Italian scientist, Riccardo Ciusa (1877–1965), began his own related investigations. In his case, Ciusa was attempting to polymerize various heterocycles (pyrrole, thiophene, and furan) in order to generate graphitic

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FIGURE 1.13  (a) Angeli's proposed basic unit for the structure of pyrrole black and (b) modern resonance structures describing oxidized polypyrrole. (From Angeli, A. 1918. Sopra i neri di pirrolo. Nota. Gazz. Chim. Ital. 48(II):21–25.)

FIGURE 1.14  Ciusa’s proposed structures for [C4NHI]n. (From Ciusa, R. 1925. Su alcune sostanze analoghe alla grafite. Gazz. Chim. Ital. 55:385–389.)

analogues. (From Rasmussen, S. C. 2015. Early History of Polypyrrole: The First Conducting Organic Polymer. Bull. Hist. Chem. 40: 45–55.).

1.4.2 Ciusa and ‘Graphite’ from Pyrrole At the University of Bologna, Riccardo Ciusa began investigating the thermal polymerization of tetraiodopyrrole in 1921 as a potential route to a material that could be considered a type of pyrrole-based graphite [82–84]. These efforts produced a black material with a graphitic appearance by heating tetraiodopyrrole under vacuum at 150–200°C. This material gave an elemental composition corresponding to [C4NHI]n, which Ciusa concluded to be an intermediate in the formation of the desired pyrrole ‘graphite’. He later proposed the two structures given in Figure 1.14 as possible representations of this intermediate species [84]. This intermediate material was then reheated at higher temperature (described as incipient red), which gave a black material with an appearance similar to graphite flakes and an elemental composition consistent with [C4NH]n [83]. These methods were repeated with thiophene and furan to obtain similar results before Ciusa finally investigated the thermal polymerization of hexaiodobenzene. This final effort produced a graphite material that he described to be similar to ordinary graphite, although exhibiting a higher resistivity that was approximately six times that of true graphite [84]. Unfortunately, Ciusa never reported the resistivity of the hetreocyclic graphites and thus it is unknown how they might have compared to either the synthesized or native graphites. Of course, beyond appearances and elemental composition, Ciusa really didn’t report any characterization of these materials. Nearly 40 years later, however, Ciusa’s research became the basis of efforts to produce conductive organic polymers by Donald Weiss and coworkers in Melbourne, Australia [18, 20].

1.4.3 Weiss and Conducting Polypyrrole Beginning in 1959, a team of CSIR (Council for Scientific and Industrial Research) researchers led by Donald Weiss (1924–2008, Figure 1.15) began investigating semiconducting organic polymers as potential electrically-activated and easily regenerated adsorbents [18]. Although initial efforts focusing on xanthene polymers failed to produce materials with low enough resistivity [85], Weiss came across

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FIGURE 1.15  Donald Weiss (1924–2008). (Courtesy of Robert Weiss.)

Ciusa’s reports on pyrrole ‘graphite’ in the process. Based upon Ciusa’s results, Weiss thought this might provide the route to a new type of organic material with the desired conductivity [18]. As Ciusa had not characterized the electrical properties of his pyrrole ‘graphite’, Weiss began with reproducing the preparation of the material in order to study its structure and relative conductivity. The CSIR team, however, did not follow Ciusa’s exact methods of [82–84], but instead used modified conditions in which tetraiodopyrrole was heated under a flow of nitrogen in a rotating flask at temperatures as low as 120°C. The application of the nitrogen flow provided both an inert atmosphere and allowed the transfer of iodine vapor away from the reaction. These conditions were reported to produce black, insoluble powders, which Weiss described as “polypyrroles” comprised of [14]: ... a three-dimensional network of pyrrole rings cross-linked in a nonplanar fashion by direct carbon to carbon bonds. Analysis of the products led to the conclusion that the materials contained “adsorbed molecular iodine,” as well as nonreactive iodine presumed to be “iodine of substitution” [14]. Various descriptions of the polymeric material given by Weiss and coworkers are consistent with the hypothetical structure given in Figure 1.16. Measurement of the polypyrrole resistivity (R) as pressed pellets revealed values of 11–200 Ω cm at 25°C, corresponding to conductivities (1/R) of 0.005–0.09 S cm−1 [16, 86]. Determination of the temperature dependence for the resistivity also revealed a temperature profile consistent with a standard semiconductor. Overall, the measured conductivities were well below that of carbon black, but they were significantly greater than the previous xanthene polymers and represented the highest known conductivities to date for a non-pyrolyzed organic polymer. Weiss described the conductivity of the material as follows [16]: However it is apparent that the polymers are relatively good conductors of electricity. Since no polarization was observed during the measurement of the electrical resistance, even over substantial periods of time, it is assumed that the conductivity is of electronic origin.

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FIGURE 1.16  Hypothetical structure of Weiss's polypyrrole.

Perhaps of greater importance than the magnitude of the polymers conductivity was the discovery that removal of the adsorbed molecular iodine via solvent extraction [16], chemical or electrochemical reduction [14, 15], or thermal vacuum treatment [16], all resulted in a corresponding increase in material’s resistance. Further study of this relationship via electron spin resonance (ESR) revealed evidence of a strong charge-transfer complex between the polymer and iodine [15], from which Weiss made the following conclusion [16]: Charge-transfer complexes of strength sufficient to cause partial ionization induce extrinsic [semiconductor] behaviour by changing the ratio of the number of electrons to the number of holes. Furthermore, he went on to state [14]: The presence of the oxidant iodine, and in its absence oxygen, facilitates oxidation of the polymer. Of course, this oxidative process describes what is now called p-doping of the polymer and was ultimately determined to be the key in producing highly conductive organic polymers [2, 9, 10]. Weiss admitted, however, that the full role of the iodine oxidation in determining the polymer conductivity was not realized at the time [18, 20].

1.4.4 Pyrrole Black at the University of Parma As Weiss and his team were concluding their work with their polypyrrole-iodine materials, a new resurgence in the study of Angeli’s pyrrole black was occurring in northern Italy at the University of Parma, which involved a collaboration between Luigi Chierici (d. 1967), Gian Piero Gardini (d. 2001), and Vittorio Bocchi [20]. While the specifics of this collaboration are unclear, Chierici appears to have been the guiding force in these efforts as he was studying pyrrole black before the others had arrived at Parma. These three did not work together for very long, however, as Chierici died in 1967, leaving Gardini and Bocchi to continue these efforts. The majority of these efforts concentrated on identifying the intermediates and byproducts generated during the oxidative polymerization of pyrrole via peroxide [87–89]. However, the most critical results were generated via a collaboration with Parma’s Institute of Physics that focused on ESR studies of pyrrole black. Initially, these studies utilized materials produced via Angeli’s initial H2O2/acetic acid conditions [90]. By 1968, however, these efforts utilized polymers obtained via electrolysis [91],

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thus representing the earliest known example of an electropolymerized polypyrrole. These methods produced a laminar film on the Pt electrode through the application of a constant current of 100 mA in a H2SO4 solution of pyrrole over a period of two hours. The generated polypyrrole film was then rinsed with distilled water and dried under a vacuum [91]. Characterization of the film by X-ray analysis indicated an amorphous material, and electronic measurements gave a room temperature conductivity of 7.54 S cm−1, considerably higher than that for the thermally-produced polypyrrole-iodine materials reported by Weiss [16]. Chierici and Bocchi went on to analyze the composition of the electropolymerized products via oxidative degradation [92]. Similar to the pyrrole blacks produced via H2O2 oxidation, the major degradation product of the electropolymerized materials was pyrrole-2,5-dicarboxylic acid; although, some additional products of unknown composition were also detected. These collective results led to the conclusion that all of the pyrrole blacks consisted of chains of α,α′-linked pyrroles [92]. In 1975, Gardini then started the first of multiple stays as a visiting scientist at the IBM Research Laboratory in San Jose, California [20], where he began working with Arthur Diaz.

1.4.5 Diaz and Electropolymerized Polypyrrole Films Arthur F. Diaz (b. 1938) arrived at IBM in the mid-1970s, where he was given the task of developing a new project of significant impact. As IBM was interested in building capabilities in electrochemistry, a new project with this focus was favorable, ultimately leading to research on modified electrodes [20]. As conducting polymers were a hot topic at the time, Diaz considered using such materials to modify electrodes, but he was unsure how polyacetylene might be successfully used in this way. The solution was then provided by Gardini, who was visiting IBM at the time, when he shared with Diaz the work being done at Parma on pyrrole black, particularly the most recent electropolymerization efforts [20]. Diaz found the combination of the material’s intractability and conductivity attractive and thus began investigating electropolymerized polypyrrole films. Ultimately, he developed controlled electropolymerization conditions which allowed the repeatable production of strongly adhered polymer films on platinum electrodes [93]. In the process, it was found that deoxygenated aprotic solvents resulted in better material properties [93, 94] than the previous aqueous conditions utilized at Parma [91]. Optimum conditions for the production of polypyrrole were found to be galvanostatic polymerization of the films on Pt from pyrrole in a 99:1 CH3CN-H2O mixture with Et4NBF4 as a supporting electrolyte [93, 94]. It was also found that film adherence was affected by the solution water content, with the absence of water giving poorly adhering, non-uniform films, while increased water content improved overall film adherence [94]. The films were analyzed via elemental analysis, which was consistent with a composition comprised of a backbone of coupled pyrrole units, as well as BF4− anions, in a ratio of ~4:1 [93, 94]. The proposed pyrrole-linked structure was then confirmed by Raman and reflective IR analysis, and it was concluded that this polymer backbone carried a partial positive charge that was balanced by the BF4− anions [95, 96]. Finally, analysis of the film by electron diffraction suggested low crystallinity, exhibiting only diffuse rings corresponding to a 3.4 Å lattice spacing [95]. Free-standing polypyrrole films from 5 to 50 µm thicknesses were evaluated via four-point probe to give room temperature conductivities between 10 and 100 S cm−1, much higher than previously achieved at Parma [93–96]. It was thought that this improvement in conductivity was at least partially due to higher quality films as the result of slower growth rates and very limited thicknesses [96]. This viewpoint is consistent with the modern view that the structural order of electropolymerized films decrease with the corresponding film thickness, resulting in decreased conductivity in the film. The electropolymerized polypyrroles were also characterized via temperature-dependent conductivity measurements to reveal a temperature profile consistent with a classical semiconductor [95, 96].

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FIGURE 1.17  Giulio Natta (1903–1979). (Courtesy of the Giulio Natta Archive.)

1.5 Polyacetylene Although doped polyacetylenes are most often presented as the origin of conducting polymers, polyacetylenes are actually fairly late examples of conjugated polymers [19, 24]. The polymerization of acetylene dates back to 1866 with the work of the French organic chemist Pierre Eugène Marcellin Berthelot (1827–1907), but these efforts did not generate a conjugated polymer. Instead, all early polymerization efforts resulted in a crosslinked three-dimensional material later known as cuprene [24, 97]. The linear conjugated polymer would have to wait until 1955 when the Italian polymer chemist Giulio Natta (Figure 1.17) reported the first successful generation of polyacetylene. In addition to its relatively late entry in the history of conjugated polymers, polyacetylene is also distinctive in that it was the first conjugated polymer not generated via oxidative polymerization.

1.5.1 Natta and the Polymerization of Acetylene After successes in the application of catalytic polymerization to α-olefins and diolefins in the early 1950s, Giulio Natta (1903–1979) expanded his scope and began investigating the application of the previously successful metal-based polymerization catalysts to acetylenes [98–100]. These efforts generated an initial Italian patent in 1955, which encompassed the polymerization of acetylene and its derivatives using organometallic catalysts of group 4–8 transition metals [98]. Natta then presented some of his initial results on acetylene polymerization in July of 1957 at the XVI International Congress for Pure and Applied Chemistry in Paris, a summary of which was included as part of a report of the meeting published in Angewandte Chemie later that same year [99]. This initial report was then followed by a full 1958 publication detailing the successful catalytic polymerization of acetylene via combinations of

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FIGURE 1.18  Natta’s optimized conditions for the polymerization of acetylene.

triethylaluminium (Et3Al) and titanium alkoxides [100]. As outlined in Figure 1.18, the best results were obtained via the addition of acetylene to a heptane solution of Et 3Al and titanium(IV) propoxide at 75°C over a period of 15 hours, with a catalyst molar ratio (Al:Ti) of 2.5. These conditions resulted in a 98.5% conversion of monomeric acetylene to give a black, crystalline polymer that was completely insoluble in organic solvents [100]. Characterization of powder samples by X-ray diffraction revealed a crystalline content of ~90–95% and it was determined that the collected X-ray data were consistent with linear chains of polyacetylene in which the double bonds were concluded to be predominantly trans in configuration [100]. The combination of the black color, metallic luster, and the relatively low electrical resistivity (~1010 Ω cm, compared to 1015–1018 Ω cm for typical polyhydrocarbons) led to the conclusion that these polyacetylene products consisted of long sequences of conjugated double bonds and thus were structurally identical to a very long polyene. Shirakawa later stated, however, that this conclusion was not accepted widely at the time [101, 102]. Although the samples exhibited poor solubility, the polyacetylene materials were found to be fairly reactive, particularly with oxidants such as O2 and Cl2 [100]. For example, the materials rapidly absorbed atmospheric O2 at elevated temperatures to give more lightly colored products. In a similar manner, reaction with Cl2 resulted in the production of a white solid that was found to be amorphous by X-ray characterization. It was found that heating this white chlorinated product at 70–80°C caused a rapid loss of HCl accompanied by a darkening of the polymer. Alternately, treatment with elemental potassium caused removal of chlorine to give a black amorphous solid. Although Natta stated that his 1958 report represented only an initial communication with additional publications planned [100], no additional studies on polyacetylenes were ever published by Natta and his coworkers. Other groups, however, did not hesitate to continue this work, resulting in a gradual replacement of the older term, polyene, with the formal name, polyacetylene [101, 102]. These continuing efforts focused not only on the further study of the polymerization reaction and the materials generated, but on new catalytic systems as well. The most significant of these polyacetylene studies were carried out by Masahiro Hatano and Sakuji Ikeda at the Tokyo Institute of Technology [24].

1.5.2 Tokyo Institute of Technology and Continued Studies of Polyacetylene It was at the Chemical Resources Laboratory of the Tokyo Institute of Technology that Masahiro Hatano (b. 1930) reported the first detailed study characterizing the semiconducting properties of polyacetylene in 1961 [103]. This began with study of the Ziegler-Natta catalyzed polymerization conditions, but with a focus on the effect of various reaction parameters on polymer crystallinity and electronic properties. The most important conclusions were that titanium(IV) propoxide gave more crystalline materials than titanium(IV) chloride and that crystallinity generally increased with the reaction temperature [103]. Hatano then characterized the polymer products via electron spin resonance (EPR) spectroscopy and pressed-pellet DC conductivity measurements to show that samples of increased crystalline nature exhibited a greater concentration of unpaired electrons and lower resistance (as low as 1.4 ×104 Ω cm). Finally, the resistance was determined over the range of ~20–125°C, which revealed a temperature dependence typical of intrinsic semiconductors. The following year Hatano and coworkers continued these efforts with the study of new types of polymerization catalysts, with the goal of producing highly crystalline polyacetylene [104]. Hatano then published two additional papers [105, 106], the first of which focused on the effect of pressure on the conductivity of polyacetylene. Held at a constant pressure, the conductivity of highly crystalline samples

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was measured at temperatures over the range of room temperature to that of liquid nitrogen [105]. The current-voltage linearity exhibited an Ohmic relationship, and the temperature dependence was typical of intrinsic semiconductors as previously observed at atmospheric pressure. However, as the pressure was increased, the overall conductivity also increased. This led to the conclusion that the charge carrier density changed under pressure, but that the charge carrier mobility did not change [105]. Hatano’s final polyacetylene paper attempted to use his previous research to present an improved model of the polymer structure, a polymerization mechanism that could account for these structural aspects, and a summary of the known structure-conductivity relationships [106]. Much of this involved a number of detailed investigations of both polyacetylene and its deuterated analogue via IR spectroscopy, which led to the conclusion that the polymer consisted of both cis and trans content and that the relative cis/trans content was dependent on the polymerization temperature [106]. Finally, a proposed polymerization mechanism was presented in which addition of acetylene to the catalyst produced a cisalkene unit and sequential additions resulted in a trimeric intermediate structure, which could either cycle back to eliminate benzene from the catalyst or continue in the growth of a cisoid open chain. This cisoid conformation could also undergo isomerization to a corresponding transoid conformation, although this process was believed to have a sufficiently large activation energy and was still thought to be mediated by the catalyst [106]. Hatano left Tokyo Institute of Technology in 1967 to move to Tohoku University, after which he did not continue further studies of polyacetylenes. His colleague, Sakuji Ikeda (1920–1984), however, continued the tradition of polyacetylene studies at Tokyo Tech, beginning in 1963 with further study of the polymerization mechanism via isotopic labeling [107–109]. Ikeda also continued the previous efforts of Hatano to study temperature effects in determining the relative cis vs. trans configuration of the polyacetylene backbone. Analysis of samples polymerized at temperatures over the ranges of −78 to 80°C revealed a clear trend of increasing trans configuration with increasing temperature [110]. More significantly, however, Ikeda also correlated the relative amount of benzene formed at each temperature and found corresponding temperature effects for benzene formation. Of course, the most well-known work to have come from Ikeda’s group was the discovery and study of a new form of polyacetylene. Reported in the early- to mid-1970s, this work was carried out primarily by Ikeda’s new assistant, Hideki Shirakawa [19, 24].

1.5.3 Shirakawa and Polyacetylene Films Hideki Shirakawa (b. 1936, Figure 1.19) joined Ikeda’s group in April of 1966 as a research associate [18, 19, 24, 111, 112] and was focusing on the study of polyacetylene by the fall of 1967. The discovery that polyacetylene could be produced as lustrous, silvery films has since achieved near mythical status that has been retold many times over the years, with the story rarely told the same way twice [24]. What is known is that in October of 1967 [111, 113], Hyung Chick Pyun (b. 1926), a visiting Korean scientist from the Korea Atomic Energy Research Institute (KAERI) [24], was preparing a sample of polyacetylene using conditions nearly identical to those previously reported by Natta [114] and Hanato [103]. However, instead of generating the typical powder samples usually produced, these efforts produced ragged pieces of a polymer film [111, 115]. In order to understand what had occurred, Shirakawa reviewed the experimental conditions used by Pyun to ultimately find that the catalyst concentration used had been 1000 times greater than intended [101, 102, 111, 112]. Shirakawa posed the following as an explanation for the error [111]: I might have missed the “m” for “mmol” in my experimental instructions, or the visitor might have misread it. In contrast, Alan MacDiarmid has given a quite different account, stating [116]: ... and he [Shirakawa] replied that this occurred because of a misunderstanding between the Japanese language and that of a foreign student who had just joined his group.

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FIGURE 1.19  Hideki Shirakawa (b. 1936). (Adapted from Hall, N. 2003. Twenty-five years of conducting polymers. Chem. Commun. 1–4. With permission of the Royal Society of Chemistry.)

Interestingly, Alan Heeger has given nearly the same account as MacDiarmid, stating [117]: Then he [Shirakawa] had a Korean visitor who misunderstood what he said in Japanese and instead of making the catalyst in the millimolar concentration, he made it in molar concentration and out came something very different. Pyun had grown up in Korea during the years that the country was under Japanese occupation (1910– 1945), however, and it has been confirmed that he spoke fluent Japanese [24, 115]. As such, this casts significant doubt on the versions given by MacDiarmid and Heeger [24]. Whatever the actual reason, the high catalyst content accelerated the rate of polymerization to the point that, rather than occurring in solution to give the typical black precipitate, the acetylene polymerized at the air-solvent interface or along the wetted walls of the vessel to give the observed silvery films [111–113]. Still, whatever the reason of the error, refinement of these new conditions ultimately allowed Shirakawa to reproducibly generate silvery plastic polyacetylene films via polymerization of acetylene on the surface of unstirred, concentrated catalyst solutions [114, 118–122]. Consistent with the previous studies of polyacetylene powders [110], the backbone configuration of the polymer films was strongly temperature dependent. An irreversible isomerization of the cis to trans forms was found to occur above 145°C and thus careful control of the polymerization temperature allowed the isolation and characterization of films consisting of either the all-cis or all-trans structures (Figure 1.20). Copper-colored films of the all-cis form exhibited conductivities of 10 –9–10 –8 S cm–1, while silver-colored films of the all-trans form exhibited higher conductivities (10 –5 –10 –4 S cm–1) [122], though no greater than that previously reported for highly crystalline polyacetylene powders [103]. This was somewhat surprising as the films were expected to provide increased order, and previous studies had shown that conductivity increased with crystallinity [101]. Characterization of the polymer films by X-ray diffraction [119] gave results nearly identical to Natta’s previous studies [100].

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FIGURE 1.20  Temperature dependence of acetylene polymerization.

1.5.4 Smith, Berets, and Doped Polyacetylene Shortly before the accidental discovery of the first polyacetylene films, Donald J. Berets (1926–2002) and Dorian S. Smith (1933–2010) began studying the effects of various gaseous additives on the conductivity of polyacetylene powders at the American Cyanamid Company [123]. These efforts initially focused on the effect of oxygen on the conductivity of polyacetylene pressed pellets, finding that lower resistivity was observed in samples with lower oxygen content. In the process, however, they observed a surprising result [123]: On admission of 150 mm pressure of oxygen to the measuring apparatus (normally evacuated or under a few cm pressure of He gas), the resistivity of polyacetylene decreased by a factor of 10. If the oxygen was pumped off within a few minutes and evacuation continued at 10–4 mm pressure for several hours, the original electrical properties of the specimen were restored. This ultimately led to the conclusion that polyacetylene first adsorbed oxygen in a reversible manner, causing a reduction in resistivity. Upon prolonged exposure, however, the normally observed increase in resistivity occurred due to an irreversible chemical reaction between the polymer and oxygen. The study was then expanded to study the effects of other gases on the polymer conductivity. Various electron acceptors (BF3, BCl3, Cl2, SO2, NO2, O2, etc.) were all found to give decreases in resistivity, although oxidizing gases such as oxygen and chlorine ultimately caused chemical reaction with the polymer. In contrast, electron donors (NH3, CH3NH2, H2S, etc.) had the opposite effect. Of the gases investigated, the most dramatic results were obtained using BF3 to cause the conductivity to increase by three orders of magnitude (to ~0.0013 S cm–1). Berets and Smith explained these results as follows [123]: The effect on conductivity of the adsorbed electron-donating and electron-accepting gases is consistent with the p-type nature of the specimens... If holes are the dominant carriers, electron donation would be expected to compensate them and reduce conductivity; electron acceptors would be expected to increase the concentration of holes and increase conductivity; this is observed. Although they did not completely understand the interaction of the gaseous additions with polyacetylene, they quite clearly state that the “electrical conductivity... depended on the extent of oxidation of the samples” [123]. However, these results did not seem to generate much interest at the time, and the authors never followed this initial paper with any further studies.

1.5.5 MacDiarmid, Heeger, and Poly(sulfur nitride) Nearly a decade after the Smith and Berets report, a related study involving the addition of gaseous Br2 to the inorganic polymer poly(sulfur nitride), (SN)x, was being investigated at the University of

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FIGURE 1.21  Alan G. MacDiarmid (1927–2007) and Alan J. Heeger (b. 1936). (Adapted from Hall, N. 2003. Twentyfive years of conducting polymers. Chem. Commun. 1–4. With permission of the Royal Society of Chemistry.)

Pennsylvania [124–127]. The basis of this work was a collaboration that began in 1975 after Alan J. Heeger (b. 1936, Figure 1.21) had become intrigued by reports of the metallic nature of poly(sulfur nitride) [112, 116, 128]. After learning that his colleague Alan G. MacDiarmid (1927–2007, Figure 1.21) had previous experience with sulfur nitride chemistry, Heeger approached MacDiarmid about working together on studies of this new polymer. Methods for the preparation of the polymer via the solid-state polymerization of S2N2 were initially developed in order to ensure access to high-quality samples of the material. The developed methods resulted in a lustrous golden material, which embodied the first reproducible preparation of analytically pure (SN)x [124, 125]. The following year, they then reported the material’s electronic properties, which exhibited conductivities of 1.2–3.7 × 103 Ω–1cm–1 [126]. Finally, as previous reports had noted that (SN)x reacted with halides, they synthesized the derivative (SNBry)x via treatment of the material with Br2 vapor. Relative to pristine (SN)x, this brominated derivative was found to exhibit a 10-fold increase in conductivity [127].

1.5.6 Doped Polyacetylene Films Not long after successfully producing high-quality samples of poly(sulfur nitride), MacDiarmid spent time at Kyoto University as a Visiting Professor [116]. While in Japan, he was invited to speak at the Tokyo Institute of Technology. After his lecture, MacDiarmid was invited to tea with the head of the chemistry department, and it was there that he met Hideki Shirakawa [112, 116]. Shirakawa had not attended his lecture and so MacDiarmid was showing him a sample of his golden (SN)x film over tea [24]. Upon seeing the film, Shirakawa mentioned that he had a similar material and returned to his lab to retrieve a sample of his silver polyacetylene film to show MacDiarmid [19, 24]. MacDiarmid was quite interested in the silver film, and he arranged for funding to bring Shirakawa to the United States to do some additional work on polyacetylene [116]. Thus, in September of 1976, Shirakawa began working with MacDiarmid and Heeger as a visiting scientist at the University of Pennsylvania [111]. The initial focus of Shirakawa and MacDiarmid was to improve the purity of the polyacetylene films [116]. As discussed above, Smith and Berets had shown that reducing oxygen content increased the polymer conductivity [123], and thus it was reasoned that increasing the purity of the films could potentially further increase its conductivity. These efforts ultimately resulted in films with purities as high as

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~98.6% [19, 24]. Surprisingly, however, it was found that the film’s conductivity decreased with enhanced purity [116]. It was then proposed that perhaps the film impurities were acting as dopants that increased the polyacetylene conductivity [116], similar to the effects previously seen by Heeger and MacDiarmid via the treatment of (SN)x with Br2 [127]. Further support for this possibility came from previous investigations by Shirakawa and Ikeda on the reaction of polyacetylene with Cl 2, during which in situ IR measurements had revealed a dramatic decrease in IR transmission [101]. This IR response suggested that the initial halogen-treated material might have unusual electronic properties, and thus it was collectively decided to study the effect of Br2 addition on the conductivity of the polyacetylene films. It was then on November 23, 1976, that Shirakawa and Dr. Chwan K. Chiang, a postdoctoral fellow working under Heeger, made the first of these measurements [101, 111]. Using a high-quality film of transpolyacetylene, the conductivity was monitored by four-point probe while being exposed to gaseous Br2 [8, 19, 24, 112]. The conductivity of the film increased rapidly under a Br2 pressure of 1 Torr, rising from 10–5 to 0.5 S cm–1 within a span of only 10 minutes. Even greater increases in conductivity (as high as 38 S cm–1) were then achieved when the measurement was repeated using I2 vapor in place of Br2 [8], optimization of which ultimately produced conductivities as high as 160 S cm–1 later that same year [9]. Although the initially reported conductivities were similar to that previously reported for polyaniline materials [72–75], these optimized results were noteworthy as this was the first report of a conducting polymer fully in the metallic range. Further refinement of the doping process found that even higher conductivities could be achieved by replacing I2 with AsF5 [9, 10]. Thus, AsF5-treated trans-polyacetylene films gave conductivities of 220 S cm–1, while AsF5-treated cis-polyacetylene gave even higher values (560 S cm–1). The significantly higher values achieved via cis-polyacetylene films then led the researchers to revisit the use of I2 on cispolyacetylene in 1978 to produce conductivity values above 500 S cm–1 [11]. At the same time, it was also demonstrated that polyacetylene could be doped with electron-donating species, such as sodium, to give conductivities of 8 S cm–1 [11]. In a final 1978 paper, the doping of polyacetylene with electron donors had been optimized to give conductivities as high as 200 S cm–1 [12].

1.6 Polythiophene Unlike the previous three parent conjugated polymers, the history of polythiophene only dates back a few decades and is the only example discussed in the current chapter that does not predate the seminal polyacetylene work of the 1970s. Several early papers described the polymerization of thiophene via treatment with various acids, but the resulting products were of quite low molecular weight and exhibited limited conjugation [129–131]. This was then followed by a 1971 patent on the electropolymerization of heterocycles, including thiophene, in aqueous mixtures of acetic acid and KOH, but no characterization was reported beyond elemental analysis of the products [132]. It was in 1980, however, that the first true reports appeared, with multiple groups reporting the preparation and characterization of polythiophene materials nearly simultaneously. The earliest of these reports came from Takakazu Yamamoto at the Tokyo Institute of Technology.

1.6.1 Yamamoto and Polythiophene via Catalytic Cross-Coupling After applying transition metal-catalyzed cross-coupling to the preparation of polyphenylene and polymethylene in the late 1970s, Takakazu Yamamoto (b. 1944) turned to the application of these methods to the polycondensation of 2,5-dibromothiophene in 1979. As reported in January of 1980 [133], this involved the addition of a single equivalent of magnesium to 2,5-dibromothiophene, followed by polymerization via catalytic cross-coupling with NiCl2 (Figure 1.22) to give a black precipitate. Low molecular weight material was then removed via sequential extraction with hot methanol and hot CHCl3. The remaining insoluble product was 78% of the initial precipitate and exhibited only a single sharp band via IR analysis at 788 cm−1. Conductivity measurements of the purified polymer gave a value of 5.3 × 10−11 S cm−1, which increased to 3.4 × 10−4 S cm−1 after doping with I2.

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Conjugated Polymers: Perspective, Theory, and New Materials

FIGURE 1.22  Catalytic polycondensation of 2,5-dibromothiophene.

This was then followed up in 1981 with the preparation of the polymer of 2,4-dibromothiophene and comparison to the previous α,α′-coupled isomeric polymer [134]. Although the conductivity of the neutral material was fairly comparable (5 × 10−13 S cm−1), I2 doping resulted in very little enhancement of the conductivity. By this point, the conductivity of the I2 doped poly(α,α′-thiophene) had been increased to 4 × 10−2 S cm−1. The following year, these methods were extended to functionalized analogues, beginning with 3-methylthiophene [135]. The overall properties were extremely similar to that of the unfunctionalized patent, although with increased solubility in CHCl3. The CHCl3 soluble fraction made up 95% of the material, which was found to have a Mn of 2400 (n = ~25). The I2-doped product was found to exhibit a room temperature conductivity of 2.8 × 10−2 S cm−1. A full paper compiling the results to date was then reported in 1983 [136], followed quickly by a more detailed study of the doping of polythiophene with both I2 and SO3 [137]. The maximum conductivity of iodine-doped polythiophenes was essentially the same as reported above, while that of the SO3-doped materials was ~1 × 10−3 S cm−1. Two additional reports of using I2-doped polythiophenes as positive electrodes in Galvanic cells were also reported [138, 139].

1.6.2 Lin and Related Catalytic Cross-Coupling Methods Approximately five months after Yamamoto’s initial report on polythiophene, John W.-P. Lin and Lesley P. Dudek at Xerox reported nearly identical methods, differing really only in the metal catalyst applied [140]. In this latter case, various metal acetylacetonate (acac) complexes were used in comparison to the NiCl2 applied by Yamamoto. Although nickel, iron, and cobalt complexes all catalyzed the polymerization, it was the use of Ni(acac)2 that gave the best results. Overall, the properties of the resulting polythiophene product was nearly identical to that of Yamamoto, although Lin was able to obtain higher conductivities for I2-doped samples (up to 0.1 S cm−1). It was also shown that the polymer conductivity could be enhanced by thermal annealing. Lin never published any follow-up papers on this initial report.

1.6.3 Polythiophene via Electropolymerization The first report of electrochemically produced polythiophenes appeared in late 1980 as a very brief communication by V. L. Afanas’ev and coworkers in Russia [141]. The experimental methods were based on those applied to the electropolymerization of pyrrole by Diaz and coworkers [93] to give oxidized polythiophene films balanced by BF4− counterions. Conductivities of 10−3 S cm−1 were reported but little additional details were provided. This was then followed a few months later by a preliminary paper of Diaz and coworkers at IBM in early 1981 [142], which provided a correlation of the potential of oxidation with the number of repeat units in oligomers and polymers of pyrroles, phenylenes, and thiophenes. Although this initial paper did not detail the preparation of the polythiophenes used, the authors stated that the polymers were previously prepared in an unreported 1980 study. A more extensive report was given in 1983, however, which provided the full details for the electropolymerization of thiophene, as well as a variety of functionalized derivatives [143]. The polythiophene films were produced via nearly identical conditions to the previous polypyrrole work to give either polythiophene-BF4 or polythiophene-PF6 films. Although a significant number of polythiophene derivatives were successfully electropolymerized, only polymers of thiophene, bithiophene, and 3-methylthiophene could be produced thick enough to provide

Early History of Conjugated Polymers

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free-standing films that could be peeled off from the Pt electrode. These films exhibited typical thicknesses of 10−4–10−3 cm with conductivities of 0.02–1 S cm−1 [143], considerably higher than that previously reported by Afanas’ev and coworkers [141], but still lower than that previously obtained for polypyrroles [93–96]. The highest conductivity was found for the poly(3-methylthiophene)-PF6 films [143]. This was then followed up by two additional papers that primarily focused on the substituent effects on the electrochemical and electrochromic properties of the films [144, 145]. Expansion of the IBM group’s work on polypyrrole to other heterocycles such as thiophene was also reported in France by Gérard Tourillon and Francis Garnier in early 1982 [146]. The methods utilized were similar to that of Diaz and coworkers [93], with the majority of films electropolymerized onto Pt from a CH3CN-Bu4NClO4 solution. However, other solvents, electrolytes, and electrode substrates were also investigated [146]. In this way, polythiophene-ClO4 films with conductivities of 10–100 S cm−1 could be generated, which were considerably higher than any of the previously reported conductivities discussed above. The authors attributed this increase in conductivity to reduced impurities in their electropolymerized products. This was then followed with a more detailed investigation of the effect of the dopant (counterion) on the properties of the electrogenerated films, as well as the study of electropolymerized materials of 3-methyl- and 3,4-dimethyl-thiophene [147]. In the process, it was determined that the conductivity increased with the doping level, although the maximum conductivity was concluded to be limited by the polymer morphology. Fully de-doped polythiophenes were determined to have a conductivity of ~10−7 S cm−1. A second 1983 paper then focused on the enhanced stability of polythiophenes in comparison to other known conducting polymers such as polyacetylene and polypyrrole [148]. Also in 1982, but some six months after the initial report of Tourillon and Garnier, came further reports of electropolymerized polythiophenes by Keiichi Kaneto and coworkers in Japan [149]. Again, this work was based on the previous efforts of Diaz and coworkers on polypyrrole [93] and utilized conditions very similar to those described above, using AgClO4 as the electrolyte and an indium-tin oxide (ITO)-conducting glass-working electrode. Electropolymerization resulted in the formation of a dark greenish film that could be easily removed from the electrode. The thickness of the film was ~10 µm with a room temperature conductivity of 0.6 S cm−1 [149]. Varying the solvent-electrolyte combinations applied and further optimization of the polymerization conditions then resulted in films with enhanced conductivities of 20–106 S cm−1, similar to that reported by Tourillon and Garnier [146, 147]. It was found that metallic films could be grown from LiBF4/PhCN [150] and that all of the polythiophene films could also be either electrochemically or chemically de-doped to give neutral films with a band gap of 2.0 eV and a conductivity of 2 × 10−8 S cm−1 [150, 151].

1.6.4 Polythiophenes via Chemical Oxidation Following the initial reports of the electropolymerization of thiophene, Gerhard Koßmehl (b. 1934) and Georg Chatzitheodorou reported the first example of polythiophene via the chemical oxidation of thiophene in the fall of 1981 [152]. These efforts began with the AsF5 doping of polythiophenes produced via Yamamoto’s polycondensation methods [133] to achieve conductivities up to 1.4 × 10−3 S cm−1 [152]. However, it was found that conducting polymeric materials could also be produced by the reaction of either thiophene or 2,2′-bithiophene with AsF5, with conductivities up to 0.021 S cm−1. These materials could then be de-doped with 25% aqueous ammonia to give materials with conductivities similar to the neutral polythiophenes generated via polycondensation. Characterization of the de-doped materials via mass spectroscopy suggested that the neutral polymer was made up of thiophene oligomers up to the nonamer [152]. Potential polymerization mechanisms were also proposed beginning from the radical cation of thiophene. This was then followed by a second paper on polythiophene-AsF5 materials in 1982 [153] before expanding the scope of this approach by replacing AsF5 with the nitrosonium salts, NOSbF6 and NOBF4, in 1983 [154]. These efforts generated conductivities as high as 0.07 S cm−1 and analysis by mass spectroscopy, which suggested these materials to be of higher molecular weight than the previous AsF5 analogues, with measurable oligomers up to n = 16.

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Conjugated Polymers: Perspective, Theory, and New Materials

1.7 The Rise of Synthetic Metals and a Developing Field of Conductive Polymers As illustrated by the various histories outlined above, conjugated and conducting polymers have a very long history, with significant development before the seminal polyacetylene work of the late 1970 that is most often viewed as the point of origin for these materials. As such, it is hard to justify claims that Shirakawa, MacDiarmid, and Heeger discovered electrically conductive polymers. This is not to say, however, that their work in the 1970s was not a critical turning point in the history of these materials. Although oxidized forms of polyaniline had approached the threshold of metallic conductivity by 1969 [75], it was the doped polyacetylenes that provided the first concrete example of metallic organic polymers. Another important aspect of the polyacetylene work was that this was really the beginning of the multidisciplinary approach to the study and development of conjugated/conducting polymers, with the bulk of the previous work all carried out by various industrial and academic chemists. Finally, it is quite clear that the plastic films of polyacetylene combined with both their silver appearance and metallic conductivity captured the collective interest of the scientific community in such a way that none of the previous results had been able to [18, 23]. As a consequence, it was the collective work of Shirakawa, Heeger, and MacDiarmid that sparked the significant growth in the study of conjugated and conducting polymers that followed. In the process, the previously isolated efforts on either polyaniline or polypyrrole joined the growing work on polyacetylene to become the nucleus of a new field of electronic materials.

1.7.1 Synthetic Metals During the development of this new field, doped conjugated polymers became synonymous with the term synthetic metals, as more and more of these materials were exhibiting electrical conductivities in the metallic range [21, 22, 47, 48, 155]. This is aptly illustrated by Alan MacDiarmid’s Nobel lecture entitled “‘Synthetic Metals’: A Novel Role for Organic Polymers” [48]. The term predates conducting polymers, however, dating back to 1911 when it was first introduced by Herbert N. McCoy (1870–1945) at the University of Chicago, who had concluded [21, 22, 156]: ... that it is possible to prepare composite metallic substances, which may be termed synthetic metals, from constituent elements, some of which at least are nonmetallic. The first practical and stable examples of synthetic metals were various graphite intercalation compounds reported by Alfred R. Ubbelohde (1907–1988) beginning in 1951 [22]. During the early 1970s, however, organic charge-transfer salts, metal chain compounds, and poly(sulfur nitride) had all been discovered to exhibit metallic conductivity and thus added to the growing class of synthetic metals. Polyacetylene had been added by 1979 [157], and by 1991, MacDiarmid and Arthur Epstein had included doped polyparaphenylene, poly(phenylene vinylene), polypyrrole, polythiophene, and polyaniline in a review of conducting polymers as synthetic metals [155]. As research on this growing class of materials spanned a range of scientific disciplines and geography, it was desired to develop a setting to bring these interdisciplinary researchers together to share results and insights. This, then, resulted in the organization of a workshop held in Siofok, Hungary during the summer of 1976 [158]. This workshop became the seed that germinated into the organization of a longstanding international conference on these conducting materials, including the various theoretical and technological issues relating to their study. Given the official title of the International Conference on the Science and Technology of Synthetic Metals, this meeting is most commonly referred to by those in the field as just ICSM and still represents one of the primary international conferences for the field of conjugated and conducting polymers. In fact, it was at the second ICSM conference in New York City (ICSM ‘77) that Shirakawa, MacDiarmid, and Heeger’s November 1976 production of highly conductive materials from the Br2 and I2 doping of free-standing films of polyacetylene was first reported [22, 158].

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The conference was held annually from 1976 to 1982 and has been held biennially ever since. By 1999, half of the presentations covered conjugated and conducting polymers [159].

1.7.2 Dedicated Literature As the field of conducting polymers and synthetic metals continued to develop, a dedicated venue for the year-round dissemination of results was needed. Thus, a new journal was launched by Elsevier in October 1979, aptly titled Synthetic Metals (Figure 1.23) [22]. In the introduction of the first issue [160], Editor F. Lincoln Vogel of the University of Pennsylvania described this publication as ... a new international journal for the publication of research and engineering papers on graphite intercalation compounds, transition metal compounds, and quasi one-dimensional conducting polymers.

FIGURE 1.23  The cover of the first issue of Synthetic Metals, published October 1979. (Synthetic Metals, Vol 1, Issue 1, Copyright Elsevier and used with permission).

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Conjugated Polymers: Perspective, Theory, and New Materials

Initial Associate Editors of the journal included future Nobel laureate Alan J. Heeger, and Wayne L. Worrell (1937–2012), best known for his work in solid electrolytes [22]. In 1984, Heeger would also go on to take over duties from Founding Editor Vogel to become the journal’s Editor-in-Chief, until he ultimately relinquished those duties to Arthur Epstein in 2001. The initial Editorial Board also included future Nobel laureate Hideki Shirakawa, as well as Alfred Ubbelohde. The initial issue of the journal also contained Ubbelohde’s final published paper on intercalated graphite. To date, this is still the only journal dedicated primarily to organic conducting materials. The field then obtained one of its first critical reference texts in February of 1986 with the publication of the first edition of the Handbook of Conducting Polymers, published by Marcel Dekker, Inc. (Figure 1.24). Edited by Terje A. Skotheim of the Conducting Polymer Group at Brookhaven National Laboratory, this handbook consisted of two volumes totaling 1417 pages and arose from a desire to assemble knowledge of the various aspects of the chemistry and physics of conjugated polymers in a single reference source. As stated by Skotheim in the preface [161]: The purpose of these volumes is to provide an overview of the status and direction of this diverse field by presenting the ideas of leading research workers in the various subfield of conducting

FIGURE 1.24  The cover of the first edition of the Handbook of Conducting Polymers, Marcel Dekker, Inc., 1986.

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polymer research. Introductory material is included to make the subject matter usable for scientists of various backgrounds... It is my hope, therefore, that these volumes will provide a reference point for future developments. As the field continued to grow and develop, revised and updated editions of the Handbook followed, with the second edition published in November of 1997. The second edition also saw an expansion of the editorial team to include Ronald L. Elsenbaumer (University of Texas at Arlington) and John R. Reynolds (then at the University of Florida). A third edition was then published in January of 2007, with Skotheim and Reynolds returning as Editors [2]. This third edition was published by CRC Press, as Marcel Dekker, Inc. had been purchased by Taylor and Francis in 2003. The current 4th edition of the Handbook continues the previous tradition of an updated edition each decade, with the addition of Barry C. Thompson (University of Southern California) to the previous editorial team of Skotheim and Reynolds.

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70. Constantini, P., G. Belorgey, M. Jozefowicz, and R. Buvet. 1964. Préparation, propriétés acides et formation de complexes anioniques de polyanilines oligomères. C. R. Seances Acad. Sci. 258:6421–6424. 71. Jozefowicz, M., G. Belorgey, L. T. Yu, and R. Buvet. 1965. Oxydation et réduction de polyanilines oligomères. C. R. Seances Acad. Sci. 260:6367–6370. 72. Jozefowicz, M. and L. T. Yu. 1966. Relations entre propriétés chimiques et électrochimiques de semi-conducteurs macromoléculaires. Rev. Gen. Electr. 75:1008–1013. 73. Yu, L. T. and M. Jozefowicz. 1966. Conductivité et constitution chimique pe semi-conducteurs macromoléculaires. Rev. Gen. Electr. 75: 1014–1018. 74. De Surville, R., M. Jozefowicz, L. T. Yu, J. Perichon, and R. Buvet. 1968. Electrochemical chains using protolytic organic semiconductors. Electrochim. Acta 13:1451–1458. 75. Jozefowicz, M., L. T. Yu, J. Perichon, and R. Buvet. 1969. Proprietes Nouvelles des Polymeres Semiconducteurs. J. Polym. Sci. Part C: Polym. Symp. 22:1187–1195. 76. Angeli, A. 1915. Sopra il nero del pirrolo. Nota preliminare. Atti Accad. Naz. Lincei Cl. Sci. Fis. Mat. Nat. Rend. 24:3–6. 77. Angeli, A. and L. Alessandri. 1916. Sopra il nero pirrolo II. Nota. Gazz. Chim. Ital. 46(II):283–300. 78. Angeli, A. and G. Cusmano. 1917. Sopra i neri di nitropirrolo. Nota. Atti Accad. Naz. Lincei Cl. Sci. Fis. Mat. Nat. Rend. 26(I):273–278. 79. Angeli, A. 1918. Sopra i neri di pirrolo. Nota. Gazz. Chim. Ital. 48(II):21–25. 80. Angeli, A. and C. Lutri. 1920. Nuove ricerche sopra i neri di pirrolo. Nota. Atti Accad. Naz. Lincei Cl. Sci. Fis. Mat. Nat. Rend. 29(I):14–22. 81. Angeli, A. and A. Pieroni. 1918. Sopra un nuovo modo di formazione del nero di pirrolo. Nota. Atti Accad. Naz. Lincei Cl. Sci. Fis. Mat. Nat. Rend. 27(II):300–304. 82. Ciusa, R. 1921. Sulla scomposizione dello iodolo. Atti Accad. Naz. Lincei Cl. Sci. Fis. Mat. Nat. Rend. 30(II):468–469. 83. Ciusa, R. 1922. Sulle grafiti da pirrolo e da tiofene (Nota preliminare). Gazz. Chim. Ital. 52(II):130–131. 84. Ciusa, R. 1925. Su alcune sostanze analoghe alla grafite. Gazz. Chim. Ital. 55:385–389. 85. McNeill, R. and D. E. Weiss. 1959. A xanthene polymer with semiconducting properties. Aust. J. Chem. 12:643–656. 86. Weiss, D. E. and B. A. Bolto. 1965. Organic polymers that conduct electricity. In: Physics and Chemistry of the Organic Solid State, Vol. II. New York, NY: Interscience Publishers, pp 67–120. 87. Chierici, L. and G. P. Gardini. 1966. Structure of the product C8H10N2O of oxidation of pyrrole. Tetrahedron 22:53–56. 88. Bocchi, V., L. Chierici, and G. P. Gardini. 1967. Structure of the oxidation product of pyrrole. Tetrahedron 23:737–740. 89. Bocchi, V., L. Chierici, G. P. Gardini, and R. Mondelli. 1970. Pyrrole oxidation with hydrogen peroxide. Tetrahedron 26:4073–4082. 90. Dascola, G., D. C. Giori, V. Varacca, and L. Chierici. 1966. Rèsonance paramagnètique èlectronique des radicaux libres crèès lors de la formation des noirs d’oxypyrrol. Note. C. R. Seances Acad. Sci. Ser. C 267:433–435. 91. Dall’Olio, A., G. Dascola, V. Varacca, and V. Bocchi. 1968. Resonance paramagnètique èlectronique et conductiviè d’un noir d’oxypyrrol èlectrolytique. Note. C. R. Seances Acad. Sci. Ser. C 267:433–435. 92. Chierici, L., G. C. Artusi, and V. Bocchi. 1968. Sui neri di ossipirrolo. Ann. Chim. 58:903–913. 93. Diaz, A. F., K. K. Kanazawa, and G. P. Gardini. 1979. Electrochemical polymerization of pyrrole. J. Chem. Soc. Chem. Commun. 635–636. 94. Kanazawa, K. K., A. F. Diaz, W. D. Gill, P. M. Grant, G. B. Street, G. P. Gardini, and J. F. Kwak. 1980. Polypyrrole: An electrochemically synthesized conducting organic polymer. Synth. Met. 1:329–336.

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95. Kanazawa, K. K., A. F. Diaz, R. H. Geiss, W. D. Gill, J. F. Kwak, J. A. Logan, J. F. Rabolt, and G. B. Street. 1979. ‘Organic Metals’: Polypyrrole, a stable synthetic ‘Metallic’ polymer. J. Chem. Soc. Chem. Commun. 854–855. 96. Diaz, A. F., A. Martinez, K. K. Kanazawa, and M. Salmon. 1981. Electrochemistry of some substituted pyrroles. J. Electroanal. Chem. 130:181–187. 97. Rasmussen, S. C. 2017. Cuprene: A historical curiosity along the path to polyacetylene. Bull. Hist. Chem. 42:63–78. 98. Natta, G., P. Pino, and G. Mazzanti. 1955. Polimeri ad elevato peso molecolore degli idrocarburi acetilenici e procedimento per la loro preparozione. Italian Patent 530,753 (July 15, 1955). Chem. Abst. 1958, 52:15128b. 99. Natta, G., G. Mazzanti, and P. Pino. 1957. Hochpolymere von Acetylen-Kohlenwasserstoffen, erhalten mittels Organometall-Komplexen von Zwischenschalenelementen als Katalysatoren. Angew. Chem. 69:685–686. 100. Natta, G., G. Mazzanti, and P. Corradini. 1958. Polimerizzazione stereospecifica dell’acetilene. Atti Accad. Naz. Lincei Mem. Cl. Sci. Fis. Mat. Nat. 25:3–12. 101. Shirakawa, H. 2001. The discovery of polyacetylene film: The dawning of an era of conducting polymers (Nobel Lecture). Angew. Chem. Int. Ed. 40:2574–2580. 102. Shirakawa, H. 2002. The discovery of polyacetylene film. The dawning of an era of conducting polymers. Synth. Met. 125:3–10. 103. Hatano, M., S. Kanbara, and S. Okamoto. 1961. Paramagnetic and electric properties of polyacetylene. J. Polym. Sci. 51:S26–S29. 104. Kanbara, S., M. Hatano, and T. Hosoe. 1962. 遷移金属アセチルアセトナート-トリエチルア ルミニウム系によるアセチレンの重合 (Polymerization of acetylene by transition metal acetylacetonate-triethylaluminium system). J. Soc. Chem. Ind. Japan 65:720–723. 105. Shimamura, K., M. Hatano, S. Kanbara, and I. Nakada. 1967. Electrical conduction of poly-acetylene under high pressure. J. Phys. Soc. Japan 23:578–581. 106. Hatano, M. 1967. アセチレン重合体の構造と電気的性質 (Structures and electrical properties of acetylene polymers). Tanso 50:26–31. 107. Ikeda, S. and T. Akihiro. 1963. Syntheses of benzene-14C6 and benzene-2H6 using a Ziegler-catalyst. Radioisotopes 12:368–372. 108. Ikeda, S., T. Akihiro, and Y. Akira. 1964. Measurement of C2 component in acetylene-14C polymerization system by Ziegler catalyst by radio gas chromatography. Radioisotopes 13:415–417. 109. Ikeda, S. and T. Akihiro. 1966. On the mechanism of the cyclization reaction of acetylene polymerization. J. Polym. Sci. B Polym. Lett. Ed. 4:605–607. 110. Ikeda, S. 1967. チグラー触媒によるエチレンおよびアセチレン重合の立体化学 (Stereochemistry of ethylene and acetylene polymerization by Ziegler catalyst). J. Soc. Chem. Ind. Japan 70:1880–1886. 111. Shirakawa, H. 2001. Hideki Shirakawa. In: Les Prix Nobel. The Nobel Prizes 2000, ed. T. Frängsmyr. Stockholm: Nobel Foundation, pp 213–216. 112. Hall, N. 2003. Twenty-five years of conducting polymers. Chem. Commun. 1–4. 113. Shirakawa, H. 1996. Reflections on “Simultaneous Polymerization and Formation of Polyacetylene Film on the Surface of Concentrated Soluble Ziegler-Type Catalyst Solution,” by Takeo Ito, Hideki Shirakawa, and Sakuji Ikeda, J. Polym. Sci.: Polym. Chem. Ed., 12, 11 (1974). J. Polym. Sci. A Polym. Chem. 34:2529–2530. 114. Shirakawa, H. and S. Ikeda. 1971. Infrared spectra of poly(acetylene). Polym. J. 2:231–244. 115. Hargittai, I. 2011. Risking reputation: Conducting polymers. In: Drive and Curiosity: What Fuels the Passion for Science. Amherst, NY: Prometheus Books, pp 173–190. 116. MacDiarmid, A. G. 2001. Alan G. MacDiarmid. In: Les Prix Nobel. The Nobel Prizes 2000, ed. T. Frängsmyr. Stockholm: Nobel Foundation, pp 183–190. 117. Hargittai, B. and I. Hargittai. 2005. Alan J. Heeger. In: Candid Science V: Conversations with Famous Scientists. London: Imperial College Press, pp 411–427.

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118. Shirakawa, H., T. Ito, and S. Ikeda. 1973. Raman scattering and electronic spectra of poly(acetylene). Polym. J. 4:460–462. 119. Ito, T., H. Shirakawa, and S. Ikeda. 1974. Simultaneous polymerization and formation of polyactylene film on the surface of concentrated soluble Ziegler-type catalyst solution. J. Polym. Sci. Polym. Chem. Ed. 12:11–20. 120. Ito, T., H. Shirakawa, and S. Ikeda. 1975. Thermal cis-trans isomerization and decomposition of polyacetylene. J. Polym. Sci. Polym. Chem. Ed. 13:1943–1950. 121. Ito, T., H. Shirakawa, and S. Ikeda. 1976. ポリアセチレンのシス― トランス組成と固体構造 (Cistrans composition and solid structure of polyacetylene). Kobunshi Ronbunshu 33:339–345. 122. Shirakawa, H., T. Ito, and S. Ikeda. 1978. Electrical properties of polyacetylene with various cistrans compositions. Makromol. Chem. 179:1565–1573. 123. Berets, D. J. and D. S. Smith. 1968. Electrical properties of linear polyacetylene. Trans. Faraday Soc. 68:823–828. 124. MacDiarmid, A. G., C. M. Mikulski, P. J. Russo, M. S. Saran, A. F. Garito, and A. J. Heeger. 1975. Synthesis and structure of the polymeric metal, (SN)x, and its precursor, S2N2. J. Chem. Soc. Chem. Commun. 476–477. 125. Mikulski, C. M., P. J. Russo, M. S. Saran, A. G. MacDiarmid, A. F. Garito, and A. J. Heeger. 1975. Synthesis and structure of metallic polymeric sulfur nitride, (SN)x, and its precursor, disulfur dinitride, S2N2. J. Am. Chem. Soc. 97:6358–6363. 126. Chiang, C. K., M. J. Cohen, A. F. Garito, A. J. Heeger, C. M. Mikulski, and A. G. MacDiarmid. 1976. Electrical conductivity of (SN)x. Solid State Commun. 18:1451–1455. 127. Chiang, C. K., M. J. Cohen, D. L. Peebles, A. J. Heeger, M. Akhtar, J. Kleppinger, A. G. MacDiarmid, J. Milliken, and M. J. Moran. 1977. Transport and optical properties of polythiazyl bromides: (SNBr0.4)x. Solid State Commun. 23:607–612. 128. Heeger, A. J. 2001. Semiconducting and metallic polymers: The fourth generation of polymeric materials (Nobel Lecture). Angew. Chem. Int. Ed. 40:2591–2611. 129. Meisel, S. L., G. C. Johnson, and H. D. Hartough. 1950. Polymerization of thiophene and alkylthiophenes. J. Am. Chem. Soc. 72:1910–1912. 130. Armour, M., A. G. Davies, J. Upadhyay, and A. Wassermann. 1967. Colored electrically conducting polymers from furan, pyrrole, and thiophene. J. Polym. Sci. Part A-1: Polym. Chem. 5:1527–1538. 131. Curtis, R. F., D. M. Jones, and W. A. Thomas. 1971. The ‘Trimer’ and ‘Pentamer’ from the polymerisation of thiophen by polyphosphoric acid. J. Chem. Soc. C 234–238. 132. Louvar, J. J. 1971. Polymerization of Heterocyclic Compounds. Patent No. US 3,574,072. 133. Yamamoto, T., K. Sanechika, and A. Yamamoto. 1980. Preparation of thermostable and electricconducting poly(2,5-thienylene). J. Polym. Sci. Polym. Lett. Ed. 18:9–12. 134. Yamamoto, T., K. Sanechika, and A. Yamamoto. 1981. Preparation of poly(2,4-thienylene) and comparison of its optical and electrical properties with those of poly(2,5-thienylene). Chem. Lett. 10:1079–1082. 135. Yamamoto, T. and K. Sanechika. 1982. Preparation and properties of p-conjugated poly(3-methyl2,5-thienylene). Chem. Ind. 9:301–302. 136. Yamamoto, T., K. Sanechika, and A. Yamamoto. 1983. Preparation and characterization of poly(thienyene)s. Bull. Chem. Soc. Jpn. 56:1497–1502. 137. Yamamoto, T., K. Sanechika, and A. Yamamoto. 1983. Formation of adducts of poly(thienylene)s with electron acceptors and electric conductivities of the adducts. Bull. Chem. Soc. Jpn. 56:1503–1507. 138. Yamamoto, T., Z. Masanobu, and A. Yamamoto. 1984. Li│LiI│iodine galvanic cells using iodinepoly(2,5-thienylene) adducts as active materials of positive electrodes. Chem. Lett. 13:1577–1580. 139. Yamamoto, T., Z. Masanobu, and A. Yamamoto. 1985. Secondary cells using poly(2,5-thienylene)s and poly(2,5-pyrrolylene)s as materials of positive electrodes. Zn│ZnI2│I2 secondary cell. Chem. Lett. 14:563–566.

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140. Lin, J. W.-P. and L. P. Dudek. 1980. Synthesis and properties of poly(2,5-thienylene). J. Polym. Sci. Polym. Chem. Ed. 18:2869–2873. 141. Afanas’ev, V. L., I. B. Nazarova, and M. L. Khidekel. 1980. Polythiophene – An electrically conductive organic compound. Izvestiia Akademii nauk SSSR. Seriia khimicheskaia (7):1687–1688. 142. Diaz, A. F., J. Crowley, J. Bargon, G. P. Gardini, and J. B. Torrance. 1981. Electrooxidation of aromatic oligomers and conducting polymers. J. Electroanal. Chem. 121:355–361. 143. Wattman, R. J., J. Bargon, and A. F. Diaz. 1983. Electrochemical studies of some conducting polythiophene films. J. Phys. Chem. 87:1459–1463. 144. Waltman, R. J., A. F. Diaz, and J. Bargon. 1984. Electroactive properties of polyaromatic molecules. J. Electrochem. Soc. 131:740–744. 145. Waltman, R. J., A. F. Diaz, and J. Bargon. 1984. Electroactive properties of polyaromatic molecules. J. Electrochem. Soc. 131:1453–1456. 146. Tourillon, G. and F. Garnier. 1982. New electrochemically generated organic conducting polymers. J. Electroanal. Chem. 135:173–178. 147. Tourillon, G. and F. Garnier. 1983. Effect of dopant on the physicochemical and electrical properties of organic conducting polymers. J. Phys. Chem. 87:2289–2292. 148. Tourillon, G. and F. Garnier. 1983. Stability of conducting polythiophene and derivatives. J. Electrochem. Soc. 130:2042–2044. 149. Kaneto, K., K. Yoshino, and Y. Inuishi. 1982. Electrical properties of conducting polymer, polythiophene prepared by electrochemical polymerization. Japan J. Appl. Phys. 21:L567–L568. 150. Kaneto, K., Y. Kohno, K. Yoshino, and Y. Inuishi. 1983. Electrochemical preparation of a metallic polythiophene film. J. Chem. Soc. Chem. Commun. 382–383. 151. Kaneto, K., K. Yoshino, and Y. Inuishi. 1983. Electrical and optical properties of polythiophene prepared by electrochemical polymerization. Solid State Commun. 46:389–391. 152. Koßmehl, G. and G. Chatzitheodorou. 1981. Electrical conductivity of poly(2,5-thiophenediyl)AsF5-complexes. Makromol. Chem. Rapid Commun. 2:551–555. 153. Koßmehl, G. and G. Chatzitheodorou. 1982. Electrical conductive AsF5-complexes of poly (2,5-thiophenediyl). Mol. Cryst. Liq. Cryst. 83:291–296. 154. Koßmehl, G. and G. Chatzitheodorou. 1983. Synthesis and electrical conductivities of poly(2,5-thiophenediyl) salts by the action of nitronium or nitrosonium salts. Makromol. Chem. Rapid Commun. 4:639–643. 155. MacDiarmid, A. G. and A. J. Epstein. 1991. ‘Synthetic Metals’: A novel role for organic polymers. Makromol. Chem. Macromol. Symp. 51:11–28. 156. McCoy, H. N. 1911. Synthetic metals from non-metallic elements. Science 34:138–142. 157. MacDiarmid, A. G. 1979. Metallic and semiconducting materials derived from non-metallic elements. Microstruct. Sci. Eng. Technol. 13.1–13.8. 158. Reynolds, J. R. and A. J. Epstein. 2000. ICSM 2000: Over twenty-five years of synthetic metals. Adv. Mater. 12:1565–1570. 159. Bernier, P. 1999. Preface. Synth. Met. 101:xxv. 160. Vogel, F. L. 1979. Editorial introduction. Synth. Met. 1:1. 161. Skotheim, T. A., ed. 1986. Handbook of Conducting Polymers, New York: Marcel Dekker, Inc., pp. iii–iv.

2 Recent Advances in the Computational Characterization of π-Conjugated Organic Semiconductors 2.1 Introduction.........................................................................................38 2.2 Density Functional Theory for Organic Electronics......................38

2.3

Jean-Luc Brédas, Xiankai Chen, Thomas Körzdörfer, Hong Li, Chad Risko, Sean M. Ryno, and Tonghui Wang* *

This chapter is dedicated to Professor Alan J. Heeger, whose seminal contributions to the field of organic electronics, guidance, and friendship have been a constant source of inspiration for over thirty years, and to the memory of Professor Robert J. Silbey, our dear friend and mentor, who has molded us into the scientists we have become and whose legacy will continue to shape future generations of scientists.

2.4

2.5

2.6

The Electronic-Structure Method of Choice for Organic Electronic Materials • A Brief Introduction to DFT and TD-DFT • Challenges in DFT Applications and Recent Advances in Functional Development

Noncovalent Interactions and Polarization in the Condensed Phase.................................................................................47 Noncovalent Interactions: Solid-State Packing, Miscibility, and Processing • Polarization and Site Energies in the Bulk and at Interfaces: Impact on Charged-State Characteristics

A Theoretical Description of Organic Emitters for LightEmitting Diodes Exploiting Thermally Assisted Delayed Fluorescence����������������������������������������������������������������������������������������56 Theoretical Description of Reverse Intersystem Crossing • Relationships of the Spin-Orbit Couplings with the Excitation Characteristics • Role of Non-Adiabatic Coupling in the Reverse Intersystem Crossing Process • Novel Molecular-Design Strategies for TADF Emitters.

Molecular Dynamics Description of Organic-Organic Interfaces and Polymer Pure Phases.................................................66 Interfaces Between Layers of Small Molecules: Interfacial Mixing • π-Conjugated Polymer Pure Phases: Main-Chain Conformation and Inter-Chain Packing • Polymer-Fullerene Packing and Interfaces in the Mixed Regions

Characterization of the Interfaces between an Organic Layer and a Metal or Conducting Oxide Surface........................... 74 Description of the Change in Surface Workfunction upon Deposition of an Organic Layer • Brief Description of the Computational Methodology • Surface Defects • Charge-Transfer Characteristics for Donor/Acceptor Molecules Physisorbed on Metal-Oxide Surfaces • Characterization of the Binding Modes of the Surface Modifiers

Acknowledgments...........................................................................................84 References.........................................................................................................84 37

38

Conjugated Polymers: Perspective, Theory, and New Materials

2.1 Introduction The past ten years have witnessed a flurry of advances in the field of organic electronics, as many of the chapters in this book will highlight. This has clearly been the case on the applications side with: organic light-emitting diodes (OLEDs), now a multi-billion-dollar industry, taking a significant part of the high-end market for smart-phone and television displays and entering the solid-state lighting market; organic solar cells (OSCs) on the verge of breaking the 15% power conversion efficiency (PCE) mark; organic field-effect transistors (OFETs) demonstrating charge-carrier mobilities over 10 cm2/(V·s); and the growth of flexible and wearable organic electronics and organic bioelectronics. In parallel, major progress has also taken place on the theoretical and computational side, which constitutes the focus of the present chapter. Density Functional Theory (DFT) has now firmly emerged as the electronic-structure method of choice for organic electronic materials. In Section 2.2, we present a brief overview of recent advances in the DFT description of organic electronic systems, as well as a discussion of the most pressing challenges and of important caveats. Section 2.3 is devoted to the progress achieved in the theoretical understanding of how the chemical nature and molecular architecture of π-conjugated chromophores determine: (i) the intermolecular noncovalent interactions that impact the molecular packing arrangements in thin films or crystals; and (ii) polarization phenomena in the solid state, both of which govern the material electronic and optical properties. Commercially available OLEDs are currently based on second-generation emitters consisting of coordination complexes of heavy metal atoms, such as Ir or Pt, with organic ligands. Since a few years ago, much attention has been paid to the design of a new generation of purely organic emitters whose efficiency relies on thermally assisted delayed fluorescence (TADF). Section 2.4 introduces the basic theoretical concepts behind the design of such TADF emitters. A ubiquitous component of organic electronic devices is the presence of multiple interfaces. In OSCs, the active layers are usually bulk heterojunctions, i.e. blends of an electron-donor component and an electron-acceptor component; in this instance, organic-organic interfaces between donors and acceptors appear at the meso-scale or nano-scale. Section 2.5 introduces the molecular dynamics approaches that have been recently refined to characterize the local morphology in these blends. Finally, in Section 2.6, we highlight the theoretical understanding that we are reaching of organicinorganic hybrid interfaces, such as those formed between organic surface modifiers and transparent conducting oxides used as charge-injection or -collection electrodes.

2.2 Density Functional Theory for Organic Electronics 2.2.1 The Electronic-Structure Method of Choice for Organic Electronic Materials Basic material research in the field of organic electronics covers an extensive and diverse spectrum of materials, ranging from small molecules such as TCNE (tetracyanoethylene) or TTF (tetrathiafulvalene), fullerenes such as PC61BM (phenyl-C61-butric acid methyl ester), to polymers and polymer mixtures such as PEDOT:PSS (poly(3,4-ethylenedioxythiophene) polystyrene sulfonate). Devices such as OLEDs and OSCs typically feature several layers of metal-to-semiconductor and semiconductorto-semiconductor interfaces, which can be organic-organic or organic-inorganic, and sometimes crystalline vs. amorphous in nature. A coherent computational characterization of π-conjugated organic semiconductors therefore requires a theoretical approach that allows the assessment on equal footing of the structural, electronic, and optical properties of this large variety of materials and respective interfaces. Over the past ten years, DFT1,2 has become unrivaled among all electronic-structure methods when it comes to the computational characterization of π-conjugated organic semiconductors. This is

Recent Advances in Computational Characterization

39

due to its excellent accuracy-to-numerical-cost ratio, its ease of implementation and application, its ability to describe both finite (small molecules, oligomers) and periodic (polymers, crystals) systems and their interfaces, and because it allows not only characterization of the ground-state and excited-state electronic and structural properties but also molecular dynamics simulations at the ab initio level. The interest in DFT calculations has been growing steadily over the past 20 years. According to a Web of Science topic search,3 more than 27,000 papers that dealt with DFT calculations were published in 2016 alone. As shown in Figure 2.1, almost half of these publications used one of the two most prominent exchange-correlation (XC) functionals, that is, the semi-local PBE functional,4 commonly used in physics for periodic calculations on solids, and the global hybrid functional B3LYP,5–8 frequently used in chemistry for the computational description of molecules. Generally speaking, the choice of the XC functional determines both the numerical costs as well as the accuracy of a DFT calculation. Despite their success, standard semi-local and global hybrid functionals also display a number of issues, the most prominent ones being the lack of nonlocal correlation effects such as dispersion interactions9–13 and the appearance of electron self-interaction errors.14–16 Depending on the context, the latter are often also referred to as localization/delocalization errors.17–19 A vast amount of publications discuss these problems in detail and a variety of solutions have been suggested;20–27 a detailed account of all these aspects is, therefore, out of the scope of this chapter. Instead, we aim here to provide insights into some of the most prominent challenges encountered in the description of organic electronic materials in recent years,28 and to discuss some of the recent advances in DFT functional development that we have found particularly useful in our own work on π-conjugated organic semiconductors. However, before we do so, as many readers of this book might not be very familiar with computational chemistry, it appears useful to provide a brief introduction into the basic concepts of DFT and time-dependent (TD)-DFT. More detailed reviews can be found in Refs.2,29–32

FIGURE 2.1  The number of DFT citations has exploded over the last 20 years. PBE and B3LYP represent the number of citations of References 4 and 5, respectively, in each year. Other represents papers that do not cite either of the two but have DFT or TD-DFT as topics. All numbers are estimations based on a Web of Science3 search.

40

Conjugated Polymers: Perspective, Theory, and New Materials

2.2.2 A Brief Introduction to DFT and TD-DFT First, a distinction has to be made between electronic ground-state and excited-state calculations. In ground-state DFT, one solves a set of one-particle Schrödinger equations of the form:33,34



é 1 ù ê - 2 D + v H éën ùû ( r ) + vion éën ùû ( r ) + vˆ xc ú ji (r ) = i ji (r ) (2.1) ë û n(r ) =

N

å f j (r ) i

i

2

(2.2)

i =1

N



å f = M (2.3) i

i =1

Here, n ( r ) denotes the electron density at point r in space and M the total number of electrons. Solving Equation (2.1) leads to a set of (occupied and unoccupied) orbitals, and eigenvalues, εi. In a way similar to Hartree–Fock (HF) theory, the classical electron-electron interactions and electron-ion interactions are incorporated into the local potentials v H [n ]( r ) and vion [n ]( r ), respectively. Note that in Equation (2.1) we have chosen the notation used in Generalized Kohn-Sham (GKS) theory,34 which means that the XC effects are integrated into the operator vˆ xc; this operator has a general expression combining a local XC contribution with a non-local, HF-like exchange contribution:

vˆ xc = b vˆ HF + (1 − b ) v x [n ]( r ) + v c [n ]( r ) (2.4)

In standard Kohn-Sham (KS) theory, 33 b is set to 0 and the XC operator reduces to a local potential. This is the approach followed in the original XC functional, the local-density approximation (LDA), 33 as well as in generalized gradient approximations (GGAs) such as the PBE functional4 and meta-GGAs.35 Using a global admixture of HF exchange, as expressed in Equation (2.4), leads to so-called global hybrid functionals. An example of the latter is the B3LYP functional, where b = 0.2. Importantly, there also exist other ways of mixing DFT and HF exchange, some of which will be discussed in more detail below. At this stage, however, it is key to realize that this mixing of the local XC functional with non-local HF-like exchange leads to a paradigm shift, with the underlying theory changing from standard KS33 to GKS34 theory. Also, it can be sometimes necessary to include a nonlocal correlation contribution, that is, to mix the local v c [n ]( r ) correlation potential with non-local correlation contributions, such as those obtained from second-order Møller-Plesset perturbation theory (MP2) or random phase approximation (RPA) calculations; such functionals are referred to as double-hybrid functionals.36,37 Ground-state DFT calculations can be used to calculate a number of important parameters, including molecular geometries, vibrational frequencies, electronic band-structures, electronic densities, and density differences, or dipole moments. However, it is by construction an electronic ground-state theory; thus, the calculation of charged excitation energies such as ionization energies, electron affinities (EAs), and charge-transfer (CT) gaps, which has increasingly become a central aspect of the computational characterization of π-conjugated materials, can be a major challenge, as will be further discussed below. The computation of (neutral) optical excitations requires a TD-DFT treatment.38–40 Intuitively, this can be understood based on the picture that, upon optical excitation, the dipole moment starts to oscillate at a frequency that corresponds to the excitation energy, that is, the electron density becomes TD.

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Since an explicit treatment of this time-dependence is numerically demanding, one usually relies on linear-response theory. Here, the objective is to assess, in the linear regime, the system’s response due to a perturbation from a TD electric field, which is assumed to be small compared to the electron-nuclear interactions. The application of this linear-response formalism in TD-DFT leads to Casida’s equations38 (where spin indices and occupation numbers are omitted for clarity): A  B* 



B   Xω  1 =ω A*   Yω  0

0  X ω  (2.5) 1   Yω 

Here, in the general notation for hybrid functionals in the spirit of Equation (2.4):

Aia , jb = δij δab ( εa − εi ) + (ia|jb ) − bHF (ij|ab ) + (1 − bHF ) (ia f xc jb ) (2.6)



Bia , jb = (ia|bj ) − bHF (ib|aj ) + (1 − bHF ) (ia f xc bj ) (2.7)

where, the two-electron integrals are given in Mulliken notation:

(ia|jb ) : = ∫ ∫ ϕi* (r )ϕa (r )

(ia f

xc

jb ) : =

1 ϕ j ( r ′) ϕb* ( r ′) drdr ′ (2.8) r − r′

∫ ∫ ϕ (r)ϕ (r) f * i

a

xc

(r , r ′) ϕ j (r ′) ϕb* (r ′) drdr ′ (2.9)

with f xc representing the XC kernel, which is derived from the XC functional E xc [n ] via:

f xc ( r , r ′) =

δv xc ( r ) δE xc = (2.10) δn ( r ) δn ( r ′ ) δn ( r ′ )

The ϕi terms denote the occupied (indices i and j) and unoccupied (indices a and b) orbitals and the εi terms the corresponding eigenvalues, obtained by a ground-state DFT calculation. Solving Casida’s equations provides the optical (neutral) excitation frequencies ω as the eigenvalues of the Casida matrix. From the corresponding eigenvectors, it is possible to derive the oscillator strength for each of the excitations. Just as in the case of ground-state DFT, the quality of the results obtained from TD-DFT largely depends on the choice of the XC functional, which enters Casida’s equations both directly and indirectly: directly through the XC kernel of Equation (2.10) and indirectly through the orbitals and eigenvalues obtained from the ground-state DFT calculation. A frequently used variant of linear-response TD-DFT is provided by the Tamm–Dancoff approximation (TDA),41 which is obtained by setting the B-matrix in Equation (2.5) to zero. We note that the TDA can be considered to be the TD-DFT equivalent of the (post-HF) configuration interaction (CI) singles approach, with the main difference that the ground- and excited-state Slater determinants are now set up from (G)KS orbitals and not from HF orbitals. As we will discuss in more detail at the end of this section, the TDA can be particularly useful for the calculation of triplet states, in particular in situations prone to triplet instabilities.42,43 It is important to realize that TD-DFT per se does not describe excitations as simple single-particle transitions between individual molecular orbitals. Instead, each individual TD-DFT excitation is characterized as a transition between the single-determinant DFT ground-state and a specific linear combination of singly excited Slater determinants. For example, the lowest optical excitation in organic π-conjugated molecules is often mainly characterized by a transition from the highest occupied molecular orbital (HOMO) to the lowest unoccupied molecular orbital (LUMO); however, it also typically carries contributions from other single-particle transitions, e.g. from HOMO-n to LUMO or HOMO

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to LUMO+n. As a consequence, simply considering the HOMO and LUMO wavefunctions can often be misleading when attempting to characterize the nature of an excitation. A much sounder alternative is to plot the so-called natural transition orbitals (NTOs), which are obtained by a singular value decomposition of the transition density matrix obtained for the TD-DFT excitation of interest. More simply, NTOs are a graphical representation of the linear combination of the occupied (hole-NTO) and unoccupied (electron-NTO) orbitals that contribute to a particular optical excitation. In Section 2.4, we will employ such an NTO analysis for the characterization of excitations in TADF emitters for OLEDs.

2.2.3 Challenges in DFT Applications and Recent Advances in Functional Development As mentioned earlier, PBE (for periodic solids) and B3LYP (for molecules) have been the most prominent XC functionals over the last 20 years, which is primarily due to their usually very good accuracy-to-numerical-cost ratio, wide availability, and ease of implementation. This also holds true in the field of organic π-conjugated systems, where these and other semi-local and global hybrid functionals have long been the work horses for computational chemists and physicists alike. In recent years, the ease of application of modern quantum chemistry codes has further opened up the field to scientists with limited background in quantum chemistry. Hence, the popularity of these XC functionals and the ever-growing number of researchers using these approximations make it critical to point out the most important drawbacks and errors when applying these standard functionals to organic electronic materials. 2.2.3.1 Condensed Phases and the Problem of Dispersion Corrections in DFT A very well-known error associated with standard XC functionals is the lack of non-local correlation effects such as dispersion (i.e. induced dipole–induced dipole) interactions. This error originates from the semi-local nature of the most commonly used correlation functionals, which is inherently not suitable for describing long-range, non-local correlation. In the field of organic electronics, this problem becomes particularly important in the computational assessment of polymer chain conformations, condensed-phase systems, and molecule/substrate interactions, where the evaluation of the structural, vibronic, and electronic properties strongly depends on a correct description of dispersion interactions.36,44–59 The problem of missing correlation effects in standard DFT functionals can be addressed in various ways.9,11,60 One obvious way, though, numerically, a rather expensive solution, is to directly include nonlocal correlation in an XC functional derived, for example, from a non-local kernel61 or via RPA.62,63 A simpler yet often sufficiently accurate and numerically efficient approach is to add an empirical dispersion correction on top of a converged DFT calculation.9,60 This straightforward approach, referred to as DFT-D, has been consistently improved over the years. For example, the latest dispersion correction suggested by Grimme et al., DFT-D3,64 includes atom-pairwise specific dispersion coefficients and cutoff radii that are computed from first principles, and the different local chemical environments of individual atoms are considered when determining the empirical van der Waals coefficients. In highly parameterized XC functionals, that is, functionals with a large number of empirical parameters that are typically fitted to experimental and theoretical test sets, it can be especially beneficial to re-optimize the parameters after the empirical dispersion correction has been added. This is the idea behind the ωB97X-D functional,65 which we have found particularly useful in our recent work on organic π-conjugated materials, since it combines the accuracy of highly parameterized functionals with the benefits of using a long-range corrected (LRC) hybrid functional (see below) and an empirical dispersion correction in the spirit of Grimme’s D236 approach. Since such pairwise-based dispersion corrections can overestimate the stability of large complexes, especially those with considerable overlap of the π molecular orbitals,66,67 further improvements can

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come from incorporating many-body effects of the relevant chemical environment into the pair-potentials through a self-consistent procedure.67–69 The inclusion of non-local correlations based on the adiabatic fluctuation dissipation theorem,70–72 via connecting DFT with many-body theory, has allowed the development of new functionals within the random phase approximation.73–80 Non-local components have also been directly included in DFT functionals (e.g. vdW-DFT) in order to improve the description of noncovalent interaction strengths.81–84 While further improvements of dispersion corrections in DFT remain a very active field of research, current approximations have reached a level of accuracy, efficiency, ease of application, and availability that allows researchers interested in organic electronic materials to describe reliably the structural and vibronic properties of organic materials and molecules on surfaces. For more detailed reviews of firstprinciple models for dispersion interactions, the reader is referred to Refs.11,60,85 2.2.3.2 Self-Interaction Errors and Tuned Long-Range Corrected Hybrid Functionals One of the oldest and still prominent problems of standard XC functionals such as LDA, semi-local functionals such as PBE, and standard global hybrid functionals such as B3LYP is the so-called electron self-interaction error.14 While its origin can be easily understood, it has proven to be challenging to solve. The self-interaction problem arises from the fact that LDA, and with it all semi-local functionals such as GGAs and meta-GGAs, derive their approximation for exchange and correlation interactions from the exact solution for the homogenous electron gas. As a consequence, these functionals typically work well in situations with many electrons and a slowly varying electronic density. In contrast, the worst-case scenario for such approximations corresponds to a single, localized electron. The key aspect here is that, while in HF theory the exchange contribution exactly cancels the classical Coulomb interaction of one-electron densities, this is not the case for semi-local XC functionals. As a consequence, for one-electron systems, a spurious, repulsive electron-electron interaction remains, which corresponds to the electron self-interaction error. Among the most important consequences of the appearance of this repulsive self-interaction error are: (i) the destabilization of bound electrons (for instance, HOMO energies are too high in energy); (ii) a general tendency for the electron density to overly delocalize; (iii) a dramatic underestimation of CT excitation energies; and (iv) an overestimation of conjugation effects, which results in an overestimation of rotational barriers86 and an underestimation of the degree of bond-length alternation.87 We note that the tendency of standard DFT approaches to significantly underestimate the charge-transport gaps is often attributed to the self-interaction error. However, the issue of how to interpret HOMO–LUMO gaps in DFT is, in fact, an intricate problem, which we will address separately below. While the self-interaction error can be quantified and corrected easily in the case of a single electron, the issue becomes much harder to solve in many-electron systems. The reason is that electron motions are correlated; thus, we cannot distinguish individual electrons and, consequently, we cannot separate proper electron-electron interactions from electronic self-interactions. One way to circumvent this problem is to focus on the delocalization error. This error can be formally quantified by coupling the system of interest to a bath of electrons and allowing an exchange of a fraction of an electron with that bath (see Figure 2.2).17,18 It can be demonstrated that, in an exact theory, the total energy of the system must evolve linearly with the fractional particle number (red curve).88 Semi-local and global hybrid functionals, however, spuriously stabilize situations with fractional charges, leading to a concave behavior of the total energy (blue curve). In contrast, the HF approach spuriously destabilizes fractional charges, such as to favor situations with integer charges (green curve). In other words, while HF has the tendency to localize charges and electron densities, semi-local and hybrid functionals tend to overly delocalize electron densities. This localization/delocalization error is a footprint of the electron self-interaction error in many-electron systems. Figure 2.2 also suggests a way to correct for the localization/delocalization error, namely by mixing HF exchange and semi-local XC functionals such that a straight-line behavior can be obtained. A mean

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FIGURE 2.2  Illustration of the localization/delocalization error. Total energy as a function of a fractional number of electrons when going from a neutral molecule (N) to its cation (N-1), as calculated by standard XC functionals and HF theory in comparison to the exact result. Fractional charges are spuriously destabilized in HF theory and spuriously stabilized in standard DFT.

to achieve this goal is to rely on non-empirically tuned LRC hybrid functionals.89,90 The basic idea91 of LRC hybrid functionals is to partition the 1/r Coulomb operator (with r being the electron-electron distance) into a short-range and a long-range part via the standard error function:

1 erf ( ωr ) 1 − erf ( ωr ) = + (2.11) r r r

The error function varies smoothly from erf(0) = 0 to erf(∞) = 1, and the range separation is determined by the parameter ω. By treating short-range and long-range electron-electron interactions on a different footing, the range-separation scheme allows one to incorporate, on the one hand, semi-local or standard hybrid DFT within the short-range part of the Coulomb operator and, on the other hand, full HF exchange plus semi-local correlation in the long-range component. The main benefit of LRC hybrid functionals is that they restore the correct 1/r asymptotic behavior of the XC functional, which is lacking in semi-local and global hybrid functionals. The remaining problem is the choice of the range-separation parameter, for which two options exist. One is to empirically fit the range-separation parameter once and for all to a given test set of molecules. This approach has the advantage of being size extensive and no further fitting for specific systems of interest is required. In the case of organic π-conjugated systems, however, it has been demonstrated that the optimal range-separation parameter, i.e. the one that leads to the correct straight-line behavior in Figure 2.2 and therefore minimizes the localization/delocalization error, can vary significantly with the size and degree of conjugation of the system at hand.92 Consequently, the range-separation parameter has often to be non-empirically tuned in such a way that the localization/delocalization error is minimized. This tuning can be achieved by choosing the range-separation parameter that minimizes the difference between the HOMO eigenvalue and the computed ionization potential (IP):

∆ IP ( ω ) = −εωHOMO − ( E gs ( ω, N ) − E gs ( ω, N − 1) ) (2.12)

This IP-tuning procedure89,90 has been shown to improve the description of properties related to the IP and the fundamental gap for a range of systems. We will demonstrate several applications of this methodology throughout this chapter.

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2.2.3.3 Charged Excitation Energies and the Physical Interpretation of Gaps in DFT The charged excitation energies of materials such as the IPs and EAs are among the key factors that determine the suitability of π-conjugated materials for use in organic electronic devices (for instance, in the case of OLEDs, they directly impact the hole- and electron-injection barriers from the electrodes). Being by construction a ground-state theory, the reliable prediction of IPs and EAs is a challenge for DFT. When one is only interested in the first IP and/or EA of a single molecule, the ∆SCF approach offers a simple solution. Here, the ground-state energies of the neutral molecule and its respective cation and anion are calculated and the first IP/EA as obtained as total-energy differences. In a recent benchmark study for a test set of 24 small-molecule organic acceptors,93–95 it was shown that, using the ∆SCF approach, the PBE functional yields mean absolute errors of 0.67 eV and 0.21 eV for IPs and EAs, respectively. The results can be further improved by using a global hybrid functional such as B3LYP, which yields mean absolute errors of 0.37 eV and 0.20 eV for IPs and EAs, respectively. Whether such results can be considered to be sufficiently accurate depends on the type of applications of interest. However, as mentioned in the Introduction, situations are often encountered in organic electronics where we are interested in the charge transfer between two molecules and/or polymer chain segments, be it in the context of donor-acceptor interfaces, small-gap polymers, or dopant-introduced semiconductor thin films. Here, the key is the ability to predict the level alignments, the extent of charge transfer across an interface, the interface dipoles, and the energies of the CT and charge-separated excited states. Since the ∆SCF approach is not applicable in these instances, alternative approaches are required. An accurate method for the evaluation of charged excitation energies is many-body perturbation theory in the GW approximation.96,97 However, due to large basis set requirements, these calculations are numerically much more expensive than DFT, even when carried out in a non-self-consistent way at the G0W0 level. As a result, when studying large π-conjugated molecules or interfaces, researchers often take the simple approach of physically interpreting the HOMO and LUMO eigenvalues and the corresponding (G)KS gap, for example, via a projected density-of-states approach. This approach, however, can be highly misleading. Indeed, it is important to realize first and foremost that, in contrast to HF theory, Koopman’s Theorem98 does not hold for DFT. In exact (G)KS theory, there is in fact only one eigenvalue that carries a physical meaning. According to the IP-Theorem,88,99 given the exact XC functional, the HOMO eigenvalue equals the exact vertical IP. However, there exists no similar theorem for any of the other eigenvalues; in particular, the LUMO eigenvalue does not equal the vertical EA. This is further illustrated in Figure 2.3, which highlights what we refer to as the energy-gap dilemma in DFT.100 We recall that, from an experimental standpoint, the fundamental gap, ∆Efund, is the difference between IP and EA of the material (this gap is also referred to as the transport gap). The optical gap, ∆Eopt, is the energy of the first optical excitation, which differs from the fundamental gap by the exciton binding energy, Eb. In π-conjugated systems, Eb is on the order of several tenths of 1 eV or larger. The KS gap, ∆EKS, i.e. the difference between the HOMO and LUMO energies in exact KS DFT, equals neither ∆Efund nor ∆Eopt. The difference, ∆ xc, between the fundamental gap and the KS gap is called the derivative discontinuity.88 This name can be understood from Figure 2.2, where ∆ xc manifests as the discontinuity of the derivative of the total energy (red line) at integer occupation. By introducing HF exchange into the XC functional, the underlying methodology is changed from KS to GKS, thereby changing the eigenvalues. In exact GKS theory, the LUMO eigenvalue and, therefore, the HOMO-LUMO gap increases with the amount of non-local HF-like exchange in the XC functional. Consequently, there typically exists an amount of HF exchange for which the GKS gap, ΔEGKS, can equal the optical and/or the fundamental gap. For B3LYP (20% HF), for example, it is often found empirically that the GKS gap is very close to the optical gap for many organic π-conjugated molecules. However, the optimal amount of HF exchange for the GKS gap to equal the fundamental gap, which is the one relevant for chargeseparation processes, is different for different systems; the reason is that the exciton binding energy and the derivative discontinuity are system-specific.

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FIGURE 2.3  The band-gap dilemma. The gaps of exact KS and GKS theories do not generally agree with the experimental optical or fundamental gaps.

The consequence is that the optimal amount of HF exchange needs to be determined independently for each system. This can be achieved by the non-empirical IP-tuning procedure introduced in Equation (2.12), since it guarantees the approximate functional to fulfill the IP-Theorem and, consequently, gives a physical meaning to the corresponding eigenvalues. If the IP-tuning is applied to global hybrid functionals in the spirit of Equation (2.4), one typically obtains b values in the range 0.7–0.9. However, such large amounts of HF exchange lead to an unbalanced functional and, consequently, to inaccurate results for many structural and electronic properties. Therefore, it is more appropriate to rely on LRC hybrid functionals, in which the amount of HF exchange is kept low for short-range interactions, which govern chemical binding, and goes to full HF exchange for long-range interactions, which govern CT and charged excitation energies. In the benchmark study mentioned above,94 non-empirically tuned LRC hybrid functionals have been demonstrated to yield HOMO and LUMO eigenvalues that, in comparison to the benchmark, yield mean absolute errors of 0.23 eV and 0.16 eV for IPs and EAs, respectively. Thus, these functionals not only yield HOMO/LUMO eigenvalues and GKS gaps that are physically interpretable, but they are also more accurate in predicting charged excitation energies than non-selfconsistent G0W0 approaches based on a semi-local or standard global hybrid functional starting point.95 Importantly, in the solid state, the fundamental gap is significantly reduced compared to individual molecules or polymer chains, due to polarization effects from the surrounding medium, as will be detailed in Section 2.3. In the context of non-empirically tuned LRC hybrid functionals, this gap renormalization can be accounted for by considering the dielectric constant of the medium. In practice, this can be achieved either via an explicit DFT calculation, in which the LRC hybrid functional is adapted to take into account the screening of the Coulomb operator,101 or, in the context of a much simpler and therefore also more approximate ansatz,102 by carrying out the non-empirical tuning in a polarizable environment modeled by a continuum solvation model. 2.2.3.4 Optical Excitation Energies, Charge-Transfer Excitations, and Triplet States Just as for ground-state DFT, the accuracy of linear-response TD-DFT calculations strongly depends on the choice of the XC functional.103–106 While modern global hybrid functionals typically yield sufficiently

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accurate optical excitation energies for organic π-conjugated molecules at comparatively low computational costs, several issues that can significantly hamper an accurate description of organic electronic materials still remain. For example, the commonly used adiabatic approximation in TD-DFT only allows the description of single excitations; thus, excited states with a significant double-excitation character, such as the 2 1Ag state in polyenes, cannot be properly characterized.107,108 In addition, commonly used semi-local and global hybrid functionals have failed in the description of Rydberg,109 CT,110–112 and triplet valence113,114 excited states. A number of solutions to these issues have been proposed, among them the recently developed local hybrid functionals, which have shown promising results but, to date, remain not widely available.113 A different approach, which we have found particularly useful in our recent work on the computational characterization of organic π-conjugated materials, is related to the non-empirical tuning of LRC hybrid functionals we introduced above.89,90 By incorporating full HF exchange for long-range electron-electron interactions and providing HOMO and LUMO energies that approximate IPs and EAs, this class of functionals is particularly useful when it comes to the computational assessment of CT excitations—for example, in donor-acceptor complexes—and optical excitation energies that carry a significant CT character—for example, in low gap materials such as donor-acceptor polymers. Also, LRC hybrid functionals manage to maintain the accuracy of modern global hybrid functionals for singlet valence excitations, since these are governed by short-range interactions. On the other hand, the high amount of exact exchange included in long-range interactions can exacerbate a problem that is well-known in TD HF theory, that is, the triplet instability problem.42,43 Triplet instability arises from the fact that HF theory (and XC functionals with a large amount of HF exchange) overestimates the stability of triplet states.115 As a consequence, such functionals often provide for an unbalanced treatment of singlet and triplet excitations. This constitutes a severe problem for the computational characterization of materials of interest in the case, for instance, of singlet fission or of TADF. Hence, to ensure a balanced description of CT, valence singlet, and triplet excited states in organic π-conjugated materials, a most useful option is to exploit a combination of LRC hybrid functionals with non-empirical tuning and the TDA approximation to TD-DFT.43 Indeed, by removing the so-called de-excitations from Casida’s equations, TDA has been demonstrated to prevent the triplet instability problem.42,115 This combination of methodologies has been applied to the case of TADF, as will be highlighted in Section 2.4; it has also been exploited in the DFT calculations, referred to in Section 2.5, of CT singlet and triplet excited states present at organic-organic donoracceptor interfaces.

2.3 Noncovalent Interactions and Polarization in the Condensed Phase In this section, we are interested in describing what has been learned recently regarding how the chemical composition and molecular architecture of organic semiconductors direct noncovalent interactions and polarization in the condensed phase. As described in other chapters of the Handbook, advances in molecular design, material processing—including innovative solution formulations, printing techniques, and post-processing practices—and device architectures have enabled significant advancements in device performance. From the standpoint of organic semiconductor design, there tends to be a general focus on tuning molecular-scale electronic and optical properties. However, even subtle chemical changes imparted to π-conjugated chromophores to regulate these characteristics can have profound, and often unintended, impact on the solid-state packing arrangements that effectively govern the eventual material electronic and optical properties. Hence, there has been a clear need to develop a better understanding of how molecular chemistry impacts the noncovalent interactions and the preference for specific packing arrangements in crystals or thin films of organic semiconductors.116–134

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Also, the electronic-state energies related to reduced or oxidized molecules, which are relevant in the context of charge-carrier transport or to the CT complexes formed at materials interfaces (see Section 2.5), are governed by the solid-state environment. The impact of the surrounding medium, described by the solid-state polarization, is determined by the individual molecular electrostatic moments, the strengths of the intermolecular electrostatic interactions, and the natures of the charged/ excited states. The polarizable environment is challenging to describe due to the long-range character of the interactions among electrostatic moments, and requires theoretical methods that can span multiple length scales in an efficient manner. Here, we present recent advances in theoretical approaches to uncover the physicochemical relationships that are essential to describing solid-state noncovalent interactions and polarization effects. We begin this section with an overview of what is meant by noncovalent interactions, and how electronicstructure and molecular dynamics techniques can be used to determine their impact on materials packing. We then follow a similar course for solid-state polarization, showcasing progress in the description of this multiscale phenomenon. Looking forward to the next advances in organic electronics, our discussion here can provide inspiration to the synthetic materials chemists and process engineers and enable novel design paradigms by explicitly correlating noncovalent interactions and polarization phenomena to device performance.

2.3.1 Noncovalent Interactions: Solid-State Packing, Miscibility, and Processing We begin with a clarification of nomenclature for the purpose of this discussion. One often finds reference to ambiguous “π-π interactions” as being significant attractive forces that stabilize solid-state molecular arrangements in organic semiconductors; these designations are then used, for instance, to discuss the possibility to create materials with perfectly co-facially packed π-conjugated backbones. However, there actually exist clear physical definitions for noncovalent interactions. When appropriately considered, these classifications provide detailed chemical insight that can be used to understand why molecular and polymer-based materials take on certain solid-state packing arrangements, and can be implemented in new materials design guidelines. While there remains intense debate as to what π-π interactions are and what role they play, as a function of chemical composition and molecular architecture, in determining molecular packings,56,135–146 it is in fact preferable to discuss π-π interactions in the context of electronic interactions, i.e. the overlap of π wavefunctions on neighboring molecules that determine the extent of intermolecular electronic couplings. To consider more precise definitions of noncovalent interactions,147 we begin with the force that gives matter volume: exchange-repulsion. A consequence of the Pauli exclusion principle, which negates two electrons (fermions) from occupying the same quantum state and, therefore, the same region of space,148,149 the repulsion between electrons on adjacent molecules limits the degree of molecular contact. Exchange-repulsion has significant implications for the maximum electronic couplings that can be achieved among molecules (or polymer chain segments) in the solid state. As shown for two co-facially stacked tetracene molecules (Figure 2.4), both the intermolecular exchange-repulsion and electronic coupling increase exponentially with decreasing intermolecular distance.128,132 However, the increase for the exchange-repulsion is significantly steeper at shorter distances—and in particular at intermolecular distances found in organic semiconductors—providing a physical limit to the ability to create molecular materials with tight and full co-facial packing arrangements and strong intermolecular electronic couplings. We note that chemical considerations can be invoked to overcome the physical limitations set forward by exchange-repulsion. For instance, one can modify the chemical composition to alter the dispersion interactions, which constitute a first attractive component. Dispersion, also referred to as London dispersion or London forces, is often the leading contributor to intermolecular stabilization in π-conjugated molecular and polymeric materials. Interactions arising from instantaneous fluctuations

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FIGURE 2.4  For a co-facially stacked tetracene dimer as a function of intermolecular separation: intermolecular electronic couplings (tH, closed, circles), determined from the absolute value of the HOMO:HOMO overlap at the B3LYP/cc-pVDZ level of theory, and exchange-repulsion energies (open, circles), determined at the SAPT0/juncc-pVDZ level of theory. The inset represents the tetracene dimer. (Adapted from Sutton C., et al. The interplay between intramolecular and intermolecular interactions determines the planarization of its tetracene core in the solid state. Journal of the American Chemical Society 2015, 137(27), 8775–8782.)

of the electron density (i.e. induced dipoles) give rise to the R −6 dependence of the dispersion energy that forms the attractive part of an interatomic potential. Strong dispersive interactions are usually present amongst the π-conjugated backbones, the alkyl side-chains often appended to make these systems more soluble, and their combinations.133 A second attractive force is due to electrostatic interactions, which result from interactions among the permanent multipole moments centered on each molecule—e.g. dipole-dipole, dipole-quadrupole, quadrupole-quadrupole, and the like. Such electrostatic interactions are the determining contributions to the electrostatic energy at large intermolecular distances, typically over 4 Å. Stronger electrostatic interactions turn on at short distances (less than 4 Å), where considerable overlap can be expected among the molecular charge densities, which would lead to a large exchange-repulsion interaction. However, the overlapping charge density on one molecule experiences significant attraction from the nuclei of the other molecule, and vice versa; this large attractive contribution is referred to as charge penetration.131,150,151 Given that intermolecular distances in many organic semiconductors are less than 4 Å, contributions from charge penetration can be substantial. We stress that analyses based on more simple approaches to account for electrostatics, such as atom-centered charges, electrostatic potential plots, or even distributed multipoles, are unable to grasp such charge penetration effects and can, thus, incorrectly predict too repulsive electrostatics. The final attractive force, induction, is the energetic stabilization due to the electronic relaxation of the charge distribution on one molecule—i.e. formation of induced electrostatic moments—in response to the presence of another molecule and its permanent multipoles. Contributions from induction are commonly small in magnitude, though they can become significant when considering charged molecules (i.e. interactions with monopole moments). Supermolecular approaches are often implemented to compute the interaction energy ( Eint ) between two systems. For instance, in the case of a complex formed by two molecules (denoted “molecule 1” and “molecule 2”), each with their own energy as isolated molecules ( E1 and E2 , respectively), an indirect

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determination of the interaction energy comes from the difference of the complex energy ( Ecomp ) compared with the isolated molecular energies at fixed geometry:

Eint = Ecomp − ( E1 + E2 ) (2.13)

For solids, it is the cohesive energy that is evaluated to determine the stabilization of a molecule/polymer in the solid-state environment. For a crystalline system, for instance, the cohesive energy ( Ecoh ) is defined as:

Ecoh =

E solid − Eisolated (2.14) N

where Eisolated is the energy of the isolated (gas phase) molecule, E solid is the energy of the crystalline (solid) unit cell, and N is the total number of molecules in the unit cell. By these definitions, the interaction and cohesive energies are negative for stable interactions. For the computational practitioner, there have been tremendous strides in methods development to explore the strength and nature of noncovalent interactions, though the task does remain difficult.52,54,57,59,130,147,152–156 Indeed, the application of the supermolecular approach requires very accurate methods capable of calculating extremely small differences in total energies, as the total energy of a system is several orders of magnitude larger than noncovalent interaction energies. Moreover, noncovalent interactions arise from the correlated motions of electrons.157 Hence, high-level wavefunction methods, for example the so-called gold-standard CCSD(T) method—the coupled-cluster method with single, double, and perturbative triple electron excitations—is needed in combination with a large basis set to capture these electron correlation effects.137,158–162 These methods impose a substantial computational cost and are therefore limited to small molecular systems.163,164 Perturbation theories, for example the MP2 approach, can be used to compute the strength of noncovalent interactions. While MP2 tends to overestimate interaction energies, especially those arising from dispersive π-interactions,165 error cancellations can negate some of these issues, allowing MP2 methods to provide acceptable descriptions for the strength of intermolecular interactions in aromatic systems.166,167 Moreover, care has to be taken regarding basis set inconsistency in the context of supermolecular calculations, which can lead to basis set superposition errors (BSSE).168,169 DFT methods, of course, remain a workhorse for the investigation of the strength of noncovalent interactions, both in terms of the supermolecular approach and the determination of crystal cohesive energies. Significant improvements in the determination of these interactions have been endowed by the addition of corrections that account for long-range dispersion interactions; see Section 2.2 of this chapter for a discussion of some of these corrections. Such supermolecular and cohesive energy approaches, however, do not provide for a decomposition of the interaction energy into the physical components described above for noncovalent interactions. Given that noncovalent interactions are small, decomposition schemes rely on perturbation theory to determine the interaction energy from a sum of various energy terms. Symmetry-adapted perturbation theory (SAPT)151,161,170–174 is one such perturbative approach that has been widely adopted. The SAPT method decomposes the total interaction into contributions arising from exchange-repulsion [Eexch], dispersion [Edisp], electrostatic moments [Eelec], and induction [Eind], with the total interaction energy being given by the sum of the individual components:

Eint = Eexch + Eelec + Eind + Edisp (2.15)

The so-called SAPT0 approach (where the “0” indicates that intra-monomer electron correlation is neglected) has been shown to give accurate stacking energies for a wide array of noncovalent

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systems, especially when combined with an appropriately sized basis set (such as jun-cc-pVDZ).159,171–173 Importantly, this level of theory can provide rigorous insight into the nature of intermolecular interactions in π-conjugated systems.128,129,131–134,172–173,175 Before delving into materials-specific examples, we also note that molecular dynamics (MD) approaches have been important in determining miscibility parameters in the condensed phase, both in solution and the solid state. The Hildebrand solubility parameter (δT )176 is the square-root of the cohesive energy density (CED), where the CED is defined as the difference between the enthalpy of vaporization (∆H) and the product of the universal gas constant times temperature (RT) divided by the molar volume (V):

δT = CED =

∆H − RT = δ2D − δ2H − δ2P (2.16) V

Hansen177 further separated the Hildebrand parameter into three intermolecular contributions: dispersive interactions (δ D ), Coulombic or dipole-dipole interactions (δ P), and hydrogen-bonding interactions (δ H ). The energies to compute cohesive energy densities and Hildebrand and Hansen solubility parameters can be readily extracted from equilibrated MD simulations, which has proven to be an important simulation method to relate the strengths of noncovalent interactions to the properties of materials systems. SAPT-derived noncovalent intermolecular interactions can be related to these widely used Hildebrand and Hansen solubility parameters. Indeed, connecting SAPT interaction energies to the solubility parameters can be done in principle by evaluating solute-solute, solutesolvent, and solvent-solvent interactions, and including the average number of neighbors from MD simulations or NMR (nuclear magnetic resonance) spectroscopy.178,179 New force fields based on SAPT interaction energies are also under increasing development, providing the potential for larger-scale simulations that directly take into account such quantum mechanically determined results to evaluate solubility.180,181 We now turn to a few examples as to how the computational approaches outlined above can be used to explore materials interactions in the condensed phase. As fruit-fly materials for the organic electronics community, noncovalent interactions in the oligoacene series have been widely vetted.128,131,145,151,170,182–194 For instance, SAPT0/jun-cc-pVDZ (a level of theory that we will now simply refer to as “SAPT0”) was used to explore the potential energy surfaces for co-facially stacked acenes of varying length as a function intermolecular displacement.131 As one may expect, the total SAPT0 interaction energies fall off quickly as the intermolecular separation along the stacking axis moves beyond ca. 3.5 Å (Figure 2.5); the exchange-repulsion energy decreases rapidly, with the electrostatic, induction, and dispersion energies following suit. Dispersion is the dominant attractive interaction at intermediate intermolecular separations (i.e. greater than 4.5 Å), while at even larger separations (i.e. larger than 7 Å) electrostatic interactions dominate. These variations in the length scales of the stabilizing interactions are due to the 1/R 5 distance dependence of the quadrupole-quadrupole interactions compared to the approximately 1/R6 distance dependence of dispersion. In these dimers, slipping one molecule along the long axis affords two-dimensional potential energy surfaces that reveal large fluctuations of the total interaction energies.131 The potential energy surfaces display peaks and valleys whose numbers correspond with the number of fused rings in the acene (i.e. two minima are found for naphthalene). Interestingly, the interaction energy minima correspond to staggered arrangements that are reminiscent of those found in triisopropylsilylethynyl (TIPS)-pentacene and rubrene (5,6,11,12-tetraphenyl-tetracene).132,133,195,196 While it is generally accepted that the TIPS groups in TIPS-pentacene and the (nearly) orthogonal phenyl groups in rubrene in large part prevent the acene backbones from adopting their commonly found herringbone packing configurations, the noncovalent backbone interactions fine tune the packing; they lead to staggered, parallel-displaced arrangements that turn out to provide strong electronic couplings and favorable charge-carrier transport characteristics.

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FIGURE 2.5  SAPT0-based analysis for a parallel-displaced pentacene dimer at varied intermolecular distances. The long-axis displacement of the top pentacene molecule is 1 Å along the long axis. The SAPT0 energies are represented by: total interaction energy (inverted triangles), exchange-repulsion (squares), dispersion (triangles), electrostatics (circles), and induction (diamonds); the total HF energy is given by the cross-hatches. (Adapted from Ryno S., et al. Noncovalent interactions and impact of charge penetration effects in linear oligoacene dimers and single crystals. Chemistry of Materials 2016, 28(11), 3990–4000.)

Not only are the acenes prototypical materials for organic semiconductor applications, their structures can be considered as platforms to which synthetic chemists can add functionality through substitution. Heteroatom substitution has been widely used to modify the molecular redox and optical characteristics; however, these substitutions also have great impact on the preferred molecular packing configurations and the resulting materials properties.128–130,134 For instance, in a series of azapentacenes,128 SAPT0 results for co-facial configurations at an intermolecular distance of 3.5 Å reveal that nitrogen substitution into the acene backbone reduces the exchange-repulsion. Arising from smaller wavefunction overlap (and, in turn, smaller electronic couplings) with increasing nitrogen content, the smaller exchange-repulsion energies are evidence for contracted electron density in the N-heteropentacenes. The dispersion and electrostatic energies are also generally reduced with nitrogen substitution. Hence, even though the molecules in these series are iso-electronic and the same intermolecular separations are considered, the nature and strength of the intermolecular interactions are found to vary in significant ways with such rather subtle changes in chemistry. In the case of thiophene derivatives such as thienoacenes, positional disorder of the sulfur atoms, specifically in trialkylsilylethynyl-substituted anthradithiophenes (ADT), has been proposed as an explanation to the variations in charge-carrier transport characteristics of the molecular materials derived from these systems.134,197 As thin films of the molecular organic semiconductor materials are generally disordered, pathways of consistent molecular alignments that mitigate the impact of charge trapping sites (i.e. defect sites within the material that lead to electronic disorder) are necessary to provide efficient charge-carrier transport conduits through the active layer. SAPT0 and DFT calculations of anti-benzodithiophene (BDT) model structures (Figure 2.6) reveal important differences in the noncovalent interactions and subsequent intermolecular electronic couplings as a function of sulfur positional disorder.134 In general, it is found that contacts where sulfur atoms are in close proximity are less energetically favorable than those where the sulfur atoms are further away from each other, due to large exchange-repulsion energies; on the other

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FIGURE 2.6  SAPT0 total and decomposition energies for two anti BDT conformations as a function of displacement along the long molecular axis. The molecules comprising anti BDT dimers are separated by 3.5 Å. The closed symbols represent the dimer complex labeled 1, while the open symbols represent the dimer complex labeled 2. The SAPT0 energies are represented by: total interaction energy (inverted triangles), exchange-repulsion (squares), dispersion (triangles), electrostatics (circles), and induction (diamonds). (Adapted from Thorley K.J. and Risko C. On the impact of isomer structure and packing disorder in thienoacene organic semiconductors. J. Mater. Chem. C 2016, 4(18), 4040–4048.)

hand, the electronic couplings tend to be larger when direct sulfur contacts are present. These effects have implications for the description of the electronic properties of related π-conjugated polymers as well. For instance, combined theoretical modeling and experimental approaches were used to develop models of the (semi)crystalline polymer poly(2,5-bis (3-tetradecylthiophene-2-yl)thieno[3,2-b]thiophene) [PBTTT], which pointed to a slipped-stack packing arrangement of the PBTTT backbone, which limited sulfursulfur contacts.198,199 These results contrasted a previous cofacial theoretical model200 and modified the understanding of the PBTTT electronic characteristics. So far, we have focused on the noncovalent interactions among the π-conjugated backbones. However, information pertaining to how alkyl side-chains impact molecular packing can also be derived from electronic-structure calculations. For instance, for the class of trialkylsilylethynyl-substituted acenes, quantum-chemical evaluations describe how the slightly larger electron density contained within the volume of the TIPS moiety of TIPS-pentacene when compared to its triethylsilylethynyl (TES) pentacene counterpart is responsible at least partly for the differences in preferred packing (brickwork for TIPS-pentacene and slipped-stack for TES-pentacene).133 The in silico exploration of polymorphs suggested that TES-pentacene, if appropriately processed, could be developed into a material with improved charge-carrier transport characteristics when compared to its native form. In the case of the polymer poly(benzo[1,2-b:4,5-b′]dithiophene–thieno[3,4-c]pyrrole-4,6-dione), PBDTTPD, combined electronic-structure calculations and MD simulations revealed at the molecular scale how the nature of the alkyl side-chains (linear or branched) on the polymer direct the close contacts with PC61BM.201–203 As will be further discussed in Section 2.5, the simulations indicate that linear alkyl side-chains on the PBDTTPD backbone tend to extend away

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from the polymer backbone, while bulkier branched side-chains remain in closer proximity. This, in turn, impacts the fullerene preferential locations around the polymer backbone.203 This latter example brings us to the use of MD simulations to detail how noncovalent interactions impact materials properties. MD simulations of fullerenes—including PC61BM, phenyl-C71-butyric acid methyl ester (PC71BM), and indene-C60-mono adduct (ICMA)—as well as of the polymer poly(3-hexylthiophene) [P3HT], have illustrated the wide range over which noncovalent interactions determine the materials characteristics.204–208 Hildebrand and Hansen miscibility parameters for these materials show very good comparability with available experimental data.209,210 Furthermore, the simulated trends in thermal transitions, which provide a means to measure the strength of condensed-phase noncovalent interactions, are also well reproduced; for instance, the appearance of a cold crystallization event in PC71BM was identified through MD simulations, providing molecular details as to the nature of the solid-solid phase transformation.205 These characteristics, developed from molecular-scale interactions, have in turn provided insight into variations in the cohesive energies for polymer:fullerene blends as a function of the substituent on the fullerene, and the amount of mechanical stress that can be applied prior to film fracture. Moreover, simulations with chlorinated solvents and high-boiling-point solvent additives, which make direct connections to the variations in the fullerene:fullerene, fullerene:solvent, fullerene:additive, and solvent:additive interactions, have illustrated the effects that these additives have on the formation of fullerene aggregates in casting solutions.208 Though we have highlighted only a few examples, the above discussion underlines how aspects pertaining to molecular-scale noncovalent interactions can have profound impact on understanding materials properties. Undoubtedly, this developing knowledge starts allowing for a move away from trial-and-error approaches and lays down the foundation for an a priori design of new molecular and polymeric materials and the requisite processing protocols.

2.3.2 Polarization and Site Energies in the Bulk and at Interfaces: Impact on Charged-State Characteristics We now consider how solid-state intermolecular interactions impact processes involving charge carriers.211,212 Here, the focus is on the site and polarization energies. The site energy corresponds to the energy of the charged-state at each individual site (i.e. molecule or polymer chain segment) in the solid state. The polarization energy measures the stabilization of a charge-carrier due to the nature of the solid-state environment when compared to the gas phase. Following the definition of Lyons,213,214 the electronic polarization energy that decreases the ionization energy [increases the electron affinity] is given by:

P+ = IE solid − IE gas  P− = EAsolid − EAgas  (2.17)

where P+ [ P− ] is the polarization energy in the presence of a positive [negative] charge, IE solid [ EAsolid ] is the ionization energy [electron affinity] in the solid state, typically measured by ultraviolet photoelectron spectroscopy (UPS) [inverse photoelectron spectroscopy, IPES], and IE gas  EAgas  is the gas-phase ionization energy [electron affinity], measured by gas-phase UPS [electron detachment] techniques. Over the last few decades, there has been tremendous effort to develop models that could describe solid-state polarization fully and accurately. The long-range electrostatic interactions, charge (de)localization and band dispersion, molecular orientation, material disorder, surface electrostatic interactions, and lattice and nuclear relaxations, all contribute to the polarization energy.211,214–218 Taking into account each of these characteristics is not without challenge. Since these methodologies have been reviewed elsewhere, here we will only make a cursory overview to highlight some of the methodological differences.211,212 We note that most models localize the charge of interest on a single molecule in order to simplify the theoretical considerations.212 A variety of classical, quantum, and hybrid methods thereof have been developed to build ever more complex models that take into account greater details over larger-length scales. For instance, among

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the simpler representations is the classical electrostatic cavity model developed by Born,219,220 wherein the polarization energy is determined for a point charge in a spherical cavity embedded in an isotropic polarizable medium. Microelectrostatic models, which can explicitly include both permanent (static) and induced (dynamic) electrostatic interactions, offer opportunities to treat large-scale, complex systems.214,217,221–226 Though implementations vary, these microelectrostatic models determine the polarization energy via an evaluation of the static and dynamic intermolecular electrostatic interactions in the presence and absence of a charge-carrier (on a single molecule) and taking the difference in the total energies of these two situations. In a step beyond microelectrostatic models, charge redistribution (and charge response) models divide molecules into sections that are individually polarizable, which allow charges to reorganize (again, confined to a single molecule) in response to an electric field.227–234 Finally, significant effort has been devoted to the development of hybrid classical/quantum and fully quantum mechanical models to evaluate the polarization energy.101,235–241 The quantum mechanical methods range from semiempirical HF methods to DFT, thus allowing for great flexibility in the complexity of the systems being investigated. We now review a few examples as to how these approaches are used to examine the solid-state polarization effects. We begin with localized charges in bulk crystalline systems, whose conceptual simplicity has made them among the most widely studied; again, the oligoacenes and their derivatives have often served as test beds for theoretical investigations.217,222,242–248 For instance, even though the isolated molecular electronic structures of pentacene and TIPS-pentacene are similar, as determined by electronicstructure calculations and UPS measurements,246–248 their different solid-state packing arrangements (herringbone for pentacene and brickwork for TIPS-pentacene) lead to very different noncovalent interactions and, in turn, polarization energies.248 Notably, while the quadrupole-quadrupole interactions are stabilizing in pentacene, they are destabilizing in TIPS-pentacene. As a result, there occurs a large reduction in the calculated polarization energy of TIPS-pentacene (P+, 0.59 eV; P−, 0.69 eV) compared to pentacene (P+, 1.02 eV; P-, 0.79 eV), with the trends for P+ corresponding well with experiment.246,247 Additionally, the polarization asymmetry (i.e. the difference between P+ and P−) is larger in pentacene than TIPS-pentacene. Similarly, when comparing tetracene and rubrene, there exist substantial electrostatic and polarization differences, which again result from the variations in packing.249 Though more computationally demanding, the impact of charge delocalization over several molecules has also been considered in theoretical models of bulk systems. Here again, fullerenes and oligoacenes have served as systems of interest.250–253 In the case of oligoacenes, polarization and delocalization effects have recently been shown to be additive.254 In crystalline lattices, defects present in the bulk or grain boundaries are sources for static energetic disorder, i.e. variations in the energy of a charged molecule in the solid state as a function of differences in its environment.223,243 Microelectrostatic models have revealed how charge (monopole)-quadrupole interactions at sites along pentacene grain boundaries impede charge-carrier transport by acting as either intrinsic barriers or trapping sites. Amorphous systems, given their fluctuating distributions of molecular orientations, are obviously expected to have larger degrees of static energetic disorder. These effects, for instance, have been shown to be important in the fullerenes,207 where both the chemical nature of the substituent and fullerene size and shape can lead to major changes in the static disorder of the amorphous phase. In addition to static (positional) energetic disorder, dynamic energetic disorder arises from thermally induced molecular motions (i.e. vibrational or translational motions).217 Dynamic disorder is the sole source for site-energy variations in perfectly ordered and pure crystals made of equivalent molecules.212 Sorting out contributions to site energy disorder from static and dynamic components requires that such molecular motions be taken into account, often requiring to rely on MD techniques. As shown in the cases of both acenes and fullerenes,207,255 dynamic disorder can lead to significant broadening of the range of ionization energies or EAs and have considerable impact on charge transport. Recent experimental and theoretical efforts, for instance, have shown that controlling dynamic disorder through appropriate substitution of alkyl side-chains along the rod-like structures of acenes and related

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molecules can significantly limit the impact of solid-state molecular motions, adding a new feature to the molecular design paradigm.256,257 With the exception of grain boundaries, which can be viewed as interfaces within a homogeneous material, we have so far neglected the impact of interfaces. Heterogeneous interfaces, of course, are ubiquitous in organic electronic devices, whether they are formed at the point of contact between the organic semiconductor and dielectric, or electrodes (or electrode-boundary layers), or within active layers formed from multicomponent, phase-separated materials. The methods outlined above have been implemented and modified to explore the impact of such heterogeneous interfaces on the processes of charge-carrier transport, photoinduced charge transfer, and charge separation and recombination.212,252,258,259 Heterogeneous interfaces with insulators, dielectrics, and metals, for instance, are expected to have a broad impact on the polarization effects and site energies relevant for charge-carrier transport. Even the interface with vacuum has been calculated in the case of tetracene to lead to a decrease of 0.1 eV for the polarization of holes when compared to the bulk.245,249 Similar “free-standing” models of organic semiconductors are often sufficient to treat interfaces with a non-polar insulator surface.260 However, when the organic semiconductor rests on a polar substrate (e.g. dielectric or self-assembled monolayer),261–263 the electrostatic effects from the underlying material need to be explicitly accounted. Finally, the physico-chemical characteristics of metals (or conducting oxides) require special attention when interfaced with organic semiconductors. In particular, the large metal polarizability and the potential for vacuum-level shifts are of major interest, as will be detailed in Section 2.6.212,264–269 In the case of OSCs, theoretical models have been critical in understanding how the polarization effects and site energies change as a function of distance from organic-organic heterojunction interfaces and, in turn, impact the charge transfer, separation, or recombination processes. Clearly, the electrostatic interactions at organic-organic heterojunction interfaces can vary significantly.224,225,259,270–277 Here as well, acenes and fullerenes have served as model systems to explore organic-organic heterojunctions.225,275 Investigations of the pentacene(001) or pentacene(0–11) interfaces with C60 reveal that each donor and acceptor site (molecule) at the interface resides in a distinct electrostatic environment, which is a function of charge-quadrupole and direction-dependent induced-dipole interactions. In addition, the dynamic fluctuations in the electrostatic environment due to the thermally induced molecular motions need to be considered (Figure 2.7).225,275 A number of groups have explored the development and impact of band bending, the creation of electrostatic environments that can favor or negate charge separation, and the significant energetic disorder in IE and EA that can result in large barriers to charge separation.224,225,272,275–276,278 Though conceptually simple model interfaces have generally been the focus of these theoretical investigations, the results allow the emergence of a better picture as to how intermolecular multipole interactions, polarizability, and static and dynamic disorder at the many interfaces in organic semiconductor devices impact the materials characteristics that control meso-scale device function.

2.4 A Theoretical Description of Organic Emitters for Light-Emitting Diodes Exploiting Thermally Assisted Delayed Fluorescence Much attention has been paid, both in academia and industry, to OLEDs since the pioneering works of Tang and Vanslyke279 on small-molecule emitters, and of Friend and co-workers on conjugated polymer emitters.280 However, for these first-generation organic emitters, electrical excitation generally results in the formation of 25% singlet excitons and 75% triplet excitons, due to spin statistics. The radiative decay from the lowest singlet excited state (S1) to the ground state (S0) contributes to a prompt fluorescence emission, while the phosphorescence emission from the lowest triplet excited state (T1) to the ground state is nominally spin forbidden given the negligible spin-orbit coupling (SOC) in purely organic molecules. As a result, the triplet excitons decay non-radiatively, which implies a 75% energy loss. Assuming a light outcoupling efficiency of 20%, this means that the maximum external quantum efficiency (EQE) is limited to 5%.

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FIGURE 2.7  Polarization energy variation at selected sites near the pentacene(001) (circles): C60 (squares) interface as a function of time. The polarization energies are computed in 0.5 ps increments. (Adapted from Ryno S.M., et al. Polarization energies at organic-organic interfaces: impact on the charge separation barrier at donor-acceptor interfaces in organic solar cells. ACS Applied Materials & Interfaces 2016, 8(24), 15524–15534.)

These energy losses are the reason why, since the 1990s, much effort has been devoted to utilize triplet excitons in order to break the 5% efficiency limit of the first-generation OLED devices. In 1998, Forrest, Thompson, and co-workers started developing organometallic coordination complexes of heavy transition metals (e.g. iridium or platinum) as the emitters in the active layers of OLEDs.281,282 In these secondgeneration OLEDs, the triplet excitons can be harvested through efficient intersystem crossing (ISC) from the T1 to S0 state, enabled by the strong spin-orbit coupling induced by the heavy metal atoms. Internal quantum efficiencies (IQE) can reach up to 100% and, through active layer and outcoupling optimizations, EQEs over 30% have recently been demonstrated.283,284 However, the high cost of phosphorescent OLEDs, due to expensive transition metals, significantly limits their practical industrial applications. While these second-generation emitters are those exploited in the OLEDs currently on the market, there is strong impetus to find new ways of exploiting purely organic emitters in order to: (i) lower the cost associated with heavy metals; and (ii) increase the ability to fine tune the electronic and optical properties via the synthetic flexibility associated with organic compounds. In 2012, Adachi and co-workers proposed a promising approach to harvest triplet excitons in organic emitters by promoting reverse ISC (RISC) from T1 to S1 via simple thermal activation.285 This process, illustrated in Figure 2.8, gives rise to thermally activated delayed fluorescence (TADF) in metal-free organic conjugated compounds (see Figure 2.8). Since then, a large number of experimental and theoretical investigations have been devoted to these purely organic TADF third-generation OLED emitters.286–289 Impressive photo-physical properties and device performances have been reported with, in some instances, IQE reaching nearly 100% and EQE as high as 41.5%.290 However, there remain challenging issues, such as how to best balance the need for a small single (S1) – triplet (T1) energy gap with the need for a large S1 to S0 oscillator strength, the limited stability of the blue TADF emitters, and relatively low efficiencies of red and near-infrared TADF emitters. It is thus desirable to develop a better understanding of the basic TADF mechanisms. Our objective in this section is then to give a brief introduction to some recent advances in the theoretical understanding of TADF emitters.

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FIGURE 2.8  Schematic diagram of electroluminescence processes in OLED devices utilizing TADF emitters. h+/e−: hole/electron generated under electrical excitation; S0: ground state; S1: the first singlet excited state; T1: the first triplet excited state; ∆E: singlet-triplet energy gap; SOC: spin-orbit coupling; F: fluorescence, DF: delayed fluorescence, RISC: reverse intersystem crossing, E: energy.

Figure 2.8 can be used to describe the basic theoretical aspects that apply to TADF emitters. The RISC process depends on the nature of the lowest singlet and triplet excited states, which determine both the singlet-triplet energy gap, ∆EST, and the extent of SOC; as we will describe, this can also be strongly impacted by molecular vibrations and vibronic coupling effects. Additionally, the efficiency of both prompt and delayed fluorescence processes depends on the strength of the radiative decay to the ground state. Figure 2.9 displays some representative TADF emitters. As can be observed, these emitters usually combine electron-rich donor and electron-poor acceptor moieties, meaning that the lowest excited states are expected to have strong CT character. As explained in Section 2.2, this is the reason why ensuring the reliability of the computational results requires the use of TD-DFT methodologies based on LRC functionals within the TDA framework, possibly in conjunction with a polarizable continuum model.

2.4.1 Theoretical Description of Reverse Intersystem Crossing The RISC process can be considered as an electron transfer reaction from the T1 state to the S1 state. Its rate can be cast in the framework of the Fermi Golden rule and Marcus-Levich-Jortner theory as291

kRISC =

2π S1 H˘ SO T1 

2

1 4ππ M kBT

∞

∑ n=0

e −S

 ( ∆E ST + λ M + nω eff )2  Sn  (2.18) exp  − n! 4λ M kBT    

where S1 Hˆ SO T1 denotes the spin-orbit coupling between the S1 and T1 states; λM denotes the Marcus reorganization energy related with the intermolecular and intramolecular low-frequency vibrations; kB, the Boltzmann constant; T, temperature; ωeff , the effective energy of a mode representing the nonclassical high-frequency intramolecular vibrations ( ωeff kBT >> 1) and S, the effective Huang-Rhys factor corresponding to this mode. A first lesson to be learned from the rate expression is that the singlet-triplet energy gap, which is generally a positive quantity (as ES1 > ET1), needs to be as small as possible. This realization is the basis for the common strategy of designing TADF emitters as a combination of donor-acceptor moieties with a large twist (dihedral) angle between them. In such systems, it is expected that the HOMO is localized on

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FIGURE 2.9  Molecular structures of typical TADF emitters.

the donor and the LUMO on the acceptor. Then, given the simplistic assumption that a single electronic configuration describes both S1 and T1 states and corresponds to an electronic transition from HOMO to LUMO, the electron exchange energy vanishes and ∆EST goes to zero. In this instance, the TDA excitation energies for the S1/T1 states, and ∆EST can be expressed as292:

(

) (

)



  1 E S1 = ( ε L − ε H ) + HL f xc↑,↑ HL + HL f xc↑,↑ HL + 2  HL HL  (2.19) r  



ET1 = ( ε L − ε H ) + HL f xc↑,↑ HL − HL f xc↑,↓ HL (2.20)



  1 ∆EST = 2  HL + f xc↑,↓ HL  (2.21) r  

(

) (

)

where ε H ε L denotes the HOMO (H)/LUMO (L) orbital energy; f xc , the XC kernel; ↑ / ↓ , electron spin up/down; and r, the inter-electron distance. According to Equation (2.21), ∆EST is directly proportional to the spatial overlap between the HOMO and LUMO wavefunctions. Thus, if S1 and T1 both have strictly a HOMO-LUMO CT character, with the HOMO and LUMO spatially separated, ES1 and ET1 become nearly equal. In reality, the lower singlet and triplet excited states possess a more complex, and

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distinct, electronic-configuration description, which necessitates going well beyond the simple HOMOLUMO picture (we note that, if that HOMO-LUMO picture were valid, the absence of HOMO-LUMO spatial overlap would quench any luminescence between S1 and S0). Figure 2.10 illustrates the NTOs (see Section 2.2) for the S1, T1, and T2 states of the PXZ-TRZ and ACRFLCN molecules depicted in Figure 2.9. In both instances, the S1 state has a very strong CT character. In PXZ-TRZ (where the dihedral angle between the D-A moieties reaches nearly 90°), the T1 state has also a dominant CT nature. As a consequence, the singlet-triplet splitting in PXZ-TRZ is as low as 0.06 eV.293 A number of theoretical investigations have demonstrated that local-excitation (LE) triplet states are generally more stable than their CT-dominated counterparts, which is consistent with the fact that the exchange energy increases with the spatial overlap of the relevant wavefunctions.291 Thus, when the energy of the triplet state localized on the donor or acceptor moiety lies below that of the CT triplet state, the T1 state is expected to be an LE-dominated state. For example, in the ACRFLCN molecule, which has a spiro structure and in which the S1 state keeps a substantial CT excitation character (see Figure 2.10), both the NTO hole and electron wavefunctions describing the T1 state are localized on the cyano-substituted fluorene unit. The T2 state, however, corresponds to a CT excited state similar to the S1 state. As a result, the S1-T1 ∆EST in ACRFLCN is significantly larger than in PXZ-TRZ, ca. 0.24 eV.291 The theoretical investigations of Kim and co-workers have demonstrated that, in such twisted D-A TADF molecules, the nature of the lowest triplet states depends on two factors: (i) the energies of the frontier molecular orbitals of the donor and acceptor fragments; and (ii) the electronic couplings between these molecular orbitals. This provides flexibility in selecting the donor and acceptor units and tuning of the molecular structures to optimize the electronic and optical properties of TADF emitters.291,294

FIGURE 2.10  Natural Transition Orbital (NTO) analysis for the S1 and T1 excited states in PXZ-TRZ and ACRFLCN, as calculated by the LC-ωPBE/6-31+G(d) method with the non-empirical tuned range-separation parameter ω. Hole (h) and electron (e) wavefunctions with the largest weight, v, are placed below and above the arrows, respectively. Hydrogen atoms are omitted for the sake of clarity. (Reproduced from Samanta P.K., et al. Up-Conversion Intersystem Crossing Rates in Organic Emitters for Thermally Activated Delayed Fluorescence: Impact of the Nature of Singlet vs Triplet Excited States. Journal of the American Chemical Society 2017, 139(11), 4042–4051. With permission.)

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2.4.2 Relationships of the Spin-Orbit Couplings with the Excitation Characteristics In addition to the singlet-triplet energy gaps, Equation (2.18) shows that the SOCs also play a critical role in T1 → S1 reverse intersystem crossing rates. Given the purely organic nature of the TADF emitters, the SOCs are calculated to be extremely weak, ≤1 cm−1;291,295,296 in contrast, they can reach over 100 cm−1 in second-generation OLED emitters, such as Ir(ppy)3.297 In twisted, CT-type TADF emitters, where both S1 and T1 states show a substantial CT excitation character, the positive outcome is a small ∆EST; however, the drawback is that the SOC between these excited states vanishes, as demonstrated by the calculations of Samanta et al.291 The reason is that, within the one-electron approximation, the spin-orbit operator acts on both the spin magnetic quantum number of the electron and its spatial angular momentum quantum number. Consequently, SOCs between singlet and triplet states with the same spatial orbital occupation are formally zero as any change in spin cannot be compensated by a corresponding change in angular momentum, and thus the total angular momentum would not be conserved. In the case of ACRFLCN, the S1 state has a predominant CT nature, while the T1 state is predominantly a LE state, see Figure 2.10. This difference in the nature of the excited states gives rise to a non-negligible SOC of 0.46 cm−1. On the other hand, since S1 and T2 both have a significant CT character, their SOC goes down to 0.01 cm−1. PXZ-TRZ exhibits a significant SOC (1.54 cm−1) between the S1 and T2 states, which again comes from their marked difference in nature (CT vs. LE state, respectively, see Figure 2.10).291 Note that large spin-orbit couplings between S1 and T2 states can also facilitate RISC processes in TADF materials, as discussed in the following section. Based on Equation (2.18), Samanta et al. evaluated the rates of T1 → S1 RISC processes as a function of ∆EST and SOC for two reasonable values of reorganization energy.291 The results are reproduced in Figure 2.11. When ∆EST  0.25 eV, the RISC rates are reduced to the point that triplet harvesting becomes very difficult. However, recent experimental investigations have

FIGURE 2.11  RISC rate constant (k RISC) as a function of ΔEST for different spin-orbit coupling matrix elements (SOCME) and Marcus reorganization energy (λM) values. (Reproduced from Samanta P.K., et al. Up-Conversion Intersystem Crossing Rates in Organic Emitters for Thermally Activated Delayed Fluorescence: Impact of the Nature of Singlet vs Triplet Excited States. Journal of the American Chemical Society 2017, 139(11), 4042–4051. With permission.)

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FIGURE 2.12  (a) Chemical structures of TADF molecules (C-DBTD and CPT-3) with large singlet-triplet energy gaps. (b) Diagrammatic illustrations of the displacement vectors for the representative low-frequency vibrational normal modes for the molecule C-DBTD. (c) Potential energy surfaces of the excited states T1, T2, S1, and S2 for the molecule C-DBTD. (Reproduced from Chen X.-K., et al. Nature of highly efficient thermally activated delayed fluorescence in organic light-emitting diode emitters: Nonadiabatic effect between excited states. The Journal of Physical Chemistry C 2015, 119(18), 9728–9733. With permission.)

demonstrated that TADF properties can be observed in organic small-molecular emitters with large ∆EST values of ca. 0.30 eV, and good OLED performance can be achieved.298,299 In order to rationalize the data reported for such TADF molecules (e.g. the molecules shown in Figure 2.12), Chen et al. relied on a more general, non-radiative transition rate formula derived from TD second-order perturbation theory, which includes both non-adiabatic and spin-orbit couplings295:



kT1 →S1

2π = 

∑∑ P

vT1

v S1

vT1

H S′1vS1 ,T1vT1 +

∑ nµ

2

H S′1vS1 ,nµ H n′ ,T1vT1 δ ( ET1vT1 − ES1vS1 ) (2.22) ET1vT1 − Enµ

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ˆ′=H ˆ SO + H ˆ BO , where H ˆ BO represents the non-adiHere, the interaction Hamiltonians are written as H ˆ SO represents the abatic coupling (NAC) due to the breakdown of the adiabatic approximation and H

spin-orbit coupling. E n ( n ≠ S1 and T1 ) denotes the sum of the energies of the nth adiabatic electronic state and its vibrational state. vS1 and vT1 denote the quantum numbers of vibrational states corresponding to the S1 and T1 states. As discussed above, the SOCs between the S1 and T1 states in purely organic molecules are usually very small; thus, H S′1 ,T1 in Equation 22 (i.e. the first-order term in perturbation theory) has limited contribution to the RISC rate. When the D-A twist angles in such molecules (e.g. carbazole-dibenzothiopheneS,S-dioxide and carbazole-phenyl-triazine derivatives in Figure 2.12, where referred to as C-DBTD and CPT-3, respectively) approach 90° through molecular vibrations, several low-lying excited states (i.e. S1 and S2, as well as T1 and T2) start playing an important role in the RISC process via their contributions to the second term in perturbation theory in Equation 22; this second-order term can be rewritten as:

 c j SOCS1T2  ∂ T1 +  T2 ∂Q j T1 − ET2 

∑  E j

∑S

1

j

 c j SOCS2T1 ∂ S2  ∂Q j  ET1 − E S2

  (2.23) 

∂ Φm is the NAC between adiabatic electronic states m and n with the same spin multi∂Q j plicity; Q j denotes the coordinates of the jth vibration normal mode. SOCS1T2 [ SOCS2T1 ] represents the spin-orbit coupling between the S1 [S2] and T2 [T1] states. The NAC (or vibronic coupling) between two adiabatic excited states is mediated by molecular vibrations. In the C-DBTD molecule, several vibrational modes mainly related to rotations of the carbazole groups, are particularly relevant; these vibrational modes have energies of  0) leads to an increase in positive charge density at the top of the adsorbate, resulting in a decrease in the electrostatic potential energy felt by an electron away from the interface, and contributes to a reduction in surface workfunction. Conversely, a negative dipole layer, corresponding to an increased negative charge density on top of the adsorbate, increases the

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FIGURE 2.20  Comparison of the workfunction change evaluated by DFT (dark gray) and measured by UPS (light gray) as a function of the calculated μz.mol values, for the six PA molecules listed from left to right in the same order as the data points. Note that the calculated ΔΦ values are those very crudely extrapolated for a coverage of ∼8.4 × 1013 molecules/cm2 by tripling the ΔΦ values calculated for coverages of ∼2.8 × 1013 molecules/cm2. (Adapted from Hotchkiss P.J., et al. Modification of the surface properties of indium tin oxide with benzylphosphonic acids: a joint experimental and theoretical study. Adv. Mater. 2009, 21(44), 4496.)

workfunction. The workfunction change ∆Vmol induced by such a dipole layer can be estimated using the Helmholtz equation:

∆Vmol = −

eN µ z .mol (2.25) εr ε0

where N denotes the surface coverage (the number of adsorbed molecules per unit area); µ z ,mol , the average adsorbate molecular dipole moment along the surface normal direction; and ε0 and εr, the vacuum and relative permittivity of the isolated adsorbate molecular layer. To illustrate the impact of the ∆Vmol term, we have considered a series of six fluorobenzyl-substituted PA molecules on the ITO surface. In this case, the ΔVint.dip and ∆Vgeom.rel components are identical since

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FIGURE 2.21  Change in the total workfunction (∆Φ) and its three contributing components (a) and the average molecular dipole and interface dipole per molecule along the surface normal direction (b), as a function of surface coverage for trifluorophenyl-PA modified ITO surface. (Adapted from Li H., et al. Theoretical study of the surface modification of indium tin oxide with trifluorophenyl phosphonic acid molecules: impact of coverage density and binding geometry. J Mater Chem 2010, 20(13), 2630–2637.)

they depend on the binding mode of the PA moieties to ITO, which is the same across the series. Thus, according to Equation (2.25), the workfunction modification should follow a linear relationship with respect to the surface normal component of the PA molecular dipole (note that the magnitude of this component obviously depends on the orientation of the benzyl group with respect to the ITO surface). Figure 2.20 illustrates that the linear relationship indeed holds very well, both experimentally and theoretically. Experimental data show that by varying the number and locations of the fluorines on the PA benzyl group, workfunction changes by as much as 1.2–1.3 eV can be achieved. The discrepancy between the slopes of the linear theoretical and experimental evolutions originates in the too small coverages initially considered in the calculations (2.8 × 1013 molecules/cm2), which prevents an explicit account of the impact of intermolecular interactions on the dipole moments. Subsequent DFT studies of an ITO surface modified by trifluorophenyl-PA molecules investigated coverages ranging from 2.8 × 1013 to 1.0 × 1014 molecules/cm 2. When the molecules are more densely packed, the intermolecular interactions act to significantly decrease the average molecular dipole moment along the surface normal direction, see Figure 2.21.401 This corresponds to the depolarization effect in densely packed molecular layers, which has also been studied for substituted biphenylthiol and (fluoro)methylthiol SAMs chemisorbed on the Au(111) surfacee.266,403,404 The impact of such depolarization effects is the sublinear evolution of ∆Vmol and ∆Φ as a function of surface coverage, see Figure 2.21a. We now turn to the interface dipole, ΔVint.dip, whose contribution can similarly be evaluated by applying Helmholtz equation:

∆Vint.dip =

eN µ int.dip (2.26) εr ε0

where µ int.dip is the average interface dipole per adsorbed molecule. Figure 2.21b shows that the average µ int.dip only slightly decreases with an increase in surface coverage. This indicates that the interface dipole and the corresponding charge redistribution induced by the chemisorption is a more local effect that is less affected by the presence of neighboring molecules. As such, the ΔVint.dip contribution to the workfunction change is expected to display a more linear correlation with surface coverage than ∆Vmol. This is indeed observed in Figure 2.21a, especially when the coverage is less than three molecules per surface unit cell (for more details, see the discussion of Table 2 in Reference 401).

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The third component, ∆Vgeom.rel, is generally found to have negligible contribution to the workfunction change in the case of metal surfaces. However, its contribution to the workfunction change in conducting oxide surfaces modified by chemically or physically adsorbed organic layers can be significant and sometimes comparable to the other two factors.401,405,406

2.6.2 Brief Description of the Computational Methodology Both the bare surface and the modified surface are usually modeled by a slab containing a few atomic layers of the substrate, plus the adsorbate layer when investigating an interface. Periodic boundary conditions in three dimensions are usually considered, which means that along the direction perpendicular to the slab a large enough empty (“vacuum”) space is added in order to prevent any interaction among the repeated slabs, the so-called repeated slab approach. (As a side note, in some DFT packages, such as CRYSTAL and the BAND module in ADF, the periodic boundary conditions are applied only in two dimensions within the surface plane; in these instances, the surface slab can thus be treated as a “real” two-dimensional system). Geometry optimizations are needed to relax the surface or interface system to its minimum energy state. In this step, the adsorbate layer and a few atomic layers of the top surface are relaxed while the surface bottom layers are fixed at their positions in the bulk. In these optimizations, it is important to scan as many adsorption configurations as possible to eventually reach the global minimum. In recent years, computational methods combining DFT calculations, interatomic potential modeling, Monte Carlo simulations, and machine-learning algorithms have been developed to accelerate such calculations.407,408 In addition, since van der Waals forces can play a significant role in the intermolecular and molecule/substrate interactions, it is important for them to be included in the DFT functionals adopted for the geometry optimizations. Various ways to do this have been discussed in Section 2.2. The ∆Vmol and ΔVint.dip terms can be calculated individually by partitioning the interface into its constituents, that is, the molecular layer and the bare surface, both kept at the optimized geometry for the interface. In the case of the molecular layer, ∆Vmol is calculated as the plane-averaged electrostatic energy difference for an electron on the two sides of the molecular layer. The ΔVint.dip term is calculated by solving Poisson’s equation:

d 2V ( z ) 1 = − ∆ρ ( z ) (2.27) dz 2 ε0

where V(z) denotes the electrostatic potential energy and ∆ρ(z), the change in the plane-averaged charge density at the interface, induced by the molecular layer adsorption; the latter is calculated as the charge density difference between the interface and the two isolated components, all at the optimized geometries for the interface. The detailed analysis of the interface dipole provides information on the dipole direction and whether there occurs a net charge transfer across the interface. For the sake of illustration, Figure 2.22 shows the results obtained in the case of the ITO-F4TCNQ interface.409 In the top panel, a large peak in the charge density difference appears at the level of the F4TCNQ layer. This points to a significant electron accumulation into F4TCNQ (as can be expected from its electron-acceptor nature) and a corresponding electron depletion on the ITO surface, leading to a negative interface dipole along the surface normal. Here, the amount of electron transfer from ITO to F4TCNQ is about 0.8 e, which results in a workfunction increase of 1.5 eV. The modeling of metal-oxide surfaces is more complex than that of metal surfaces, due to the possibility of large variations in surface stoichiometry and termination related to different growth conditions, the presence of intrinsic or extrinsic surface defects, and the appearance of surface reconstructions. Taking common transparent conducting oxides such as ITO and ZnO as examples, their surfaces are expected to be covered with hydroxyl groups, whose coverage densities depend on the surface pre-treatment methods. Also, ZnO has surface terminations that can be polar (for instance, the Zn-terminated

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FIGURE 2.22  Plane-averaged charge density difference, charge transfer, and interface dipole for the ITOF4TCNQ interface. (Adapted from Li H., et al. Surface modification of indium-tin-oxide via self-assembly of a donor-acceptor complex: a density functional theory study. Adv. Mater. 2012, 24(5), 687.)

(0001) and O-terminated (000–1) surfaces) or non-polar (for instance, the (10–10) and (11–20) surfaces); various intrinsic defects (for instance, zinc interstitials (Zni) or oxygen vacancies (VO)) can be present and act as possible source of n-type doping; in addition, the polar surfaces can easily undergo surface reconstructions, which brings even more complexity. The polar surfaces of ionic crystals usually have non-zero dipole moments within a surface repeat unit and are categorized as Tasker type-III surface. The modeling of this type of surfaces can be problematic since the superposition of individual dipole layers can induce a diverging surface energy. In repeated slab DFT calculations, this problem has caused spurious electron transfer from the anionic side of a slab to the cationic side due to the much lower electrostatic potential energy for an electron on the cationic side. This has led to claims of the presence of metallic surface states for ZnO (0001).410–412 The issue can be solved by modifying the charge state of the oxygen in the bottom oxygen layer, as done in the work of Kresse and co-workers,413 or by saturating the oxygen layer with a layer of virtual hydrogen atoms containing ½|e| nucleus charge and -½|e| electron charge.414 A rigorous way to determine the most stable surface structure for a system with different surface stoichiometry is to compare their surface energy defined as:

γ(T ,{ pi }) =

1  Gslab (T ,{ pi }, N i ) − 2 A 



∑ n µ (T , p ) (2.28) i i

i

i

where G(T ,{ pi }, N i ) denotes the Gibbs free energy (G = E+PV-TS) of a given slab with two equivalent surfaces of area A; and ni and µi (T , pi ) are the number of atoms and the chemical potential of each constituting species i in the system. In standard DFT calculations, the Gibbs free energy is usually taken as the total energy E of the system, which neglects the temperature and pressure dependent terms. We

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note that, over the past ten years, first-principles thermodynamics methodologies have been developed to investigate redox reactions under high temperature and pressure conditions, and these two terms have been taken into account through different approximations.415–419 The chemical potential of each constituting species is usually constrained by equilibrium conditions related to the crystal growth process. For instance, assuming that each species satisfies the equilibrium condition between its gas phase and elemental bulk phase, the chemical potential must satisfy the condition µi ≤ µbulk . On the other hand, since the product bulk solid should also be in equilibrium with all i the constituting species in the gas phase acting as chemical reservoirs, the chemical potential of each species must satisfy:

∑n µ = ∑n µ

bulk i i

i i

i

+ ∆H f (2.29)

i

where ∆H f is the heat of formation of the product bulk solid. Thus, these two conditions provide upper and lower limits for the chemical potential of each species as:

µbulk + ∆H f ≤ µi ≤ µbulk (2.30) i i

2.6.3 Surface Defects In most instances, surface defects are detrimental to the performance of opto-electronic devices. For example, when the defects lead to the appearance of deep trap states, these can act as non-radiative chargerecombination centers through the Shockley-Read-Hall recombination scheme. In the early studies of polymer/semiconductor nanocrystal hybrid solar cells that were based on semiconductors such as CdSe, CdTe, PdS, and ZnO, surface defects have been identified as the main reason contributing to the poor PCE.420 Thus, when modeling conducting oxide surfaces exploited in opto-electronic applications, it is important to carefully evaluate the various intrinsic and extrinsic surface defects, their impact on the workfunction, and their electronic levels with respect to the surface Fermi level and band edge states. The presence of surface or near-surface defects is susceptible to substantially alter the surface electrostatic potential, and thus to impact the workfunction. Our DFT studies modeling the ITO,402 ZnO,414,421 and MoOx422 surfaces have indeed confirmed that the workfunctions of conducting oxide surfaces are highly dependent on their respective surface stoichiometry. For instance, in the case of ITO, we initially considered a fully hydroxylated surface model with a hydroxyl (OH) coverage as high as 6.8 × 1014 OH groups per cm2; the workfunction is then evaluated to be as low as 3.2 eV,396 which is much lower than the experimentally reported values in the rage of 4.0–5.2 eV.394,395 However, when the number of surface hydroxyl groups is gradually reduced, the workfunction of the model ITO surface increases up to 4.2 eV when the OH coverage is made to correspond to the experimental coverage of about one-eighth of the full coverage.395,402 In the case of polar and non-polar ZnO surfaces, DFT calculations have examined the impact of various intrinsic surface defects, such as zinc vacancies (VZn), oxygen vacancies (VO), or zinc interstitials (Zni).414,421,423 We have found that the presence of zinc interstitials as n-type dopants can reduce the workfunction of both the polar and non-polar ZnO surfaces by ca. 1 eV; this is related to the change in the surface electrostatic potential as well as the shift of the Fermi level for the n-doped ZnO surfaces. These results are consistent with the experimental observations that the workfunction of the ZnO(0001) polar surface can be reduced by as much as 0.75 eV by increasing the n-type charge-carrier density.424–426 On the other hand, the presence of VZn vacancies on the Zn-terminated polar surfaces (which have been found to be the energetically most stable intrinsic surface defects414) can increase the workfunction by ca. 1.2 eV. Overall, the main lesson learned from these studies is that the surface workfunction very much depends on the surface composition, which can be experimentally tuned by using different surface treatment methods.

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FIGURE 2.23  Total workfunction change and energy-level alignments upon the adsorption of PTCDI on ZnO(10-10)surface with different surface stoichiometries. (Adapted from Winget, P., et al. Defect-driven interfacial electronic structures at an organic/metal-oxide semiconductor heterojunction. Adv. Mater. 2014, 26(27), 4711–4716.)

2.6.4 Charge-Transfer Characteristics for Donor/Acceptor Molecules Physisorbed on Metal-Oxide Surfaces In OLED and OSC applications, it is common to use transparent conducting metal-oxides, such as TiO2, ZnO, MoO3, WO3, or NiOx, as an electron or hole injection or extraction interlayer between the corresponding electrode and the active electron or hole transport layer. The interactions between the electron- or hole-transport organic semiconducting layers deposited on these metal-oxide interlayers are critical in determining the interfacial charge transfer between the organic and inorganic components. An interesting example is given by the interface system consisting of a monolayer of the organic electron-acceptor 3,4,9,10-perylene-tetracarboxylicdiimide (PTCDI) deposited on the non-polar ZnO(10-10) surface. We considered the stoichiometric ZnO surface as well as non-stoichiometric surfaces containing oxygen vacancies or zinc interstitials, see Figure 2.23.421 Upon adsorption of PTCDI molecules on the stoichiometric or oxygen-deficient ZnO(10-10) surfaces, the DFT calculations result in moderate charge transfers (1 V open circuit voltages. Energy and Environmental Science 2016, 9(12), 3783–3793. 389. Gao, L.; Zhang, Z.-G.; Xue, L.; Min, J.; Zhang, J.; Wei, Z.; Li, Y., All-polymer solar cells based on absorption-complementary polymer donor and acceptor with high power conversion efficiency of 8.27%. Advanced Materials 2016, 28(9), 1884–1890. 390. Benten, H.; Mori, D.; Ohkita, H.; Ito, S., Recent research progress of polymer donor/polymer acceptor blend solar cells. Journal of Materials Chemistry A 2016, 4(15), 5340–5365. 391. Zhao, W.; Qian, D.; Zhang, S.; Li, S.; Inganäs, O.; Gao, F.; Hou, J., Fullerene-free polymer solar cells with over 11% efficiency and excellent thermal stability. Advanced Materials 2016, 28(23), 4734–4739. 392. Zhang, J.; Jiang, K.; Yang, G.; Ma, T.; Liu, J.; Li, Z.; Lai, J. Y. L.; Ma, W.; Yan, H., Tuning energy levels without negatively affecting morphology: A promising approach to achieving optimal energetic match and efficient nonfullerene polymer solar cells. Advanced Energy Materials 2017, 7(15), 1602119. 393. Marder, S. R.; Bredas, J.-L., The WSPC Reference on Organic Semiconductors: Organic Semiconductors. World Scientific, Singapore: 2016; Vol. 2. 394. Gassenbauer, Y.; Schafranek, R.; Klein, A., Surface states, surface potentials, and segregation at surfaces of tin-doped In2O3. Physical Review B 2006, 73(24), 245312. 395. Harvey, S. P.; Mason, T. O.; Gassenbauer, Y.; Schafranek, R.; Klein, A., Surface versus bulk electronic/defect structures of transparent conducting oxides: I. Indium oxide and ITO. Journal of Physics D: Applied Physics 2006, 39(18), 3959–3968. 396. Hotchkiss, P. J.; Li, H.; Paramonov, P. B.; Paniagua, S. A.; Jones, S. C.; Armstrong, N. R.; Brédas, J. L.; Marder, S. R., Modification of the surface properties of indium tin oxide with benzylphosphonic acids: A joint experimental and theoretical study. Advanced Materials 2009, 21(44), 4496. 397. Zhou, Y. H.; Fuentes-Hernandez, C.; Shim, J.; Meyer, J.; Giordano, A. J.; Li, H.; Winget, P.; Papadopoulos, T.; Cheun, H.; Kim, J.; Fenoll, M.; Dindar, A.; Haske, W.; Najafabadi, E.; Khan, T. M.; Sojoudi, H.; Barlow, S.; Graham, S.; Bredas, J. L.; Marder, S. R.; Kahn, A.; Kippelen, B., A universal method to produce low-work function electrodes for organic electronics. Science 2012, 336(6079), 327–332.

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398. Jorgensen, M.; Norrman, K.; Krebs, F. C., Stability/degradation of polymer solar cells. Solar Energy Materials and Solar Cells 2008, 92(7), 686–714. 399. Heimel, G.; Romaner, L.; Brédas, J. L.; Zojer, E., Interface energetics and level alignment at covalent metal-molecule junctions: Pi-conjugated thiols on gold. Physical Review Letters 2006, 96(19), 196806. 400. Heimel, G.; Romaner, L.; Zojer, E.; Brédas, J.-L., The interface energetics of self-assembled monolayers on metals. Accounts of Chemical Research 2008, 41(6), 721–729. 401. Li, H.; Paramonov, P.; Bredas, J. L., Theoretical study of the surface modification of indium tin oxide with trifluorophenyl phosphonic acid molecules: Impact of coverage density and binding geometry. Journal of Materials Chemistry 2010, 20(13), 2630–2637. 402. Li, H.; Winget, P.; Bredast, J. L., Transparent conducting oxides of relevance to organic electronics: Electronic structures of their interfaces with organic layers. Chemistry of Materials 2014, 26(1), 631–646. 403. Li, H.; Duan, Y. Q.; Paramonov, P.; Coropceanu, V.; Bredas, J. L., Electronic structure of selfassembled (fluoro)methylthiol monolayers on the Au(111) surface: Impact of fluorination and coverage density. Journal of Electron Spectroscopy 2009, 174(1–3), 70–77. 404. Romaner, L.; Heimel, G.; Ambrosch-Draxl, C.; Zojer, E., The dielectric constant of self-assembled monolayers. Advanced Functional Materials 2008, 18(24), 3999–4006. 405. Li, H.; Ratcliff, E. L.; Sigdel, A. K.; Giordano, A. J.; Marder, S. R.; Berry, J. J.; Bredas, J. L., Modification of the gallium-doped zinc oxide surface with self-assembled monolayers of phosphonic acids: A joint theoretical and experimental study. Advanced Functional Materials 2014, 24(23), 3593–3603. 406. Wood, C.; Li, H.; Winget, P.; Bredas, J. L., Binding modes of fluorinated benzylphosphonic acids on the polar ZnO surface and impact on work function. Journal of Physical Chemistry C 2012, 116(36), 19125–19133. 407. Deacon-Smith, D. E. E.; Scanlon, D. O.; Catlow, C. R. A.; Sokol, A. A.; Woodley, S. M., Interlayer cation exchange stabilizes polar perovskite surfaces. Advanced Materials 2014, 26(42), 7252–7256. 408. Mora-Fonz, D.; Lazauskas, T.; Woodley, S. M.; Bromley, S. T.; Catlow, C. R. A.; Sokol, A. A., Development of interatomic potentials for supported nanoparticles: The Cu/ZnO case. Journal of Physical Chemistry C 2017, 121(31), 16831–16844. 409. Li, H.; Winget, P.; Bredas, J. L., Surface modification of indium-tin-oxide via self-assembly of a donor-acceptor complex: A density functional theory study. Advanced Materials 2012, 24(5), 687. 410. Carlsson, J. M., Electronic structure of the polar ZnO{0001}-surfaces. Computational Materials Science 2001, 22(1–2), 24–31. 411. Noguera, C., Polar oxide surfaces. Journal of Physics: Condensed Matter 2000, 12(31), R367–R410. 412. Wander, A.; Schedin, F.; Steadman, P.; Norris, A.; McGrath, R.; Turner, T. S.; Thornton, G.; Harrison, N. M., Stability of polar oxide surfaces. Physical Review Letters 2001, 86(17), 3811–3814. 413. Kresse, G.; Dulub, O.; Diebold, U., Competing stabilization mechanism for the polar ZnO(0001)-Zn surface. Physical Review B 2003, 68, 245409. 414. Li, H.; Schirra, L. K.; Shim, J.; Cheun, H.; Kippelen, B.; Monti, O. L. A.; Bredas, J. L., Zinc oxide as a model transparent conducting oxide: A theoretical and experimental study of the impact of hydroxylation, vacancies, interstitials, and extrinsic doping on the electronic properties of the polar ZnO (0002) surface. Chemistry of Materials 2012, 24(15), 3044–3055. 415. Chase, M. W.; Davies, C. A.; Downey, J. R.; Frurip, D. J.; Mcdonald, R. A.; Syverud, A. N., Janaf thermochemical tables – 3rd edition. Journal of Physical and Chemical Reference Data 1985, 14, 927–1856. 416. Hu, Y.; Boudoire, F.; Hermann-Geppert, I.; Bogdanoff, P.; Tsekouras, G.; Mun, B. S.; Fortunato, G.; Graetzel, M.; Braun, A., Molecular Origin and Electrochemical Influence of Capacitive Surface States on Iron Oxide Photoanodes. Journal of Physical Chemistry C 2016; Vol. 120, pp 3250–3258. 417. Rossmeisl, J.; Qu, Z. W.; Zhu, H.; Kroes, G. J.; Norskov, J. K., Electrolysis of water on oxide surfaces. Journal of Electroanalytical Chemistry 2007, 607(1–2), 83–89.

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418. Valdes, A.; Qu, Z. W.; Kroes, G. J.; Rossmeisl, J.; Norskov, J. K., Oxidation and photo-oxidation of water on TiO2 surface. Journal of Physical Chemistry C 2008, 112(26), 9872–9879. 419. Noh, J. L., H.; Osman, O. I.; Aziz, S. G.; Winget. P.; Bredas, J. L., Impact of hydroxylation and hydration on the reactivity of α-Fe2O3 (0001) surfaces under environmental and electrochemical conditions. Advanced Energy Materials 2018, 8(21). doi.org/10.1002/aenm.201800545 420. Gao, F.; Ren, S. Q.; Wang, J. P., The renaissance of hybrid solar cells: Progresses, challenges, and perspectives. Energy and Environmental Science 2013, 6(7), 2020–2040. 421. Winget, P.; Schirra, L. K.; Cornil, D.; Li, H.; Coropceanu, V.; Ndione, P. F.; Sigdel, A. K.; Ginley, D. S.; Berry, J. J.; Shim, J.; Kim, H.; Kippelen, B.; Brédas, J.-L.; Monti, O. L. A., Defect-driven interfacial electronic structures at an organic/metal-oxide semiconductor heterojunction. Advanced Materials 2014, 26(27), 4711–4716. 422. Papadopoulos, T. A.; Meyer, J.; Li, H.; Guan, Z.; Kahn, A.; Brédas, J.-L., Nature of the interfaces between stoichiometric and under-stoichiometric MoO3 and 4,4′-N,N′-dicarbazole-biphenyl: A combined theoretical and experimental study. Advanced Functional Materials 2013, 23(48), 6091–6099. 423. Li, H.; Bredas, J. L., Comparison of the impact of zinc vacancies on charge separation and charge transfer at ZnO/sexithienyl and ZnO/fullerene interfaces. Advanced Materials 2016, 28(20), 3928. 424. Moormann, H.; Kohl, D.; Heiland, G., Work function and band bending on clean cleaved zincoxide surfaces. Surface Science 1979, 80, 261–264. 425. Moormann, H.; Kohl, D.; Heiland, G., Variations of work function and surface conductivity on clean cleaved zinc-oxide surfaces by annealing and by hydrogen adsorption. Surface Science 1980, 100, 302–314. 426. Parker, T. M.; Condon, N. G.; Lindsay, R.; Leibsle, F. M.; Thornton, G., Imaging the polar (0001 )̄ and non-polar (101 0̄ ) surfaces of ZnO with STM. Surface Science 1998, 415(3), L1046–L1050. 427. Vaynzof, Y.; Bakulin, A. A.; Gelinas, S.; Friend, R. H., Direct observation of photoinduced bound charge-pair states at an organic-inorganic semiconductor interface. Physical Review Letters 2012, 108(24), 246605. 428. Niederhausen, J.; Amsalem, P.; Wilke, A.; Schlesinger, R.; Winkler, S.; Vollmer, A.; Rabe, J. P.; Koch, N., Doping of C-60 (sub)monolayers by Fermi-level pinning induced electron transfer. Physical Review B 2012, 86(8), 081411. 429. Timpel, T. L. H.; Nardi, M. V.; Wegner, B.; Frisch, J.; Hotchkiss, P. J.; Marder, S.; Barlow, S.; Bredas, J. L.; Koch, N., Electrode work function engineering with phosphonic acid monolayers and molecular acceptors: Charge redistribution mechanisms. Advanced Functional Materials 2017, 28(15), 1801349. 430. Paramonov, P. B.; Paniagua, S. A.; Hotchkiss, P. J.; Jones, S. C.; Armstrong, N. R.; Marder, S. R.; Brédas, J.-L., Theoretical characterization of the indium tin oxide surface and of its binding sites for adsorption of phosphonic acid monolayers. Chemistry of Materials 2008, 20(16), 5131–5133. 431. Pehlke, E.; Scheffler, M., Evidence for site-sensitive screening of core holes at the Si and Ge (001) surface. Physical Review Letters 1993, 71(14), 2338–2341. 432. Köhler, L.; Kresse, G., Density functional study of CO on Rh(111). Physical Review B 2004, 70(16), 165405.

3 Perspective on the Advancements in Conjugated Polymer Synthesis, Design, and Functionality over the Past Ten Years 3.1 Introduction to this Perspective...................................................... 107 Polymer Structures • Polymer Synthesis

3.2 Advancements in Conjugated Polymer Syntheses........................114

Brian Schmatz, Robert M. Pankow, Barry C. Thompson, and John R. Reynolds

Emerging Repeat Units • New Synthetic Strategies in Conjugated •  Polymer Chemistry • Structure Property Modification of Conjugated Polymers

3.3 Future Direction and Outlook..........................................................137 Efficient Monomer and Polymer Synthesis • Polymer Properties and Applications

Acknowledgments���������������������������������������������������������������������������������������� 140 References������������������������������������������������������������������������������������������������������ 140

3.1 Introduction to this Perspective In the decades following the seminal work on conductive polyacetylene in the 1970s, there has been an immense effort and numerous accomplishments in the synthesis and development of conjugated polymers. Through the development of new conjugated polymer systems, the field evolved allowing scientists to address a large number of diverse properties and applications including light emission, charge transport, electrochromism, photovoltaics, and bioelectronics. At the heart of each of these is the development of a deep understanding of structure-property relationships. Attaining optimized structures has required a broad range of molecular and polymeric synthetic advances, which have been detailed in multiple chapters in the first three editions of the Handbook of Conducting Polymers (HBCP). Early work in the field tended to focus on specific structural polymer families, such as polyanilines, polypyrroles, polyacetylenes, poly(arylene vinylenes), and polythiophenes. As the field has matured, so has the chemist’s ability to design a desired structure, then accomplish a complicated synthesis that yields a new material with targeted and optimized properties. In this chapter, we aim to provide a 107

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perspective on the state-of-the-art in conjugated polymer chemistry. Our goal is to be complementary to the many excellent chapters that follow in this Handbook. Going beyond the simple “make it and measure it” approach, we acknowledge the importance of processing for controlled morphology, and expand on the possibilities in further reactivity of the polymer systems, which allows many new properties to be induced, fine-tuned, and ultimately optimized for practical utility.

3.1.1 Polymer Structures Conjugated polymers have progressed structurally from homopolymers with minimal functionalization to advanced copolymers where the monomers require complex syntheses. This transformative process in polymer structure is illustrated by the progression from the simplest example of polyacetylene to poly[​ 4,8-b​is(5-​(2-et​hylhe​x yl)t​h ioph​en-2-​yl)be​n zo[1​, 2-b:​4,5-b​ʹ ]dit ​h ioph​ene2,​6-diy​l-alt​-(4-(​2-eth​ylhex ​yl)3​-fluo​rothi​eno[3​,4-b]​t hiop​hene-​)-2-c​arbox​ylate​-2-6-​diyl)​] (PTB7-Th), shown in Figure 3.1.[1] By simply analyzing the changes and advancements made in conjugated polymer syntheses, one can see how the field has progressed. As shown in Figure 3.1, the structural progression starting from polyacetylene is marked by a greater focus on the inclusion of heteroatoms and alkyl substituents, as exemplified by poly(3hexylthiophene) (P3HT), altering the electronic states and providing the material melt and solution processability. Expanding the structures to conjugated copolymers ushered in a new approach to modifying electronic and physical properties, illustrated with poly[​N-9ʹ-​hepta​decan​yl-2,​7-car​bazol​e-alt​-5, 5-​(4ʹ,7​ʹ-di-​2-thi​enyl-​2ʹ,1ʹ​,3ʹ-b​enzot​hiadi​azole​)] (PCDTBT), which was state-of-the-art around the time of publication of the previous edition of the HBCP. Finally, an improved understanding of how finetuning polymeric structure can alter the properties of the materials allowed for the realization of more complex conjugated polymer structures, such as PTB7-Th and semi-random polymers, which are topics of discussion for this edition. In order to provide a context of where the field has been, and in what direction it is heading regarding polymer syntheses, the purpose of this first section is to provide a brief overview of conjugated polymers focused on at the time of the last HBCP edition—although these materials may have been prevalent within the chemical literature prior.[2] The desired impression left upon the reader is twofold. First, that the design and synthetic execution of modern-day conjugated polymers is built upon an improved foundation in synthetic methodology. This becomes apparent when comparing the polymer structures and polymerization methods between the different sections of this chapter, such as Sections 3.1 and 3.2. Many of the

FIGURE 3.1  The progression of conjugated polymers from the simplest form, polyacetylene, to those of increased complexity in terms of repeat units (PTB7-Th) or polymeric architecture (Semi-Random).

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structural motifs described in Section 3.1 are thiophene or phenylene derivatives, where those in the later sections possess different heterocyclic structures with varying functionalities. Also, polymerization methods described in Section 3.1, such as Wittig or Gilch polymerizations (Section 3.1.2), are less employed due to advancements in transition metal catalyzed polymerizations. Second, changes in conjugated polymer structure reflect changes in the desired performance for a given application, such as tailoring structure to alter light absorption, light emission, or charge transport, which can be seen when comparing the structural advancements of PTB7-Th and semi-random polymers—relative to PCDTBT (Figure 3.1) or the polythiophenes or phenylene-based polymers presented in Sections 3.1.1.1–3.1.1.4. 3.1.1.1 Polythiophene and Derivatives Polythiophenes and derivatives were, and still are, intensely studied due to the large number of structural variations and polymerizations possible, opening doors to a seemingly infinite number of potential polymers.[3] Figure 3.2 highlights some widely used materials, namely poly(3-alkylthiophenes) (P3ATs), poly(thienothiophenes), and polydioxythiophenes. The key to thiophenes’ enduring popularity is that thiophene itself can be structurally modified in nearly every position on the ring to afford materials with vastly different physical and electronic properties, illustrated in Figure 3.2 with polymers 1–4, although many other variations exist. This ease of tuning has made thiophene and related heterocycles a staple in the field of conjugated polymers. While unsubstituted polythiophene itself is insoluble and not processable, the inclusion of an alkyl chain, such as with P3HT, provides a polymer that can be solution processed and more accessibly studied. The fused, bicyclic structure of thienothiophenes provides a similar high level of structural modification, where the points of fusion for the bicyclic ring can be adjusted as shown in 5–7, and various substituents added to alter the physical and electronic properties.[3d,4] Polydioxythiophenes (8–10), and the poly(3,4-alkylenedioxythiophenes) in particular, provide a highly electron-rich monomer, due to the attached alkoxy units, thus resulting in a conjugated polymer that is both easily oxidized and stable in the oxidized form.[5] This enhanced stability of the oxidized form is not common or present with many other conjugated polymers and provides dioxythiophenes a unique electronic position within the field, particularly as stable conductors and in electrochromics. 3.1.1.2 Poly(arylene vinylenes) Poly(phenylene vinylenes) (PPVs) and their heterocyclic counterparts, shown as 11–15 in Figure 3.3, provide desirable properties for applications chiefly in organic light emitting diodes (OLED) and polymer solar cells.[6] Unlike the thiophene-based materials discussed, which are described by the aryl

FIGURE 3.2  Examples of polythiophene, polythienothiophene, and polydioxythiophene type polymers.

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FIGURE 3.3  Examples of poly(arylenevinylene) based conjugated polymers.

FIGURE 3.4  Poly(phenyleneethylene) (PPE) based polymers.

monomer(s) incorporated into the polymer backbone, the characteristic feature of these conducting polymers is the vinylic linkage between each aryl unit. Incorporation of these materials into avante garde settings, such as biological probes, shows their continuing relevance in the chemical literature.[7] However, their tunability and electronic properties are limited relative to that of P3ATs and related polymers discussed above, which is largely due to added synthetic complexity and monomer instability. 3.1.1.3 Poly(arylene ethynylenes) Poly(phenylene ethynylenes) (PPEs) are the acetylene counterpart of PPV and similar ethylene linked conjugated polymers, as shown in Figure 3.4. They are highly luminescent materials that have found extensive applications exploiting this property, including sensory and antimicrobial applications.[8] Only phenylene-based polymers are shown, but heterocycles have been incorporated into an ethynylene linked architecture.[9] The modification of the pendant groups on the phenylene moiety provides an excellent handle for tuning the physical and electronic properties of these materials. This is illustrated in Figure 3.4 with the examples of 16, an alkoxy substituted phenylene ethynylene, 17, an iptycene based copolymer, and 18, a water-soluble PPE-based polymer with carboxylic acid groups.[10]

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3.1.1.4 Narrow Bandgap Polymers The desirable optical properties of narrow bandgap polymers defined here as Eg < 1.5 eV, λonset = 827 nm, led to the intensive study of these materials, mainly in search for high-performing polymers for solar cell applications. Commonly encountered examples of these materials, which were emphasized in the previous edition of this Handbook, are shown in Figure 3.5. Briefly, tuning of the bandgap can be accomplished through the careful selection of monomers for the preparation of a copolymer, with the goal of lowering the energy of the conduction band or raising the energy of the valence band.[5a] A more detailed discussion is provided in Chapter 8 of the Handbook. One often-used approach to accomplishing a narrow band gap copolymer is to incorporate an electron-rich monomer, the donor, and an electron deficient monomer, the acceptor, to provide a donor-acceptor copolymer, as illustrated in Figure 3.5 with polymers 19 and 20.[11] Some narrow band gap polymers have already been presented, but will be repeated here since the merits of their optical properties were not previously disclosed. The examples provided in Figure 3.5 are meant to illustrate a homopolymer with an elevated valence band poly(3,4-ethylenedioxythiophene) (PEDOT), a lowered conduction band and increased valence band, poly(thineo[3,4-b]pyrazine) (PTP), and donor-acceptor copolymers, 19 and 20, where the valence and conduction bands are simultaneously elevated and lowered in energy, respectively.

3.1.2 Polymer Synthesis The method implemented for polymer synthesis is largely dictated by which polymeric structure is desired. While some conjugated polymer architectures can be prepared using various methods, many are limited to a single method or can be prepared more easily using a specific one. At the time of publication of the previous edition, many of the conjugated polymers being studied were prepared using one of methods described below. Although this is not an exhaustive list of every method for conjugated polymer synthesis, this section provides a brief overview of the different methodologies for conjugated

FIGURE 3.5  Illustrative examples of narrowing the band gap of polythiophene through structural modification of the repeat unit: adding an electron rich moiety to increase the energy of the valence band (VB), shown with PEDOT, or adding an electron deficient moiety to decrease the energy level of the conduction band (CB), shown with PTP. Narrow band gap polymers can also be obtained through the donor-acceptor approach where an electron rich aryl unit is attached to that of an electron deficient one, shown with polymers 19 and 20.

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FIGURE 3.6  A general depiction of transition metal catalyzed polymerizations for repeat units functionalized with transmetalating reagents (box), halogens (dashed circle), and reactive C–H functionalities, such as alkynes or olefins (oval).

polymer synthesis during the time of publication of the last edition. More recent advancements regarding the topics below, such as transition metal catalyzed polycondensations, are provided in subsequent sections. 3.1.2.1 Transition Metal Catalyzed Polymerizations Transition metal catalyzed polycondensation reactions were a transformative addition to the synthetic toolbox for conjugated polymers, allowing for the facile and rapid preparation of many different types of these materials.[12] As shown in Figure 3.6, the general requirements for these reactions include an aryl halide with a transmetalating reagent on the same monomer, an A-B type system, or with the halide and transmetalating reagent attached to different monomers, an A-A B-B type system. The transmetalating reagent used largely defines the differences in these named reactions. Specifically, Kumada and Negishi, which use organomagnesium and organozinc transmetalating reagents, respectively, commonly use a nickel catalyst, while Migita–Kosugi–Stille and Suzuki–Miyaura, which incorporate organoboron and organostannane transmetalating reagents, respectively, commonly use a palladium catalyst. Metathesis-, Sonogashira-, and Heck-based methods are unique in that they do not require the monomer to be functionalized with a transmetalating reagent, but instead rely on high reactivity of the C–H bonds in the pendant ethynyl or olefin groups. The scope and limitations of each method have been thoroughly investigated, and these methods work with a variety of substrates. However, Migita–Kosugi–Stille and Suzuki–Miyaura polycondensations provide certain advantages in monomer stability, scalability, and ease of synthesis. Choosing the method for polymerization comes down to the accessibility of the monomers, and the toleration of certain functional groups for a given polymerization, e.g. electrophilic functionalities may not tolerate reaction conditions set for Kumada coupling. 3.1.2.2 Electrochemical Oxidative Polymerization The electrochemical preparation of conjugated polymers, typically using anodic or oxidative polymerization techniques, is a unique method relative to others described.[12a,13] This is because the polymerization does not occur through the employment of chemical regents, but through the application of a voltage across two electrodes. Some examples of the monomers used for electrochemical oxidative polymerization include polythiophenes (21) and polydioxythiophenes (9), as shown in Figure 3.7.[2,14] Comparatively, this method is limited in that the prepared materials typically lack solubility, the polymers are isolated as insoluble films on an electrode surface, and the monomer scope is narrow due to the multiple sites for coupling present in many substrates.[3c] Since the synthesized polymers are doped from electrochemical oxidation, the polymers have to be de-doped to isolate the neutral species.

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FIGURE 3.7  Synthesis of conjugated polymers 21 and 9 using electrochemical oxidation.

FIGURE 3.8  Synthesis of conjugated polymers with vinylene linked repeat units using McMurry Coupling.

FIGURE 3.9  Knoevenagel Condensation of 24 and 25 to afford the conjugated polymer 26.

3.1.2.3 McMurry Polymerization McMurry coupling is an uncommon method for the preparation of conjugated polymers, proceeding via reduction of a pendant aldehyde and subsequent homocoupling, exemplified in polymerization of a thiophene dialdehyde (22) to the thiophene vinylene polymer (23) in Figure 3.8.[15] The use of highly reactive TiCl4 and the narrow scope of possible polymers make this method less attractive.[15,16] It does however result in vinylic linkages within the conjugated polymer backbone, providing extended conjugation. 3.1.2.4 Knoevenagel Polycondensation Knoevenagel condensation is a common method for the preparation of arylene vinylene based polymers (26).[17] This method, as with the McMurry coupling, employs a bis-aldehyde (25) to which a nucleophile generated in situ will react, generating the desired polymer. The source of the nucleophile is typically an arylene acetonitrile (24), which can be deprotonated in the presence of a strong base to generate the necessary nucleophilic species. An example of this polymerization method is shown in Figure 3.9.[6f] The reaction requires polar solvents, which may impair solubility of growing polymer chains; however, polymers with Mn of 10–20 kDa have been prepared.[17a] The major limitations with this method include a potential Michael Addition side reaction where the benzylic nitrile reacts with another nitrile leading to defects embedded within the polymer, the need for polar solvents, and the relatively harsh reaction conditions that limit the breadth of monomers that can be employed. 3.1.2.5 Gilch Polymerization The Gilch polymerization, shown in Figure 3.10a, is one example of a radical/anionic polymerization used primarily for the synthesis of PPV type polymers.[18] Other methods, such as the Wessling process,

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FIGURE 3.10  Synthesis of PPV materials using Gilch Polymerization (A) and potential incorporation of defects during the polymerization (B).

incorporate different leaving groups in place of the halide (X) used for Gilch. These can include specifically: sulfonium, xanthate, dithiocarbamate, or sulfinyl groups to name a few.[19] The reaction proceeds through the generation of intermediate 28 from 27, which can then undergo a desired head to tail propagation to form 11 from 31. Conversely, the generated intermediate can undergo undesired head to head or tail to tail couplings, which will afford the defective structures illustrated by 29 and 30 in Figure 3.10b.[18d] 3.1.2.6 Wittig Type Polycondensations The Wittig and Wittig-Horner or Horner-Wadsworth-Emmons polycondensation reactions are useful methods for the preparation of conjugated polymers that incorporate a variety of repeat units, although PPV (11) is illustrated in Figure 3.11 for simplicity.[3a,b,6a,20] Like the aforementioned polymerization reactions, such as Knoevenagel and Gilch, a strong base is required to generate the necessary phosphoniumstabilized carbanion or phosphonate-stabilized carbanion from 33 or 34, respectively, which condenses with a dialdehyde (32). This requirement presents a limitation in functional group tolerance, as with Knoevenagel or Gilch polymerizations.

3.2 Advancements in Conjugated Polymer Syntheses 3.2.1 Emerging Repeat Units The aim of this section is to highlight groundbreaking work in emerging monomers that has occurred since, or was not detailed, in the last edition of the HBCP. Focus for this section is directed more towards repeat units that have emerged due to their desirable electronic or physical properties after incorporation into a conjugated polymer architecture, or manipulation of existing repeat units to achieve the same purpose. This section seeks to highlight some of the chemistry associated with each monomer discussed, since some possess streamlined syntheses and still perform very well in organic electronics, while others are closer to the opposite.

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FIGURE 3.11  (A) Wittig polycondensation between a dialdehyde and a phosphonium ylide to afford PPV. (B) Wittig-Horner or Horner-Wadsworth-Emmons polycondensation between a dialdehyde and a phosphonate.

FIGURE 3.12  Synthesis of the bis(2-thienyl)diketopyrrolopyrrole 36. Copolymerization of this unit with a variety of donor materials affords conjugated polymers with high-performance in various organic electronic applications, such as polymers 37–39.

3.2.1.1 Amide and Imide Functionalized Repeat Units Electron deficient repeat units have been intensely studied and find frequent incorporation into donoracceptor copolymers.[21] Those functionalized with amide or imide moieties, while not necessarily new, have gained significant prevalence, largely due to their tunable structures, relative ease of synthesis, and potentially useful performance in a number of organic electronic applications, such as in organic photovoltaic (OPV) or organic field effect transistor (OFET) devices.[22] The focal point of this section is not to exhaustively show every amide and imide functionalized repeat unit, but those that are commonly incorporated into conjugated polymers. Also, the synthesis of each repeat unit is shown to highlight the relative ease with which these structures can be accessed when compared to other electron deficient structures, e.g. functionalized thienothiophene as in PTB7-Th shown in Figure 3.1. Although not shown, the aryl groups described below are typically halogenated and then incorporated into a polymer synthesis using one of the transition metal catalyzed polymerizations described in the previous section, e.g. Stille or Suzuki. For further discussion regarding topics discussed in the following section, the reader is directed to Chapters 8 and 9. Shown in Figure 3.12 with the example of 36, the 2,5-d​ihydr​opyrr​olo[4​,3-c]​pyrro​lo-1,​4-dio​ne or diketopyrrolopyrrole (DPP) core has found extensive use in OPV and OFET applications.[22a,c] Although only a few examples are presented here, DPP has been incorporated into a variety of perfectly alternating donor-acceptor copolymers, semi-random copolymers, and random copolymers.[23] These polymers typically exhibit some level of semi-crystallinity. The broad absorbance characteristics of DPP

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FIGURE 3.13  Synthesis of 2-thieno[3,4-c]pyrollo-4,6-dione (TPD), 41. Copolymerization of this unit with a variety of donor materials affords conjugated polymers with high-performance in various organic electronic applications, such as polymers 42–44.

copolymers, which extends into the near-IR, and fast charge transport properties, typically attributed to its rigid, coplanar structure, have made this unit a common choice for high-performance conjugated polymers. To illustrate this point, when 36 is copolymerized with thieno[3,2-b]thiophene, affording 37, the corresponding copolymer displays a high mobility of up to 10.5 cm2V−1s−1 in OFET devices and a power conversion efficiency (PCE) of up to 8.8% when blended with PC71BM in bulk heterojunction (BHJ) OPVs.[24] Polymers 38 and 39 also perform well in OPV cells with efficiencies of 6.9% and 5.6%, respectively.[25] The DPP core is typically functionalized with aryl groups, as shown above with the thienyl groups on 36. This is due to synthetic viability, as the preparation involves condensation between dimethyl succinate and an aryl nitrile. Having the nitrile attached to an aryl group allows for a variety of different architectures to be explored, simply through modifying or changing the aryl unit the nitrile is attached to. Illustrated in Figure 3.13 thieno[3,4-c]pyrollo-4,6-dione (TPD), 41, is typically prepared in a few steps from the 3,4-thiophenedicarboxylic acid, 40, although synthesis has been reported using only a single step.[26] Unlike DPP copolymers, described above, TPD containing polymers typically possess a narrower absorption profile that ends just before the near-IR region. It has however, found great success in OPV applications largely due to high open circuit voltages (Voc) > 0.8 eV, and the ability to modify the absorption profile by incorporating different monomers, some shown with polymers 42–44 in Figure 3.13.[27] The PCEs of 5.5% and 7% for 42 and 43, respectively, along with the high hole mobility (μ h) of 44 (1.90 cm2V−1s−1) exemplifies the versatility of TPD units in high-performance conjugated polymers. Perylene diimide (PDI) 46 has emerged over the past decade as a prevalent unit for preparing n-type conjugated polymers, and the core of PDI can be prepared in a single step from the anhydride 45, shown in Figure 3.14.[22a] It is typically halogenated and then purified using repeated recrystallizations or used as a mixture of isomers. The PDI core itself has a relatively high chemical, photo, and thermal stability compared to other repeat units, finding many uses in commercial, household products. This has translated into a high air-stability for n-type semiconductor applications when incorporated into conjugated polymers, and the subsequent materials typically display ambipolar qualities. For example, polymer 47 displays an electron mobility (μe) and an Ion/Ioff of 0.06 cm2V−1s−1 and 104 when incorporated into OFET devices.[28] It has also been shown to be a proficient polymer-acceptor when incorporated into all-polymer organic solar cells, providing a PCE of 3.45% when blended with a p-type conjugated polymer.[29] The performance of 48 is diminished relative to that of 47 with μe of 2 × 10−3 cm2V−1s−1 and Ion/Ioff of 105.[30] Interestingly, 49 displays ambipolar characteristics with μe/μ h of 4 × 10−4/2 × 10−5 cm2V−1s−1 in air.[31] Isoindigo is another broadly applied amide-based electron deficient repeat unit, characterized by its relative ease of synthesis, polymers with broad absorption to the near-IR, and excellent performance in OFET and OPV applications, shown in Figure 3.15.[22e] Similar to DPP, the aryl groups attached to

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FIGURE 3.14  Synthesis of perylene diimide (PDI) in a single step, and incorporation of this material into conjugated polymers 47–49, which all display high electron mobilities (μe) and have potential uses as n-type polymers for OFETs and polymer-acceptors in all-polymer solar cells.

FIGURE 3.15  Synthesis of isoindigo via condensation of 50 and 51 followed by alkylation. Copolymerization with various donor groups affords the donor-acceptor copolymers 53–55.

the indole core can be modified to allow for changes in electronics and physical properties, such as thienoisoindigo. However, 52 can be prepared from the commercially available 50 and 51, and for the sake of brevity only polymers incorporating 52 will be discussed. A homopolymer of 52 has been prepared and the electrochromic properties show a stable, transmissive film in the reduced state with a blue-green film in the neutral state, and it was integrated into all-polymer solar cells providing a PCE of 0.5%.[32] Polymer 55 shows excellent performance in OPV applications with a PCE of 6.3%.[33] Additionally, polymers 53 and 54 show high-performance in p-type OFET devices, with values for mobilities and Ion/Ioff of 0.42 and 0.37 cm2V−1s−1 and ~107 and 106–107, respectively.[34] 3.2.1.2 Benzothiadiazole, Quinoxaline, and Analogs Aside from amides and imides, another class of electron deficient repeat units are those derived from o-phenylene diamine (56), namely benzothiadiazole (57), benzotriazole (62), and quinoxaline (63),

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FIGURE 3.16  Synthesis of benzothiadiazole (A), benzotriazole (B), and quinoxaline (C). Exemplary polymers PCDTBT and PffBT4T-2OD which make use of these units have achieved excellent performances in OPV devices.

shown in Figure 3.16.[35] Functionalization of these materials, such as halogenation, is commonly performed so as to allow for polymerization using the commonly employed transition metal catalyzed polycondensations, such as Stille or Suzuki. Although analogs of these materials may require different starting materials than o-phenylene diamine, such as compounds 58–61, a wide-range of compounds can be generated from this low-cost starting material. When these electron deficient aryl units (57–61) are incorporated into conjugated polymers, the resulting materials typically yield a semi-crystalline morphology and a narrow band-gap. Of these, benzothiadiazole (57) and its fluorinated analog (59) have exhibited excellent performance in conjugated polymers for OPV applications through the polymers PCDTBT and poly[​(5,6-​diflu​oro-2​,1,3-​benzo​t hiad​iazol​-4,7-​diyl)​-alt-(3,3ʹʹʹ-di(2-octyldodecyl)-2, 2ʹ, 5ʹ, 2ʹʹ, 5ʹʹ, 2ʹʹʹ- quaterthiophen-5,5ʹʹʹ-diyl)] (PffBT4T-2OD), with efficiencies of 7.5% and 11.7%, respectively.[36] Unlike the amide/imide-based acceptors discussed above, a unique feature of the materials shown in Figure 3.17 is the diamine-based starting material, which can be attached to a variety of aryl units and allow for the rapid synthesis of a broad chemical library of electron deficient compounds. This point is illustrated in Figure 3.17, which conveys that the diamine functionality allows for the rapid generation of

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FIGURE 3.17  Use of aromatic diamines to generate a variety of repeat units for potential use in conjugated polymer structures.

a variety of electron deficient units with various structural motifs. Examples of these monomers incorporated into polymeric structures are provided in the subsequent section. 3.2.1.3 Fused Donors While not novel or unexplored before the last edition of the HBCP, fused-ring donors, such as fluorene, carbazole, benzodithiophene (BDT), indacenodithiophene (Figure 3.18), have garnered extensive use within the last decade with many examples already illustrated and discussed.[35a,g,37] This is largely due to the high-performance they exhibit in given applications, which relates directly to the physical and electronic tunability associated with each respective structure as illustrated schematically in Figure 3.18. In general, the fused-rings allow for enhanced photophysical properties, such as improved light absorption and charge transport properties when compared to the unfused counterparts. While more fusedring donor structures exist, the ones selected offer an introduction to these materials and are some of the most widely employed for conjugated polymer synthesis. Within this subsection, most, if not all, of the polymerizations are carried out using a transition metal catalyzed polymerization, such as Stille or Suzuki, and so specifics regarding the polymer synthesis are not detailed. However, conjugated polymer structures are illustrated with some examples of incorporating the aforementioned repeat units, and a brief synthetic outline for each fused-ring donor is shown to highlight relative synthetic simplicity or complexity associated with each of the donors. Fluorene (64) is one of the simplest fused-ring donors discussed in this section, shown in Figure 3.19. Structurally, it resembles a fused biphenyl, and when alkylated it has a quaternary center as the point of fusion, depicted as compound 65. Its major benefit is the simplicity associated with its synthesis. Many building blocks associated with this compound are commercially available, allowing for the preparation of a targeted core in a single step, as illustrated in Figure 3.19.[38] Halogenated derivatives of this compound are commercially available as well, providing an avenue for the facile synthetic manipulation to obtain a desired monomer. Polymer 66 displayed high-performance in OFET devices with a hole

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FIGURE 3.18  Illustration of some common fused-ring donor repeat units found in conjugated polymers, such as fluorene, carbazole, benzodithiophene, and indacenodithiophene.

FIGURE 3.19  A typical alkylation of fluorene to yield the core structure commonly incorporated into various conjugated polymers, illustrated with examples 66 and 67.

mobility value of 0.13 cm2V−1s−1, and polymer 67 has been explored in OPV devices, achieving a PCE of 3.70%, which includes a high Voc of 1.00 V and FF of 63%.[39] Carbazole (69) is the nitrogen analogue of fluorene, and it has also been widely employed in conjugated polymer structures. Its preparation can be easily accomplished, in a few simple steps from 68, shown in Figure 3.20. This donor has found particular success with OPV applications when copolymerized with dithienylbenzothidiazole to yield the highly studied and previously discussed PCDTBT, shown in Figure 3.20.[36a] Copolymerization with the acceptor 4,10-​bis(d​iethy​lhexy​l)-th​ieno[​2ʹ,3ʹ​:5,6]​ pyrid​o[3,4​-g]thieno[3,2-c]isoquinoline-5,11-dione (TPTI) affords polymer 70, which possesses a wide bandgap that allows for a Voc of 0.96 to be obtained when incorporated into OPV devices.[40] The isomer of 69, with the halogens in the 3,6-position rather than the 2,7-position, has also found extensive use and incorporation into various polymeric structures, finding applications mostly as emitters in OLED devices.[41] The change in connectivity leads to profound changes in electronic and physical properties for the subsequent polymers, due to changes in steric hindrance and twisting along the conjugated backbone.[42] Indacenodithiophene (72) is a fused-ring donor with a relatively high level of π-conjugation.[35g] This monomer possesses greater synthetic complexity (starting from 71) than the other donors presented in

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FIGURE 3.20  Multi-step syntheses of the carbazole core structure commonly incorporated into various conjugated polymers, illustrated with examples PCDTBT and polymer 70.

FIGURE 3.21  Multi-step syntheses of the IDT core structure commonly incorporated into various conjugated polymers, illustrated with examples 73 and 74.

this section, requiring synthetic steps that are not as approachable as some of the others detailed here, e.g. fluorene or carbazole, as outlined in Figure 3.21. However, this compound has found great success in a number of conjugated polymer copolymers, yielding OPV devices with efficiencies of 6.06% and 5.97% for polymers 73 and 74, respectively.[43] A fused-donor that has found particular success is BDT and its derivatives (77–78).[37] This structure has a high-level of tunability, largely due to the number of derivatives that can be generated from the quinoidal-intermediate, 76, shown in Figure 3.22. The quinoid itself is accessible in a few relatively simple steps (starting from 75), and it can then be arylated or alkylated to provide 77 and 78, respectively. Copolymerization of these donor units with thienothiophene or TPD derivatives affords the polymers PTB7-Th or the previously described 42, respectively. PTB7-Th displays an excellent efficiency when incorporated into OPV devices with a PCE of 10.8%.[44] The high-performance exhibited by polymers incorporating BDT has made it one of the most significant and broadly employed structural motifs over the last decade. 3.2.1.4 Heteroatom Modification Modification of the heteroatoms commonly incorporated into conjugated polymers, e.g. sulfur, nitrogen, and oxygen, offers a handle for tuning the material properties both physically and electronically.[45] Specifically, substitution to heavier, larger atoms with variable oxidation states can lead to facile tuning of the photophysical or electrochemical properties, changes in bond-lengths, polarizability, and polymer packing or chain-alignment. Larger atoms and altered bonding interactions turn many of the

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FIGURE 3.22  Synthesis of benzodithiophene both with alkyl (78) and thienyl functionalities (77), and their incorporation into conjugated polymers.

heteroarenes into a non-planar architecture, unlike what is found with thiophene, pyrrole, or furan. Also, a unique opportunity exists with exploring heteroatom modification using the less commonly incorporated atoms to discover previously unknown chemical transformations and reactivity. As realms of unknown or unconventional chemistry can sometimes be met with apprehension, many of the materials described in this section have remained less prevalent relative to previous examples. While a significant amount of work has been done regarding this topic, only a few exemplary structures shown in Figure 3.23 are provided for the sake of brevity. For further discussion regarding topics discussed in the following section, the reader is directed to Chapters 7 and 12. Of the Group 13 elements, boron has garnered the most attention with the incorporation of borondipyrromethene (BODIPY, 80)-based repeat units into conjugated polymer architectures. These are highly emissive materials and have found extensive use in OLED applications.[46] However, other elements from this group have also been incorporated, such as gallium (79).[47] Moving away from carbon, Group 14 atoms have largely been well studied and extensively incorporated into conjugated polymers, specifically silicon and germanium, which are exemplified by polymers 81 and 82.[48] From Group 15, phosphorous is the most studied, particularly when arranged in various phosphole architectures with one illustrated as 83.[49] The tunability of the phosphorous, i.e. the modification of its oxidation state or the reactivity of its lone-pair electrons, allows for a variety of different materials and derivatives to be made. Quite often, the phosphorus is incorporated into a phosphole architecture with the phosphorus center either oxidized or coordinated to a metal. Oxidation of the phosphorus center will occur over time if the material is left under ambient conditions. Bismuth has also been incorporated into a heterocyclic format to yield bismole, 84. Akin to phospholes, synthesis of these materials remains challenging relative to other heterocycles due to the high reactivity of various synthetic intermediates.[50] Selenium and tellurium are some of the most well studied heteroatoms incorporated into conjugated polymers, such as 85 and 86, outside of the typically employed chalcogens oxygen and sulfur, and for a detailed discussion the reader is referred to Chapter 7.

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FIGURE 3.23  A general overview of heteroatoms that can be incorporated in place of the commonly employed sulfur, nitrogen, and oxygen.

3.2.2 New Synthetic Strategies in Conjugated Polymer Chemistry While the previous section had a primary focus on advancements regarding the monomers incorporated into conjugated polymers, i.e. more structurally focused, this section seeks to address advancements and emerging areas related specifically to conjugated polymer synthesis. The topics described herein seek to address entirely new methodologies and polymerization conditions not present in the previous HBCP edition, or improve understandings for a known polymerization method allowing for enhancements in the control of polymer products in regard to polymer structure, e.g. molecular weight or Ð. 3.2.2.1 Polymerizations via C–H Activation A major advancement within the last decade that has taken by storm the field of conjugated polymer synthesis is the preparation of polymers using transition metal catalyzed C–H activation.[51] This has been often referred to in the chemical literature as direct arylation polymerization (DArP or DAP) and direct (hetero)arylation polymerization (DHAP), with either labeling denoting the same type of chemical transformation. Figure 3.24 illustrates the synthetic advantages that C–H activation possesses over typical transition metal catalyzed cross-coupling polymerization methods, e.g. Stille or Suzuki. A detailed discussion is provided in Chapter 5, and so the reader is referred there. Briefly, C–H activation circumvents the necessity for a transmetalating reagent to be installed on the monomer allowing for a simplified and more environmentally benign synthesis of the desired conjugated polymer. Halogenation is required for most polymerization procedures, but even this can be avoided with oxidative C–H activation methods. Shown in Figure 3.25, C–H activation-based methodologies have allowed for the synthesis and preparation of regio-regular (rr) P3HT from 87, perfectly alternating copolymers (such as 92 and 95), random copolymers (such as 89), and semi-random copolymers.[52] Polymers prepared using this methodology used to suffer from donor-donor homocouplings, acceptor-acceptor homocouplings, or branching (β) defects. However, conditions have been implemented minimize defects to undetectable levels allowing for the preparation of conjugated polymers that converge or outperform their Stille or Suzuki counterparts.[40,53]

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FIGURE 3.24  A detail of the synthetic simplicity associated with C–H activation in comparison to conventional transition metal catalyzed polymerization. The fewer number of steps needed to generate monomers makes oxidative C–H activation and C–H/C–X cross couplings attractive alternatives.

FIGURE 3.25  Preparation of P3HT, a random P3HT analog (89), and the donor-acceptor polymers 92 and 95 using C–H activation-based polymerization methodologies.

An extension of the just described methodologies above, oxidative C–H arylation methods simplify the preparation of conjugated polymers even further, shown in Figure 3.26 and illustrated with polymers 97, 99, 101, and 103.[54] Specifically, halogenation is not a requirement for monomer preparation (as shown with 96. 98, 100, and 102), and instead these molecules rely on high-acidity of the C–H bonds or coordinative effects allowing for the desired C–H bond to be activated. The necessity of a chemical oxidant, such as Cu(OAc)2 or Ag2CO3, and limited substrate scope, relative to those depicted in Figure 3.25, have made this method less applicable towards the preparation of high-performing polymers for organic electronic applications. 3.2.2.2 GRIM/Chain Transfer Polymerization (CTP) Synthetic Strategies The McCullough group’s development of Grignard Metathesis (GRIM) polymerization for polythiophenes in 1999 marked a pivotal point for the field of conjugated polymers.[55] This work enabled the room temperature synthesis of P3HT with controlled degrees of polymerization and regioregularity, providing researchers with a controlled conjugated polymer “fruitfly” that has served as the basis for understanding structure-property relationships across the field. In 2004, McCullough and Yokozawa both observed chain-growth mechanistic properties in the GRIM polymerization of P3ATs when using Ni(dppp)Cl2 as the catalyst, suggesting that GRIM could be leveraged as a living polymerization method.[3f,56] Over the past decade, both McNeil and Kiriy experimentally verified that the nickel catalyst does not dissociate after reductive elimination, but instead complexes to the π-system of the growing chain and oxidatively inserts between the C–Br bond as shown in Figure 3.27.[57]

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FIGURE 3.26  Preparation of polymers 97, 99, 101, and 103 using a variety of oxidative C–H activation methods.

FIGURE 3.27  General mechanism for Chain Transfer Polymerization (CTP) with Kumada Coupling represented in the synthetic scheme. Variations for alternate coupling mechanisms are presented on the right.

Due to the complexation and transfer of the nickel catalyst in this mechanism, polymerizations of this type have been referred to as catalyst transfer polymerization (CTP), with GRIM referred to as Kumada catalyst transfer polymerization (KCTP) because of its basis around Kumada coupling. Since these discoveries, CTP has been explored using a host of traditional polymerization methods as shown in Figure 3.27, and has allowed for the exploration of properties associated with the living growth of conjugated polymers. These include the synthesis of block copolymers via sequential monomer addition, polymer brushes, graft polymers, branched polymers, and end-capped polymers.[58] The reader is referred to Chapter 6 for a comprehensive discussion on CTP methodology and applications. Limitations of CTP include breadth of monomer scope, with the bulk of the work limited to electron-rich monomers, allowing for CTP of polythiophenes, polyselenophenes, polytellurophenes, polypyrroles, and poly(p-phenylenes) as shown in Figure 3.28.[59] CTP reactions are highly dependent on monomer-catalyst interactions, and electron-rich monomers tend to complex with the preferred nickel diamine catalyst to a greater extent than electron-poor monomers. This makes alternating copolymers a challenge, but some can be synthesized by CTP through asymmetric functionalization of the presynthesized repeat unit, as was done with the polymer P(3AT-alt-PP) shown in Figure 3.28, as well as furan-thiophene copolymers[60] This catalyst dependence also affects the synthesis of block copolymers, where monomer addition must follow in order of increasing catalyst complexation strength.[61] For this reason, the block copolymer PPP-b-P3AT, shown in Figure 3.28, could not be synthesized as P3AT-bPPP. Another concern in CTP is unintended bidirectional growth caused by random catalyst walking. Random catalyst walking was observed by Kiriy et al. and allows for the nickel catalyst to “walk” across the conjugated backbone to insert on the opposite end of the chain.[57b]

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FIGURE 3.28  Examples of polymers synthesized via CTP methods.

3.2.2.3 Continuous Flow Synthesis Continuous flow methodologies have allowed for the large-scale preparation of conjugated polymers addressing issues regarding safety, scalability, and reproducibility that tend to plague conjugated polymer synthetic methods, and a detailed discussion of this is provided in Chapter 16.[62] While an oversimplified explanation, continuous flow conjugated polymer synthesis, in general, is performed by pumping the reaction mixture (monomers, solvent, catalysts, additives) into a temperature and pressure controlled flow reactor, then through a back-pressure regulator (BPR), and extruding it and precipitating the polymer product. Setups vary depending type of flow reactor needed, and reaction mixtures can be separated into different streams and mixed only before entering the flow reactor. This methodology allows for the multi-gram (>100 g) preparation of conjugated polymers in the common organic chemistry laboratory and is further scalable for serving industrial applications. This methodology has even been applied to C–H activation-based polymerizations. However, only a handful of polymers have been prepared in this fashion with some examples (P3HT and polymers 104–107) shown in Figure 3.29. 3.2.2.4 Click-Chemistry and Multi-Component Reactions Click-chemistry has introduced new synthetic transformations to field of organic synthesis, and its application to conjugated polymer synthesis was inevitable.[63] By definition, click-chemistry allows for the rapid synthesis of desired compounds providing high yields without toxic byproducts or extensive work-up and purification steps. Of the various click-reactions described, the copper-catalyzed azidealkyne cycloaddition (CuAAC), shown in Figure 3.30A, is one of the most identifiable. This reaction occurs between alkyne and azide functionalized monomers, with a copper catalyst to help activate the alkyne moiety by enhancing its electrophilicity. This methodology was applied towards the synthesis of fluorene, phenylene, and benzothiadiazole containing copolymers, shown in Figure 3.30 as polymers 108 and 109.[64] The polymerization itself was carried out at or near room temperature in tetrahydrofuran affording high yields (>90%), albeit with modest molecular weights (Mn = 6–8.7 kDa). Another click-reaction is the thiol-yne reaction, which occurs between a thiol and an alkyne through a radicalmediated mechanism, shown in Figure 3.30B with polymer 110.[65] This methodology affords high-yielding (>95%) materials with high molecular weights (Mw = 21–61 kDa).[66] However, the radical-mechanism affords cis-trans isomers in a roughly 1:1 ratio, which may hinder the desired electronic or physical characteristics of the polymer. Another methodology akin to click-chemistry includes multi-component reactions. These reactions may not possess all of the highly regarded merits that would qualify them to be click-reactions, e.g. environmentally benign conditions, simple-purifications, and high yields, but the potential of these reactions to construct conjugated polymers from new building blocks opens an exciting door leading to a new realm of conjugated polymer chemistry and structure.[67] Also, this methodology seeks to incorporate simpler building blocks, allowing for streamlined synthetic pathways, and ultimately trimming down the amount of waste and time associated with a lengthy monomer synthesis. Shown

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FIGURE 3.29  Simplified depiction of a continuous flow reactor and examples of conjugated polymers (P3HT and 104–107) that have been successfully synthesized using continuous flow methods.

in Figure 3.30C with polymers 111 and 112, Arndtsen et al. prepared fluorene and pyrrole copolymers with different π-spacers, such as thiophene and naphthalene.[67c] The polymerization proceeds through the preparation of poly(1,3-dipoles), which then undergoes cycloaddition reactions with an alkyne. This procedure allowed for the preparation of polymers with yields up to 98% and Mn up to 40.7 kDa. 3.2.2.5 Molecular Weight and Dispersity Effects Lack of control in terms of molecular weight and dispersity made the early development of structureproperty relationships in conjugated polymers challenging, with materials having the same nominal polymer repeat unit showing orders of magnitude differences in performance in organic electronic device testing. The synthetic developments that provided increased control over molecular weight and dispersity in conjugated polymers, led by McCullough’s synthesis of regioregular P3HT, has since enabled researchers to study the effects of these variables on physical, morphological and optoelectronic properties. The most noticeable physical properties affected by MW are solubility and aggregation. High MW polymers tend to aggregate, which can lead to lower solubility in organic solvents. Increased aggregation can also be noted in the UV-vis spectra of conjugated polymers, in which higher MW polymers are often red-shifted to lower bandgaps than their lower MW counterparts. While the extent of aggregation may increase with molecular weight, the absorption profile of conjugated polymers generally remains unchanged. Slight red shifting can be observed, but typically the changes saturate at the effective conjugation length, which happens for many conjugated polymers at around ~10–20kDa.[68] For redox applications like electrochromism and charge storage, electronic properties also saturate around this MW. The most pronounced effects of MW in conjugated polymers are observed for morphology and charge mobility. An early example from the Frechet group found that mobility values in P3HT-based OFETs increased from 1.7×10−6 to 9.4×10−3 cm2V−1s−1 as the MW was increased from 3.2 to 36.5 kDa.[69] The authors hypothesized that higher crystallinity would boost charge transport, but found that the higher

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FIGURE 3.30  Synthesis of conjugated polymers using click-chemistry (A) and (B) or multicomponent polymerization (C) strategies.

crystallinity, low MW polymers exhibited lower mobility than the more amorphous high MW polymers. The low MW polymers formed well-defined crystallites in thin films, but the crystalline domains followed no preferred orientation, leading to a large amount of grain boundaries that could serve as charge traps. Brinkmann et al. subsequently determined that low MW P3HT essentially exists in a “fully extended” rigid rod state, while high MW P3HT contains both “fully extended” domains and amorphous domains that contain chain folds, chain ends, and tie-molecules.[70] The general hypothesis was that longer polymer chains could carry a charge further through a film before necessitating a hopping event to another chain, and that limiting the occurrence of hopping events increased charge mobility. This hypothesis has been reaffirmed as newer conjugated polymers, notably more disordered polymers, have emerged and followed the same trend of increased mobility with increased molecular weight. In a comprehensive study by Noreiga et al., the authors put forward the hypothesis that high mobility in conjugated polymers is obtained when a series of crystalline, aggregate domains are interconnected via long tie-molecules, as portrayed in Figure 3.31C[71] By this logic, high MW is essential to achieve this optimal morphology, and therefore to achieve high charge mobility. Building on this hypothesis, the Salleo group sought to understand whether this optimal morphology could be obtained by mixing crystalline low MW P3HT with a small amount of high MW P3HT to serve as tie-molecules.[72] Using a combination of computational modeling, XRD data, and OFET device statistics, the authors instead found that the presence of any amount of low MW P3HT (~8 kDa) blended with high MW P3HT (~61 kDa) was detrimental to the charge mobility. However, a medium MW P3HT (~29 kDa) could be blended with the high MW polymer in any ratio without affecting the

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FIGURE 3.31  Schematic representation of polymer microstructure in highly ordered films (A), amorphous films (B), and disordered films with tie-chains connecting local aggregates (C) with a magnification showing local aggregate of three polymer chains.

charge mobility values. The authors hypothesize that polymers below their disorder threshold, the MW in which changes in crystallinity and charge mobility with respect to MW begin to plateau, can act as low mobility charge traps even when connected by longer tie-molecules. This work emphasizes the importance of not only having high MW, but also of having a narrow dispersity in which the low MW fractions are still above the disorder threshold. Molecular weight effects are also seen in OPV devices of conjugated polymer and fullerene blends.[68] While many of the effects are caused by the aforementioned impact on morphology and charge mobility, the relationships become significantly more complex as variables like domain size, purity, and orientation are affected by changes in MW. Unlike neat polymers for OFETs, donor polymers for OPV applications show peak performance at an optimal MW, typically 30–60 kDa, before encountering processing issues at very high MW due to decreased solubility and improper BHJ formation when mixing the donor and acceptor phase materials.

3.2.3 Structure Property Modification of Conjugated Polymers The prior sections have focused on emerging areas in the preparation of new conjugated polymers, both with regard to new structural units and new methods of polymerization. With these advancements in the synthesis has come an enhanced ability for polymer chemists to design conjugated polymers for specific properties, functionalities, and applications. This section serves to highlight several of the ways a more highly developed synthetic toolbox has been created. 3.2.3.1 Random and Block Copolymers Synthetic advances in carbon-carbon cross-coupling polymerizations have effectively shifted the focus of the conjugated polymers community from homopolymers towards alternating copolymers. Many modern conjugated polymers are perfectly alternating, symmetric copolymers comprised of an electron-rich monomer and an electron deficient monomer to achieve an internal donor-acceptor interaction. The low bandgaps and favorable charge transport properties provided by these perfectly alternating donor-acceptor copolymers have made this design the dominant motif for semiconducting conjugated polymers, but recently researchers are exploring random, semi-random, and block copolymer structures to achieve unique properties. An early motivation for moving away from perfectly alternating copolymers was the need for broadening visible light absorption. With random and semi-random conjugated copolymers, variations in the sequence of monomers along the backbone leads to a variety of unique chromophore combinations that contribute subtly different absorption spectra, effectively broadening the overall polymer absorption profile. This method has been used within the Reynolds group to develop broadly absorbing black electrochromic polymers from a random copolymer structure of BTD and ProDOT.[73] To synthesize random conjugated polymers, chemists have utilized a dibrominated and distannylated monomer, A, along with a dibrominated and distannylated monomer, B, and vary the ratios in a Stille polymerization as shown in Figure 3.32A.[74] Alternative coupling methods, such as DHAP, can also be used to polymerize

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Br

D

Br

Br

A

Br

D

A.

+ Sn

D

Sn

Sn

D

Br

Random Sn

A

Sn

D B.

+

A

Sn

A

Br

A

Random Regiospecific

Br

A

Sn

Br

D

Sn

D

A

+

C.

Semi-Random Sn

D

Br

Regiospecific Donor-Acceptor Sequence Control

FIGURE 3.32  General synthetic methods for Stille polymerization of random conjugated copolymers (A), random conjugated copolymers with regiospecificity (B), and semi-random conjugated copolymers (C).

random copolymers.[75] These methods were successful in broadening polymer absorbance, but lacked control over sequence or regioregularity. One workaround involves synthesizing monomers with both bromide and stannane functionalities, seen in Figure 3.32B, providing the absorbance broadening effect of random copolymers but with a designed regiospecificity to ensure head-to-tail coupling.[76] However, the synthesis of these asymmetric monomers can be difficult, especially for acceptors, and represents a major barrier to an optimization of these polymers. To provide both regiospecificity and control of donor-acceptor sequence to random copolymers, the Thompson group polymerized an asymmetric alkyl thiophene monomer, a distannyl thiophene, and a dibromo BTD, leading to a semi-random copolymer as shown in Figure 3.32C.[77] By preserving the head-to-tail coupling of alkylthiophenes and preventing coupling of acceptor monomers, these semi-random copolymers achieve broadened absorbance of visible light without strongly impacting the charge mobility or crystallinity of regioregular P3HT. The ability to use a dibromo acceptor monomer also expands the toolbox of monomers that can be used for random conjugated polymers. Random and semi-random copolymers can also provide a synthetic handle over frontier energy levels through changes to comonomer ratios.[78] Aside from optoelectronic properties, random copolymers can embed other materials properties through the incorporation of functional monomers. The Thompson group was able to tune the surface energy of polythiophene films through the incorporation of thiophene monomers with oligoether or fluorinated side chains, while work from the Mei group incorporates n-propyl comonomers to break conjugation and enable meltprocessable and stretchable conjugated polymers.[79] These examples demonstrate the possibilities for functional conjugated polymers through the use of random and semi-random conjugated polymers. Conjugated block copolymers have also received significant interest in the past decade, primarily with the goal of achieving the morphological control found in traditional block copolymers. The two main synthetic strategies for creating conjugated block copolymers are polycondensation reactions of endgroup functionalized prepolymers and pseudo-living polymerizations using CTP methods as shown in Figure 3.33. For the end-group polycondensation methods, one block is grown through traditional Stille or Suzuki polycondensation reactions, and the growing chains are subsequently end-capped with

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Past Ten Years’ Advancements in Conjugated Polymers C6H13

A. Br

Br

D

Br

+

Sn

X

D

A

Sn

S S 113

Stille Br

A

D

H m C6H13

C6H13

D

Sn n

A

S n

m

B. Br

CTP

X

D

CTP

Br n X

A

D

A n

m

Br

FIGURE 3.33  General synthetic methods for synthesizing conjugated block copolymers using end-functionalized prepolymer blocks (A) and through pseudo-living chain transfer polymerization (CTP) methods (B).

another polymer block with a complementary reactive end-group (e.g. a bromide on the second block to react with a stannane on the growing first block in Stille polycondensation).[80] The key advantage of this method is the ability to work with a wide selection of monomers that are currently incompatible with CTP methods, particularly electron deficient acceptor monomers prevalent in donor-acceptor conjugated copolymers. However, the reliance on polycondensation-type polymerizations leads to higher dispersity polymers along with homopolymer side products. Moreover, a significant amount of the work around this end-group approach makes use of a bromide end-functionalized polythiophene (113) as the second block, which ultimately relies on CTP methods. Alternatively, CTP provides fine molecular weight control to afford conjugated block copolymers with narrow dispersities and minimal homopolymer impurities, as further detailed in Chapter 10.[81] With the pseudo-living characteristics of CTP, a polymer block can be grown in a controlled manner, and a change in the monomer feedstock continues the growth of the second block. However, unlike polycondensation reactions, the success of CTP is highly dependent on monomer-catalyst interactions, limiting the scope predominantly to thiophenes, selenophenes, pyrroles, fluorenes and poly(p-phenylenes).[59] But this limitation can most likely be overcome through catalyst optimization rather than serving as an inherent disadvantage to CTP. For example, the Seferos group reported the synthesis of a polythiophene-block-benzotriazole polymer enabled through the use of a nickel diamine catalyst.[82] Ongoing work is focused on catalysts with less selective association energies to broaden the scope of CTP and enable the synthesis of donor-acceptor block copolymers. Overall, conjugated block copolymers have found limited success in organic electronics applications. The main area of interest is within OPV active layers, where a single conjugated block copolymer could effectively serve as both the donor and acceptor and can be processed to adopt a lamellar morphology favorable for charge separation. While some successes show increased performance from block copolymers compared to blends of the same homopolymers, most systems still obtain PCE values below 5% when state-of-the-art OPVs are reproducibly performing at over 10% PCE.[83] These shortcomings are likely caused by the large degree of phase separation between blocks, which do not provide the optimal morphology found in OPV active layers using polymeric donors and molecular acceptors where both pure and mixed domains play important roles in charge separation. A possible remedy to this morphological problem is the use of gradient conjugated copolymers. Largely developed by the McNeil group, gradient conjugated copolymers are synthesized through gradually changing the monomer feedstock in CTP to progress from a pure first block, to an alternating or random block of the first and second monomer, to a pure second block.[84] Through this technique, conjugated polymers can be made to exhibit an intermediate morphology that is more ordered than blends of homopolymers but not as phase segregated as block copolymers. Moving away from homopolymers and perfectly alternating copolymers towards random, semi-random, and block copolymers has enabled unique properties in the field of conjugated polymers, including absorption broadening, energy level modification, and morphological control. While these classes

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of polymers can increase disorder in the solid-state and interrupt π-π packing, there are still many opportunities for these systems to outperform perfectly alternating copolymers. Specifically, the ability to incorporate functionalities that increase solubility, enable surface energy control, and provide unique materials processing properties like fusibility and elasticity are well-deserving of future research attention. 3.2.3.2 Side Chain Engineering Conjugated polymer design has steadily progressed from insoluble materials electropolymerized from a limited choice of monomers toward solution-processable materials with broad synthetic versatility. Much of this progression in the 2000s focused on the attachment of side chains, typically aliphatic in structure, onto the conjugated backbone. This modification not only afforded a final polymer that dissolved in organic solvents but provided a synthetic handle over molecular weight and dispersity via solution-based polymerization methods. Side chains expanded the toolbox for conjugated polymer chemists and led to new families of soluble polymers based on the classical systems and new aromatic monomers that were inaccessible via electropolymerization. While much of the initial intent of side chains was to provide solubility, over the past decade they have been utilized and explored for their influence on solid-state structure and morphology, backbone torsion and electron density, and reactivity that provides unique properties that go beyond optoelectronic properties as summarized in Figure 3.34. Side chains have been known for some time to affect solid-state structure, with evidence of interdigitation of long, aliphatic side chains promoting increased crystallinity in conjugated polymer thin films, notably PBTTT.[85] However as the field moved towards branched aliphatic side chains in efforts to further increase solubility, it became apparent that there is a complex relationship between solubility and solid-state structure. Replacing linear side chains with branched alternatives disrupts the ability to interdigitate, and while this greatly increases the solubility of the polymer, it also prevents the formation of ordered domains in the solid-state. A recent example in the Reynolds group shows an

Solubility Modifiers

O

O

O

Torsion & Electron Density O K

O

S S O

S

S

S

OH

C6H13

C6H13 S

O base

S

S

C6H13

C6H13

S

S

S

S

S

S

S

S S

C6H13 S

O

S

S C6H13

S

C6H13

Reactivity

S

S

Morphology & Solid State Packing

C6H13

O

O

FIGURE 3.34  Summary of the conjugated polymer properties influenced through side chain engineering, including solubility modifiers, backbone torsion and electron density, morphology and solid-state packing, and reactivity.

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isoindigo-terthiophene polymer that achieves ~5.1% in OPV devices when using linear hexyl chains but drops to ~0.3% when using 2-ethylhexyl branched chains due to the total disruption of solid-state order.[86] Despite the tendency to disrupt order, branched side chains are still ubiquitous in modern conjugated polymer design due to their ability to solubilize flat, conjugated molecular units, and their effect on solid-state packing can be minimized by pushing the branching point further away from the conjugated backbone.[87] One consideration with side chains is that their effects, beyond solubility, can be difficult to predict and are often dependent on both repeat unit structure and intended application. Copolymers containing three or four side chain sites per repeat unit are now commonplace, and each site adds another variable to contribute to the overall complexity. A notable example of the complex nature of side chain effects is the hypothesis put forward by Graham et al. that in donor-acceptor copolymers for use in PCBM blend OPVs, having bulky side chains on the donor moiety and linear chains on the acceptor moiety yield the optimum device metrics.[88] The idea is that the bulky side chains direct PCBM molecules towards the less sterically hindered acceptor moiety, thereby improving electronic interactions and facilitating charge transfer. While this hypothesis holds true for a handful of systems, there are certainly outliers, emphasizing the complex nature of structure-property relationships of side chains on conjugated polymers. Going forward, it will be important for polymer chemists to perform side chain optimization studies on new conjugated backbones and observe the effects on solid-state structure through X-ray scattering techniques like GIWAXS, GISAXS, and RSOXS. Through continued efforts in this area, the field can begin to build a deeper understanding of side chain effects on solid-state structure and move away from trial and error type research. Aside from morphology, side chains can also affect the backbone structure by donating or withdrawing electron density and altering inter-ring torsion angles. The Reynolds group has used these types side chain effects to fine tune color in electrochromic dioxythiophene copolymers.[89] Moving from the small dioxy ring in 3,4-ethylenedioxythiophene (EDOT) to a larger ring of 3,4-propylenedioxythiphene (ProDOT) increases torsional strain along the backbone and leads to a decrease in planarity and extent of conjugation, leading to higher gap systems.[90] Even more backbone strain can be introduced upon moving to the acyclic monomer, 3,4-diethylhexyoxythiophene (AcDOT). These three monomer units also vary in their ability to donate electron density into the conjugated backbone. By combining these monomer units into a variety of copolymer structures, effectively altering variables of torsion and electron density, the Reynolds group has been able to create a complete color palette of electrochromic polymers. In this manner, side chains have been widely used to modify planarity and electron density to control both the absorption profile and electronic bandgap. Side chains have been used to append additional functionality into conjugated polymer designs. Common types of functionality include crosslinkability, solubility modifiers, and side chain cleavage. For crosslinking and substrate binding, azide-terminated side chains can be incorporated to react via click-chemistry,[91] vinylenes can be thermally initiated to crosslink,[92] bromides,[93] oxetanes,[94] and acrylates[95] can offer UV crosslinking properties, and Diels–Alder units can participate in reversible crosslinking mechanisms.[96] Crosslinking conjugated polymers has been useful for achieving robust, solvent resistant films, allowing for facile orthogonal processing steps, binding the substrates and analytes, and can even be leveraged to control morphology. For solubility modifiers, oligoether side chains can be used to impart solubility in polar solvents, and have recently proved to be promising candidates for organic electrochemical transistors (OECTs) where ionic interactions play a major role.[97] Beyond solubility, oligoether chains can impart a sensing mechanism into conjugated polymers, and Hayward et al. showed that a polythiophene with oligoether side chains could be complexed with potassium to form superhelical structures.[98] Side chains terminated with ionic groups have also been used to achieve solubility in polar solvents, notably water. These conjugated polyelectrolytes (CPEs) have seen usage across both biological applications and as interlayers and workfunction modifiers in organic electronics devices.[99] Sulfonates and Carboxylates are common for anionic CPEs, and pyridinium and quaternary amines are common for cationic CPEs.

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Conjugated Polymers: Perspective, Theory, and New Materials

Ester side chains make for interesting precursors to carboxylate CPEs, as they can be polymerized via traditional solution methods using long aliphatic chains connected to the ester, and then hydrolyzed to afford the water-soluble conjugated polycarboxylate.[100] After coating thin films, the polycarboxylate can subsequently be introduced to an acidic environment to create an insoluble conjugated polycarboxylic acid. Alternatively, the Frechet and Krebs group have shown multiple examples of solid-state thermal cleavage of ester side chains through decarboxylation.[101] These polymers could be coated from solution and heated to 200°C to cleave the esters, resulting in solvent resistant conjugated polymers with pendant carboxylic acids that exhibited up to a 0.5 eV decrease in bandgap compared to their noncleaved counterparts. Ester side chains can also be oriented in such a way that hydrolysis leads to cleavage of the carboxylate, leaving behind an insoluble, alcohol functionalized conjugated polymer backbone. Using this design, polymers have been designed that can be coated from organic solution and subsequently dipped into a basic solution to remove the ester side chains.[102] Side chain cleavage provides the ability to make use of side chains for the solubilizing properties, but then remove them in the solid-state to increase the density of conjugated backbone within thin films. Alternative approaches use silyl side chains that can be cleaved with strong acids, and o-nitrobenzyl units that can be cleaved via UV light and offer the ability to photopattern conjugated polymers.[103] A recent approach combined multiple side chain functionalities, esters, and o-nitrobenzyl units to provide conjugated polymers that can be synthesized and characterized in organic solvents, hydrolyzed to a polycarboxylate that can be processed from water, and then irradiated with UV light to cleave the side chains and leave behind an insoluble conjugated polymer.[104] These examples clearly demonstrate the potential for side chains to impact more than just solubility, and over the last decade, researchers are just beginning to scratch the surface of the possibilities that exist within functional side chains. 3.2.3.3 n-Type Conjugated Polymers Prior to the last decade, the vast majority of conjugated polymers were p-type materials used in applications requiring hole transport materials and oxidative doping. However, to fully realize conjugated polymer electronics, there is a need for electron transporting, n-type conjugated polymers. Specifically, many electronic applications require complimentary hole and electron transporting materials, such as donor and acceptors in solar cells, n-channel and p-channel OFETs for integrated circuits, asymmetric supercapacitor cells for larger voltage windows, and n-type and p-type materials for thermoelectrics. Progress has been made in the past decade once naphthalene diimides (NDI) and perylene diimides (PDI), known to function as n-type organic materials, began to be incorporated into conjugated polymer backbones. One of the first examples came from the Marder group where an alkylated PDI was polymerized with dithienothiophene, 114 in Figure 3.35, and the resulting polymer exclusively operated in n-channel OFET devices, reaching electron mobility on the order of 10−2 cm2V−1s−1.[105] Soon after the Facchetti group reported a polymer consisting of an alkylated NDI monomer and 2,2ʹ-bithiophene (115) that achieved electron mobility values up to 0.85cm2V−1s−1 in n-channel OFETs.[106] The authors emphasized the importance of ambient stability for organic electronics applications, and this polymer, P(NDI2OD-T2), was both deposited and tested under ambient conditions. P(NDI2OD-T2) has since become the state-ofthe-art n-type conjugated polymer and its general backbone structure has seen usage in n-type OFETs, as acceptors in all polymer OPVs,[107] n-type thermoelectrics,[108] aqueous compatible n-type OECTs,[109] and as a charge storage material for rechargeable lithium batteries[110]. Another significant property of P(NDI2OD-T2) is its ability to be electrochemically n-doped. Because NDI is a reversible two-electron acceptor, Liang et al. classify P(NDI2OD-T2) as a π-conjugated redox polymer, a polymer consisting of both active redox sites along the backbone to accept electrons and a conjugated backbone to transport electrons.[110] Using in situ electrochemical doping experiments, the authors found that P(NDI2OD-T2) could be reversibly n-doped to a conductivity of 10−3 S cm−1, five orders of magnitude higher than its pristine conductivity.

135

Past Ten Years’ Advancements in Conjugated Polymers C10H21 C12H2 5 O

N

C14H29

C10H21

O C8H17 O S

C14H29

C10H21 N

O

S

S

N

S

F

C8H17

N

N

O

O

F

S O O O

N

O C10H21

N

O C10H21

N

C8H17

C8H17 115

S

116

F

S N

S F

C10H21

C12H25 114

O

S

N

F

F

C14H29 C14H29

117

FIGURE 3.35  Examples of n-type conjugated polymers.

Aside from PDI and NDI based polymers, recent works have focused on alternate electron deficient monomers, such as DPP and isoindigo, to create new n-type conjugated polymers. These polymers are designed to facilitate stabilization and delocalization of the LUMO associated with the π-conjugated backbone. LUMO stabilization can be achieved by incorporating electron deficient monomers with high electron affinity, and delocalization can be promoted through planarization of the conjugated backbone along with efficient π-π stacking in the solid-state.[111] With this design, the Li group has used a polymer consisting of electron deficient DPP and pyridine monomers (116) to produce n-channel OFETs with mobility values reaching 6.3 cm2V−1s−1.[112] Pyridine was chosen as a comonomer because of its higher electron affinity compared to thiophene as well as its ability to planarize with the DPP unit. Fluorination of conjugated polymers is also an effective method of stabilizing the LUMO, and has been used in both DPP and isoindigo polymers to achieve electron mobilities of ~2 cm2V−1s−1 and ~4 cm2V−1s−1, respectively.[113] An example of the highest performing n-type polymer of the study, 117, is shown in Figure 3.35. While the development of n-type conjugated polymers has lagged behind that of p-type, the field has continued to stride towards higher electron mobility, and materials are approaching mobilities useful for organic electronic applications. The future for n-type conjugated polymers will need to focus on obtaining higher electron mobility with scalable, environmentally tolerant materials. For a comprehensive overview of n-type conjugated polymers, the reader is referred to Chapter 9. 3.2.3.4 Metallopolymers Metallopolymers are an important class of conjugated polymers, offering material and electronic properties that only transition metals can provide, e.g. spin-orbit coupling, advanced architectures through metal-coordination, and magnetism, enhancing or offering new characteristics to conjugated polymers for organic electronic applications.[114] While there are many examples and various structural motifs for metal-containing polymers, the ones shown in Figure 3.36 were selected for the sake of brevity, but a more detailed discussion is offered in Chapter 11. From Figure 3.36A with polymers 118 and 119, the platinum-acetylide polymers show that copolymers can be prepared from an organic repeat unit and a transition metal complex via Hagihara coupling.[114g,h,j] These materials have been incorporated extensively into OPV and chemical sensing applications. However, diminished bonding interactions between the pπ and dπ orbitals of the conjugated, organic π-network and the metal center can inhibit overall conjugation of the π-system. Coordination polymers, such as 121 in Figure 3.36B, allow for the preparation of supramolecular structures via the addition of a ligand, such as bipyridine, which coordinates to zinc metal centers between polymer chains (120).[115] This has strong implications for the polymer structure and allows for improved charge transport properties. Polymer 121 best illustrates coordination polymers where the metal center is within the monomer but is not a member of the conjugated backbone, as with polymers 118 and 119. Another architecture is illustrated in Figure 3.36C with polymers 122 and 123, where a nickel metal center is coordinated to a salen-based ligand, showing that the coordinating ligands do not necessarily have to be covalent bonds along the polymer backbone, such as with the

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Conjugated Polymers: Perspective, Theory, and New Materials

FIGURE 3.36  (A) Hagihara coupling to afford platinum-acetylide metallopolymers. (B) Glaser coupling affords zinc-porphyrin coordination polymers. (C) Metal-salen coordination polymers.

platinum-acetylide polymers, or complex architectures as with the porphyrin. These materials exhibit interesting electrochemical properties that vary due to the steric bulk present on the polymer chain.[114a] 3.2.3.5 Conjugated Porous Polymers Another emerging area in regard to polymer structure and synthesis is that of conjugated porous polymers (CPP), shown in Figure 3.37. These materials have been incorporated into applications that cover catalysis, gas-adsorption, and small-molecule detection.[46b,116] In general, their synthesis is reliant on the use of a monomer that possesses multiple sites for reactivity to allow for an extended, multi-dimensional π-conjugated network to be produced. The pore sizes can be tuned in a decisive manner through implementing different monomers or by changing the connectivity between monomeric units, e.g. placing an ethynyl linkage between monomers as with 126. The extensive conjugated network allows for absorbance of wavelengths of light in the visible range, allowing them to be used for an extensive amount of photochemical applications. Most often these materials are implemented for use as photosensitizers, such as with photodynamic therapies, or as photocatalysts, such as with various oxidative or coupling reactions.[117] The heterogeneous nature of these materials allows for improved recyclability of the catalyst and improved chemical and thermal stability. These materials can be prepared using more common transition metal catalyzed polymerizations, such as Sonogashira (starting from 124 and 125), shown in Figure 3.37A with 126, or with more classical techniques, such as chemical oxidation with ferric chloride (starting from 127), shown in Figure 3.37B with 128.[116a,g] In one study, C–H activation was implemented to allow for the copolymerization of 129 and 130 to afford the CPP 132.[116h]

Past Ten Years’ Advancements in Conjugated Polymers

137

FIGURE 3.37  Porous conjugated polymers prepared using Sonogashira (A), oxidative polymerization (B), and C–H activation (C).

3.3 Future Direction and Outlook As this chapter has summarized where the chemistry of conjugated polymers left off ten years ago and has highlighted notable advancements made over the past decade, this section will serve as a perspective of where the field will likely be moving toward and what questions are left to be addressed. Specific attention will be given to how conjugated polymers can be made in more efficient and safe methods, and to how these polymers can be designed to best adapt to future trends in both commodity and niche application spaces.

3.3.1 Efficient Monomer and Polymer Synthesis While Section 3.2 of this chapter already covered the emergence of new monomer units and the importance of C–H activation chemistry in DArP, it is important to note the impact of direct arylation and other C–H activation chemistry on the synthesis of monomeric units. While DArP curbs the need for toxic stannylated monomers, many emerging monomers are multi-ring and heterocycle units that make use of stannanes in their production. Since the industry is generally averse to scaling up organotin

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Conjugated Polymers: Perspective, Theory, and New Materials

reactions, requiring the production, storage, and use of stannanes within any stage of the production of conjugated polymers represents a serious barrier to industry adoption. Going forward, it will be important for polymer chemists to employ new methods in C–H activation chemistry and strive to increase the efficiency and safety of their current and future monomer syntheses.[118] Aside from reducing the use of stannylated compounds, polymer chemists must also keep in mind the total complexity of their polymer syntheses. In their 2015 perspective, Po et al. provide a thorough analysis of state-of-the-art conjugated polymers for OPV active layers in terms of their synthetic complexity.[119] Defining synthetic complexity as a combination of synthetic steps, reaction yields, number of work-up steps (with an emphasis on column chromatography), and safety of the chemicals used, the authors bring to light the impact of multi-step syntheses on potential polymer scale-up, and put forward the warning that small gains in PCE are not always worth the increase in synthetic difficulty. In a field that often cites scalability and cost reduction through large-format roll-to-roll processing techniques as core advantages over alternate conducting and semiconducting materials, the idea of reducing synthetic complexity should be on the forefront of conjugated polymer research in the next decade. Another challenge that will need to be further addressed is the consistency of conjugated polymer synthesis. While monomers can be synthesized with higher yields and safer chemicals, polymerization creates specific challenges in terms of molecular weight and dispersity that can often lead to varied, unpredictable results when performed at different scales. There are already remedies for batch-to-batch variation in the literature mentioned previously in this chapter, including CTP synthetic methods to achieve predictable molecular weights with narrow dispersity and the use of continuous flow polymerization, but the scope of these efforts is currently limited. Flow polymerization has been predominantly explored through polycondensation reactions like Stille, which bring about safety concerns as previously mentioned. Direct arylation and CTP have been explored in flow polymerization setups but still need further research and optimization to become reliable methods. The scope of both direct arylation and CTP is also limited in conjugated polymer synthesis, but the toolbox of usable monomers has steadily increased with ongoing research in new catalysts and reaction conditions. The ability to harness these emerging polymerization methods and scale them up within a continuous flow reactor should continue to be a major goal of the conjugated polymer community. By making conjugated polymer synthesis efficient, scalable, and safe, the barriers to industry adoption can continue to fall over the next ten years.

3.3.2 Polymer Properties and Applications While conjugated polymer properties and applications are presented in greater detail in Volume 2 of this book, it is essential for polymer chemists to understand major trends in the field and how synthetic design can be used to assist in these efforts. Though there are a multitude of possibilities for the future of conjugated polymers, this section will highlight low-cost printed electronics and conjugated polymers for biological applications as major areas to look towards in the next decade of research. The core advantages of using electroactive polymers over inorganic materials are solution processability and the ability to fine tune properties through synthetic design. Coupled with well-developed industrial printing equipment, these qualities lend themselves to scalability and customizability, which can have significant implications if realized. Simple integrated circuits can be mass printed at low cost using polymer OFETs, benefiting the growing Internet of Things (IoT) market that seeks to connect the physical world to the digital through a myriad of low-power sensors. Electrochromic polymers can be printed onto large, irregular surfaces for next-generation low-power displays. Semi-transparent polymer OPVs can be printed onto glass windows to provide a source of energy for buildings without taking up significant space.[120] The ability to print conjugated polymers from solution has always been one of their key advantages, but work needs to be done to make conjugated polymers adapt to the printing industry. One major hurdle is the choice of solvent typically used to dissolve conjugated polymers. Chloroform and

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139

chlorobenzene are good solvents for dissolving standard alkylated conjugated polymers, but they are not good solvents for industrial printing due to safety concerns. It has been encouraging to see a recent trend of moving towards non-halogenated solvents for processing of conjugated polymers, but most solvents still tend to be aromatic and require solvent-specific additives.[36b,121] Instead of just altering the solvent, several recent examples have modified conjugated polymers with polar or ionic side chains to achieve a desired solubility.[122] Methods of obtaining water/alcohol solubility have also been discussed in the side chain engineering section this chapter. While these methods allow for printing in safer solvent systems, there is often a tradeoff in performance. Future work will look to find ways of making that tradeoff more manageable, but safety should continue to be part of the equation when developing new conjugated polymers. Another challenge is in consistency and formulation. Printed inks usually contain a cocktail of additives to provide viscosity and controlled film formation, but conjugated polymer inks may not tolerate these types of additives, and may not hold up to the printing process without them. Polymer chemists can seek to alleviate consistency issues by designing systems tolerant to variations in processing conditions. Highly disordered polymers do not require formation of highly ordered films and, as mentioned in the molecular weight and dispersity effects section, can still achieve high charge mobility through interconnected aggregates. New methodology in random and block copolymers can open up possibilities for embedding disorder into otherwise ordered conjugated backbones, and can also be used to add controlled amounts of monomers containing solubility or viscosity modifiers. Melt-processable and stretchable conjugated polymers, as seen in Chapters 7 and 17 in Volume II, can open up entirely new avenues for large-scale processing, potentially even leading to impacts in 3-D printing of electronic components. All of these opportunities in printing of conjugated polymers can be supported through molecular engineering of polymer structure, which is one of the unique advantages of conjugated polymer conductors and semiconductors. Applications of printed conjugated polymers usually derive their competitive edge from their low cost, but in biological applications these materials may prove to be ideal for interfacing electronics with biological media. For one, conjugated polymers are soft materials that can be designed to swell, contract, stretch, and heal, making them an ideal fit for growing cells or mimicking tissues. Conjugated polymers can also provide a large surface area for biological interactions, as ions can penetrate their surfaces and fully interact with the bulk of the material. Finally, conjugated polymers can be engineered for ion transport, which can be used to heighten sensitivity, or harnessed to control and regulate ionic flow. While electrochemically prepared polyaniline and polypyrrole were used for biosensing applications in the 80s and 90s, greater control over polymer synthesis has led to a rebirth in the area of conjugated polymer bioelectronics. The main efforts in this area center around an organic electrochemical transistor (OECT), which is a type of biosensor that translates ionic flux in biological media to changes in transistor current, and the organic electronic ion pump, which can be used to regulate flow of ions.[123] Both of these devices rely on the active material’s ability to transport both electronic and ionic charges, and so work over the next decade will seek to find ways to modify high charge mobility conjugated polymers to also exhibit ionic transport. As mentioned in the side chain engineering section of this chapter, oligoether side chains have been used to promote ionic interactions in conjugated polymers for OECT applications, but there are still many possibilities to explore in terms of both side chain and backbone modification to enhance ionic transport. Outside of bioelectronics, conjugated polymers are promising materials for tissue engineering and imaging applications. Soft materials are ideal for replicating tissue, but not many soft materials are electroactive. Conjugated polymers can bridge this gap, and have been used for cell scaffolding applications where oxidation of the polymer leads to more effective cell growth and differentiation.[124] For imaging applications, conjugated polymers can be engineered for high fluorescence, biocompatibility, and potentially specific binding to analytes through side chain modification.[125]

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Conjugated Polymers: Perspective, Theory, and New Materials

All of these mentioned applications in printed electronics and biotechnology can be greatly improved through continued work in conjugated polymer chemistry. The ability to affect materials’ properties through structural design has always been one of the reasons conjugated polymers are fascinating and worth studying. The field has produced more tools over the past decade to allow chemists greater freedom to create materials for specific purposes, and it will be exciting to watch how chemists use these tools to address current and new problems going forward.

Acknowledgments JRR appreciates and acknowledges funding from the National Science Foundation (1506046), the Office of Naval Research (N00014-17-1-2243; N00014-18-1-2222; N00014-16-1-2165), the Air Force Office of Scientific Research (FA9550-18-1-0184), and NXN Licensing for funding of the Georgia Tech Electroactive Polymer Program during the writing of this manuscript. BCT acknowledges funding from the National Science Foundation (MSN under CHE-1609881, CBET-1436875) and RMP acknowledges the USC Dornsife Graduate School Fellowship.

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4 Advances in Discrete Length and Fused Conjugated Oligomers 4.1 Introduction....................................................................................... 149 4.2 Oligothiophenes................................................................................. 150 4.3 4.4

Shanshan Chen, So-Huei Kang, Sang Myeon Lee, Tanya Kumari, and Changduk Yang

End-Group Modification • Conjugation Length Extension

Cyclopentadithiophene Derivatives................................................ 153 Heteroatom Modification • Regiochemistry Studies • Conjugation Length Extension • End-Group Modification

Benzodithiophene Derivatives........................................................ 162 Conjugated Length Extension • Core Unit Modification • End-Group Modification

4.5

Indacenodithiophene Derivatives................................................... 166

4.6

Rylene Diimide Derivatives............................................................. 172

Core Unit or π-Bridge Modification • Conjugation Length Extension • End-Group Modification Conjugation Length Extension

4.7 Others.................................................................................................. 179 4.8 Conclusion.......................................................................................... 186 Acknowledgments......................................................................................... 187 References....................................................................................................... 187

4.1 Introduction An important topic in materials science in recent years has been the development of organic semiconductors and their extensive use in electronics and photonics applications such as organic solar cells (OPVs), organic field-effect transistors (OFETs), sensors, and photodetectors [1–4]. These materials are promising in terms of their electronic properties, low cost, versatility of functionalization, thinfilm flexibility, easy fabrication, etc. The emergence of π-conjugated oligomers with discrete structures sequentially constructed from smaller organic building blocks is particularly inspiring. The appeal of such π-conjugated oligomers is their well-defined molecular structure, ease of purification, and good batch-to-batch reproducibility, which facilitates physical characterization and enables precise structure–property–performance relationships to be established [4–7]. Considerable efforts have been dedicated to developing novel π-conjugated oligomers for the aforementioned applications, and tremendous progress has been achieved [8–10]. For example, the use of indacenodithiophene (IDT)-based oligomers as acceptors has resulted in nonfullerene OPVs with power conversion efficiencies (PCEs) greater than

149

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14%, making these materials ideal candidates for high-performance OPVs [11, 12] In addition, highly crystalline π-conjugated oligomers generally exhibit a strong tendency to self-assemble into long-­ range-ordered structures, thereby contributing to high charge carrier mobilities in thin films [13–15]. Moreover, the electronic and structural properties of oligomers can be used as templates to yield valuable insights into the corresponding polymeric materials, such as to probe their mean conjugation length or their optical band gaps and to refine synthesis strategies [16, 17]. Therefore, π-conjugated oligomers are a promising family of semiconductor materials and have led to many recent exciting breakthroughs in the field of organic electronics. In this chapter, a brief survey of recent developments in π-conjugated oligomers and their application in OPVs and OFETs is presented. We focus on the rational design of representative oligomers based on thiophene, cyclopentadithiophene (CPDT), benzodithiophene (BDT), IDT, and rylene imide units. Correlations among chemical structures, physical properties, and resulting device performance are also proposed to provide a comprehensive understanding of topics ranging from materials design to various device applications. Furthermore, the remaining challenges and key research directions in the near future are also addressed.

4.2 Oligothiophenes Oligothiophenes are among the most highly investigated organic semiconductors because of their excellent charge transport properties, high polarizability, and simple synthesis. In 1989, for the first time, α-sexithiophene (T1) was introduced as an active component in OFETs and a carrier mobility of 10−3 cm2·V−1·s−1 was obtained [18]. The high hole mobility of T1 contributed to an initial PCE value of approximately 2.0% when blended with a C60 acceptor in OPVs (Table 4.1) [19]. However, its narrow absorption and relatively high-lying highest occupied molecular orbital (HOMO) level generated limited photocurrent generation and a large voltage loss. To achieve high-efficiency OPVs, conjugated oligothiophenes with a small bandgap and appropriate energy levels are required. One successful approach is to introduce electron-withdrawing units into the conjugated backbone to form push–pull chromophores (Scheme 4.1), which helps to extend the absorption range and enables control over the separation of the HOMO and lowest unoccupied molecular orbital (LUMO) electron density for efficient charge transfer [13]. Furthermore, variations in the π-conjugation length and substituted end groups in the molecule strongly affected their packing arrangement and solid state miscibility with fullerene acceptors and thus resulted in substantial performance differences.

4.2.1 End-group Modification In 2011, using a Knoevenagel condensation, Chen et al. synthesized a new class of oligothiophenes endcapped with electron-withdrawing alkyl cyanoacetate groups (T2, T3, and T4) (Figure 4.1) [20]. The optical bandgap of these materials as thin films was estimated to be 1.73–1.75 eV, with a gradual blue-shift of the absorption maximum when the ethyl end-group was replaced with larger octyl to 2-ethylhexyl end groups. All three oligomers demonstrate a high PCE (4.46%–5.08%) for PC61BM-based OPVs, and a PCE of 5.08% was achieved with T3:PC61BM without any post-treatment. To further improve the light-harvesting ability, T9 with 3-ethylrhodanine dye end groups was also designed [21]. The T9 film cast from chloroform shows a broad absorption from 450 to 750 nm with a reduced optical bandgap of 1.69 eV. Its stronger solar absorption contributed to a much-improved short-circuit current density (JSC) compared with that of T3, which led to a substantially higher PCE (6.10%). Later, the alkyl cyanoacetate terminal in the oligothiophene systems was replaced with various electron-withdrawing end groups such as double rhodanine (T11), 1,3-dimethylpyrimidine-2,4,6(1H,3H,5H)-trione (T5), 1,3-indanedione derivatives (T6, T7, and T8), to further investigate the correlation between these different end groups and device performance [22, 23]. In particular, T10, which contained 2-(1,1-dicyanomethylene)rhodamine as the end

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Advances in Discrete Length and Fused Conjugated Oligomers TABLE 4.1  Optical, Electronic Properties, Mobilities, and Photovoltaic Properties of Oligothiophenes

T1 T2 T3 T4 T5 T6 T7 T8 T9 T10 T11 T12 T13 T14 T15 T16

λmax (nm)a

Eg (eV)a

HOMO/ LUMO (eV)

µe (cm2V−1s−1)

µh (cm2V−1s−1)

JSC (mA cm−2)

VOC (V)

FF (%)

PCE (%)

Refs.

− 591 580 569 605 692 684 802 618 764 584 573 577 602 685 603

− 1.73 1.74 1.75 1.67 1.49 1.2 1.33 1.69 1.62 1.7 1.77 1.6 1.61 1.6 1.59

−5.3/−3.1 −5.09/−3.33 −5.13/−3.29 −5.10/−3.26 −5.12/−3.50 −4.97/−3.44 −4.90/−3.86 −5.02/−3.72 −5.00/−3.28 −4.95/−3.36 −5.09/−3.39 −5.34/−3.46 −5.16/−3.56 −5.02/−3.45 −5.22/−3.41 −4.97/−3.44

− − − − − − − − − 1.28 × 10−4 c − − − − −

1.0 × 10 4.51 × 10−4 b 3.26 × 10−4 b 1.94 × 10−4 b 0.47 × 10−4 b 1.73 × 10−4 b 0.3 × 10−4 b − 1.50 × 10−4 b 5.91 × 10−4 c 0.24 × 10−4 b − 2.61 × 10−4 c 5.77 × 10−4 c 6.54 × 10−4 c 5.11 × 10−4 c

5.6 9.94 10.74 9.91 7.54 8.56 3.14 − 13.98 14.87 6.77 0.7 10.88 10.98 15.88 13.91

0.70 0.88 0.86 0.93 0.9 0.8 0.76 − 0.92 0.91 0.92 0.9 0.92 0.87 0.92 0.82

51 51 55 49.1 60 72 28 − 47.4 68.7 39 38 59 68 69 69

2.0 4.46 5.08 4.52 4.05 4.93 0.66 − 6.1 9.3 2.46 0.24 6.33 6.5 10.08 7.86

[18, 19] [20] [20] [20] [23] [22] [22] [22] [21] [24] [23] [25] [25] [25] [25] [25]

−3 d

  Thin film.   Neat film mobilities measured via SCLC method. c  Blend film mobilities measured via SCLC method. d  Neat film mobilities measured via OFET method. a

b

SCHEME 4.1  The general synthesis procedure of end-capped oligothiophenes.

group, exhibited a narrow optical bandgap of 1.62 eV and enhanced the ground-state dipole moment that facilitates electronic coupling between neighboring molecules [24]. When blended with PC71BM as an acceptor, the optimal nanoscale interpenetrating network with highly crystalline fibrils (~10 nm in diameter) provided PCEs as high as 9.30% with a nearly 100% internal quantum efficiency in devices.

4.2.2 Conjugation Length Extension The conjugation length of T10, which is based on seven thiophene units, enables excellent OPV performance; however, the effect of the conjugation length of oligomers on their properties requires further investigation. In 2015, Chen et al. reported a series of 2-(1,​1-dic​yanom​ethyl​ene)r​hodam​ine-t​ermin​ated oligomers T12–T16, with four to nine thiophene units in their backbones, to comprehensively investigate the effect of the conjugation length (Figure 4.2) [25]. The optical band gaps of T12–T16, as estimated from the onset of film absorption, were 1.77, 1.60, 1.60, 1.61, and 1.59 eV, respectively. Because the LUMO levels are largely determinate by the electron-withdrawing end groups, their LUMO levels are similar, whereas the HOMO levels increase from −5.34 to −4.97 eV with increasing conjugation length. Notably,

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FIGURE 4.1  End-group modification for oligothiophenes.

FIGURE 4.2  Conjugation length extension for oligothiophenes.

the OPVs based on T15, T10, and T16, which have an odd number of thiophene units, generally exhibited much higher JSC values than those based on T13 and T14, which have even numbers of thiophene units; this trend is attributed to the larger dipole moment differences and better-developed fibrillar network in blends containing T15, T10, and T16. Among them, T15 exhibited an outstanding PCE of 10.10%, which is among the highest PCEs reported for a single-junction fullerene-based OPV to date.

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SCHEME 4.2  Synthesis procedure of CPDT.

4.3 Cyclopentadithiophene Derivatives Fused aromatic or heteroaromatic molecules have rigid and extended conjugation structures, enabling a decrease of the energy bandgap and promoting closer intermolecular interactions. Furthermore, the rigid fused-ring structures also can lower the reorganization energy, increase the rate of intermolecular hopping, and ultimately affect the charge transport in organic semiconductors [26]. Since Wynberg et al. [27] introduced the synthesis of 4H-cyclopenta[2,1-b:3.4-b′]dithiophene by ring-closure reaction between two thiophenes in 1968 (Scheme 4.2), the intrinsic physical and electrochemical properties and the charge transport characteristics of CPDT-based oligomers have been widely studied [28–32].

4.3.1 Heteroatom Modification In 2008, Demadrille et al. [33] incorporated CPDT derivatives as a donor unit terminated to a fluorenone building block. As shown in Figure 4.3, CPDT1 has alkyl chains protruding from its πconjugated backbone plane, whereas CPDT2 has alkyl chains within the plane, as imposed by its sp2 hybridization. The sp2 hybridization of CPDT2 narrowed the optical bandgap by slightly raising the energy levels but unfortunately led to diminished OPV performance when CPDT2 was combined with PC 61 BM. Bazan’s group studied the coordination effect of Lewis acids to electron-acceptor fragments of a CPDT-based oligomer by binding B(C6F5)3 to a nitrogen atom of benzo[2,1,3]thiadiazole (BT) (CPDT3 to CPDT5) or to a nitrogen atom of pyridyl[2,1,3]thiadiazole (PT) (CPDT4 to CPDT6) [34]. With increased electron density and less-hindered coordination sites, boron–nitrogen (B–N) interactions in the Lewis acid adduct became stronger, which led to narrowing of the optical bandgaps. Zhu et al. [35] designed CPDT-bridged donor–acceptor-π-bridge–acceptor sensitizer oligomers with different donor units (CPDT7–CPDT10) in 2015. In the oligomers, CPDT was used as an enlarged π-conjugation bridge. For example, using indoline, which exhibits strong donating power, for the donor unit (CPDT10) increased the JSC without sacrificing the open-circuit voltage (VOC). Moreover, CPDT10 showed rapid dye adsorption without co-sensitization. In 2017, two planar and symmetrical acceptor–donor–acceptor-type CPDT-based oligomers (CPDT11 and CPDT12) were synthesized [36]. The molecular structures of the donor oligomers were designed with similar characteristics to form films with energetically identical properties but different morphologies. After a solvent-vapor annealing treatment, the molecular packing of the films was reorganized. CPDT12 formed smaller molecular self-domains than CPDT11, resulting in better mixing with PC71BM and less recombination. This behavior was attributed to greater molecular aggregation as a result of the absence of two hexyl chains. The carbon bridge at the 4-position of CPDT can be substituted with another group-14 element: Si. Several groups have claimed that the charge carrier transport and the photoluminescence efficiency of Si-substituted materials could be improved dramatically. For instance, Holmes et al. [37] and Cao et al. [38–40] synthesized Si-substituted fluorene-based polymers (PSiF, P36–27SiF90 and PSiF-DBT in Figure 4.4) for blue light-emitting diodes and solar cells with enhanced material properties even though the absorption spectrum was similar to those of unsubstituted polymers. Si atoms have similarly been incorporated into CPDT-based materials. The first silicon-bridged bithiophene, dithienosilole (DTS), was introduced by Ohshita in 1999 [41]. PCPDTBT, which is known as one of the most efficient

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FIGURE 4.3  CPDT-based oligomers.

narrow-bandgap photovoltaic materials, was modified by introducing a Si atom into the CPDT unit (PSBTBT), as shown in Figure 4.5 [42, 43]. The HOMO and LUMO levels of PSBTBT were raised and the bandgap was larger compared with those of PCPDTBT. The Jsc and Voc were both increased, resulting in a substantial PCE enhancement from 3.2% to 5.1%.

4.3.2 Regiochemistry Studies In 2011, three donor–acceptor-donor–acceptor-donor type oligomers (DTS1–DTS3 in Figure 4.5) were synthesized on the basis of DTS and PT as the donor and acceptor moiety, respectively, to obtain a narrow bandgap [44]. When the end groups were changed from thiophene (DTS1) and benzothiophene (DTS2) to bithiophene (DTS3), the band gaps decreased from 1.68 eV and 1.64 eV to 1.51 eV, resulting in a highest PCE of 3.2% in OPVs. Later, Bazan et al. [45] modified the distal/distal (N atoms pointing away from the core) configuration of the PT unit in DTS3 to a proximal/proximal configuration of PT (N atoms pointing toward the core), yielding oligomer DTS4.2 The energy levels and the optical properties of DTS4 were quite similar to those of DTS3; however, the PCE was greatly increased to 4.52%. With further addition of diiodooctane (DIO) as an additive to optimize the active layer morphology, a PCE of 6.7% was achieved. To better understand the correlations among the molecular structure, dipole moment, self-assembly, and OPV performance, oligomers with proximal/distal-oriented PT (DTS5) and BT (DTS6) were synthesized [46]. The order of increasing charge transport mobilities and OPV performances was ascribed to the symmetrically disposed dipole moment and stronger nitrogen–sulfur locked conformation. Although they both involve the nonbonding electrons of the N atom on a PT unit, the synthesis methods of oligomers with proximal/proximal and distal/distal configurations differ. The

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155

FIGURE 4.4  Chemical structures of the polymers mentioned in this chapter

literature also includes several reports that the pyridyl nitrogen of PT can be protonated or react with acids to form an adduct. Thus, the acidic nature of commonly used interlayers, i.e. poly(​3,4-e​ thyle​nedio​x ythi​ophen​e)–po​ly(st ​y rene​sulfo​nate)​ (PEDOT:PSS), would give negative effect on the PT-containing donor materials [34, 47, 48]. To solve this problem, researchers replaced the PT unit in DTS4 with fluorine-substituted benzothiadiazole (FBT) as another acceptor fragment, leading to DTS7, as shown in Figure 4.6 [49]. Unlike the synthesis of DTS-PT oligomers, DTS-FBT oligomers are prepared through an opposite type of synthesis route because of the electron-withdrawing ability of fluorine, which makes a more reactive site far from the fluorine-substituted site, as shown in Scheme

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FIGURE 4.5  DTS-based oligomers with different regiochemistry.

FIGURE 4.6  Introduction of fluorinated benzothiadiazole for acidic fabrication conditions.

4.3. Compared with DTS4, the HOMO and LUMO levels of DTS7 were slightly raised and the optical bandgap was increased. To demonstrate that an FBT-containing oligomer can be more resilient against acid conditions, comparative studies of DTS4 and DTS7 were carried out. When trifluoroacetic acid (TFA) was added to the solutions, the UV–Vis spectrum of DTS4 constantly changed as acid was added, whereas the UV–Vis spectrum of DTS7 did not change even in the presence of ten equivalents of acid. As a result, the PCE of the OPVs fabricated using DTS7 and PEDOT:PSS was 7.0%. When an inverted solar cell structure with a ZnO optical spacer and a barium cathode layer were used, the PCE further was increased to 9.0% [50].

4.3.3 Conjugation Length Extension The effects of extending the discrete conjugation length of DTS-containing oligomers on their molecular properties and device performance have been studied. DTS8 and DTS9 with different numbers of thiophene spacers have been synthesized (Figure 4.7) [51]. An increase in discrete conjugation length led to a narrow optical bandgap, an increase of the HOMO level, and an increase of the thermal resistance. The effects of extending conjugation were also studied with the DTS4, DTS10, and DTS11 oligomer set [52]. Analogously, a higher degree of conjugation increased the electron delocalization and narrowed the optical bandgaps from 1.52 to 1.41 eV with enhanced absorption coefficients. DTS10 exhibited good thermal stability at temperatures greater than 200°C when applied to OFET devices; in addition, a PCE of 6.5% was attained without solvent additives or thermal annealing.

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157

SCHEME 4.3  Different synthesis route for proximal/proximal and distal/distal configuration of PT and FBT.

Based on the molecular structure of DTS10, DTS12–DTS14 with various core units were synthesized to further narrow the optical bandgap, resulting in a greater degree of conjugation and enhanced OPV performance [53]. Extending the conjugation length of the best-performing DTS10 and DTS13 oligomers led to DTS11 and DTS15, which afforded a further narrower optical bandgap and improved performance. Again, upon incorporation of an FBT unit into DTS10 and DTS11, monofluorinated DTS16, difluorinated DTS17, and double-monofluorinated DTS18 were obtained. No obvious changes in the absorption coefficients were observed; however, the optical bandgap for the fluorinated oligomers increased, as did their molecular rigidity and planarity. These changes were attributed to the fluorine–sulfur interactions between adjacent aromatic units. Next, the PT units in DTS10, DTS17, DTS11, and difluorinated DTS18 were completely substituted with FBTs, yielding DTS19–DTS22 oligomers with comparable optical bandgaps and energy levels [54]. Among them, a higher degree of fluorination led to a larger optical bandgap and enhanced thermal resistance. In a molecule with DTS as the core unit and thieno[2,3-c]pyrrole-4,6-dione as an end-group, the number of thiophene spacers was controlled [55]. Adding one thiophene spacer from DTS23 to DTS24 slightly narrowed the absorption bandgap and downshifted the LUMO level. The photovoltaic performance was better for one thiophene spacer because of the higher VOC value. DTS25 was also designed and synthesized by combining two weakly conjugated chromophores to extend the conjugation length [56]. However, the highly tilted nonplanar molecular structure of the newly synthesized molecule disrupted the electron delocalization, which led to a larger bandgap and a kinetic barrier to crystallization. Crystalline of DTS25 grown under slow solvent evaporation conditions was investigated by grazingincidence wide-angle X-ray scattering (GIWAXS) and absorption spectroscopy, yet the PCE was still limited to 1.3% even after treatment with DIO additive. This work highlights the need to develop both a new processing strategy that enables different film growth profiles and a molecular design strategy to achieve a high degree of conjugation with a planar molecular topology.

4.3.4 End-group Modification To assess the influence of the end groups of DTS-containing oligomers on their molecular properties and photovoltaic performance, DTS26–DTS29 were synthesized (Figure 4.8) [51]. With the adoption

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FIGURE 4.7  Conjugation length extension for DTS-based oligomers.

of end groups with a higher electron affinity across the molecular backbone, from DTS26 to DTS29, the optical bandgap became narrow and the HOMO and LUMO levels became deeper. The thermal properties also tended to improve from DTS26 to DTS28, whereas DTS29 exhibited thermal instability at 300°C. Kim et al. [57] investigated the effects of intermolecular interactions of electron-donating oligomers by introducing different terminal groups of ester and amide groups with different alkyl chain lengths. Compared to the ester terminal groups (DTS30–DTS32), amide terminal groups (DTS33– DTS35) exhibited stronger intermolecular interaction by hydrogen bonding. The molecular packing and orientation changed substantially with a change in the type of terminal groups. The morphology of the blend film with PC61BM was also influenced by the terminal group. The overall performance was better for the ester group with the better morphology and stronger π–π stacking.

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159

Ding et al. [58] designed and synthesized a dumbbell-like acceptor–donor–acceptor molecule, DTS36, containing DTS and two fullerene acceptor units. Sonogashira coupling of but-3-yn-1-ol and 4,7-dibromo-2,1,3-benzothiadiazole afforded BT derivative, which was reacted with DTS by Stille coupling to make an intermediate of DTS36. Esterification of the intermediate and PCBA yielded DTS36. DTS36 showed an integrated version of the intermediate and PC61BM for the absorption spectra and cyclic voltammograms, and showed a PCE of 0.4% in a single-component solar cell with a VOC of 0.79 V, a JSC of 1.75 mA·cm−2, and a fill factor (FF) of 0.27. Chen et al. [59] synthesized DTS37 with 3-ethylrhodanine end-group to make an acceptor–donor–acceptor-type structure (Figure 4.8). The molecular packing of a DTS37 and PC71BM blend film was strong, and resonant soft X-ray scattering (RSoXS) showed a suitable phase separation with a domain size of 20−30 nm after thermal annealing followed by solventvapor annealing, resulting in a high PCE of 8% for the DTS37-based device. DTS trimer-based oligomers with FBT, difluorine-substituted benzothiadiazole, and rhodanine end groups were also designed and synthesized by Cao et al. [60] resulting in DTS38–40, where the oligomers were prepared from a combination of Stille coupling and Knoevenagel condensation. The optical bandgap became narrower in the order DTS38, DTS39, and DTS40, with deeper HOMO and LUMO levels; the OPV performance was substantially high for DTS40, which exhibited a high JSC and a high FF compared to those of DTS38 and DTS39 after the solvent-vapor annealing. This result was attributed to broader absorption as well as to greater crystallinity and enhanced charge carrier transport. These results imply that proper selection of end groups also can result in a dramatic improvement in photovoltaic performance. With respect to the DTS analogs, Ge was also introduced as a bridge of bithiophene. First, Reynolds et al. [61] reported a dithienogermole (DTG)-containing polymer in 2011 (Figure 4.9). The greater bond length of C–Ge compared with that of C–Si modified the molecular conformation. Consequently, polyDTG-TPD showed a smaller bandgap, broader light absorption, and greater crystallinity than polyDTSTPD, resulting in in a PCE of 7.3%, which is 0.7% greater than that of polyDTS-TPD. In 2015, Sun et al. reported a DTG-based oligomer as a high-performing donor molecule in OPVs [62]. In contrast to the aforementioned results based on DTS and DTG polymers, the long C–Ge bond length in DTG1 did not

FIGURE 4.8  End-group modification for DTS-based oligomers.

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FIGURE 4.9  End-group modification for DTG-based oligomers. TABLE 4.2  Optical, Electronic Properties, Mobilities, and Photovoltaic Properties of CPDT-based Oligomers

CPDT1 CPDT2 CPDT3 CPDT4 CPDT5 CPDT6 CPDT7 CPDT8 CPDT9 CPDT10 CPDT11 CPDT12

λmax (nm)a

Eg (eV)a

HOMO/LUMO (eV)

413 402 646 764 626 868 533 536 546 551 651 647

1.94 1.85 1.57 1.31 1.61 1.08 − − − − 1.5 1.5

−5.24/−3.25 −5.14/−3.23 −4.86/−3.27 −5.37–4.06 −4.79/−3.18 −5.08/−4.00 −6.7/−3.8 −6.76/−3.89 −6.69/−3.76 −6.63/−3.79 −5.31/−3.60 −5.32/−3.61

µh JSC (cm2V−1s−1) (mA cm−2) VOC (V) − − − − − − − − − − 1.6 × 10−4 b 6.2 × 10−4 b

2.06 1.85 − − − − 18.24 16.23 12.32 19.69 6.25 8.25

0.75 0.67 − − − − 0.65 0.69 0.70 0.70 0.81 0.85

FF (%)

PCE (%)

Refs.

31.0 30.0 − − − − 71.8 71.6 72.7 73.1 54.0 53.0

0.48 0.37 − − − − 8.49 8.04 6.27 10.08 2.77 3.72

[33] [33] [34] [34] [34] [34] [35] [35] [35] [35] [36] [36]

  Thin film.   Blend film mobilities measured via SCLC method.

a

b

increase molecular ordering relative to that in DTS7. DTG1 (PCE ≈ 7%) showed only comparable performance to DTS7 with PC71BM in conventional and inverted-structure OPVs. Yang et al. [63] investigated the end-group effect for two DTS oligomers (DTS7 and DTS41) and two DTG oligomers (DTG1 and DTS2). Despite the similar optical or electrochemical properties between DTS and DTG oligomers, the tendency to form a crystalline domain was stronger for DTG oligomers, leading to slightly larger crystal features and rougher films. DTG-based oligomers were found to have a higher JSC with slightly lower VOC and FF compared to DTS-based oligomers. When ZnO was used, the best PCE of a DTG1 reached 7.3%. In 2016, Singh et al. [64] also synthesized DTG3 with superior crystallites as a DTG-containing DTS25 derivative and reported a high PCE of 9.1% under optimized conditions. In addition, Ohshita et al. [65] reported DTG- and PT-based photosensitizers (DTG4–6) covering the visible to the near-infrared region for dye-sensitized solar cells (Tables 4.2 through 4.4).

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Advances in Discrete Length and Fused Conjugated Oligomers TABLE 4.3  Optical, Electronic Properties, Mobilities, and Photovoltaic Properties of DTS-based Oligomers λmax (nm)a

Eg (eV)a

DTS1 DTS2 DTS3 DTS4

652 664 720 720

1.68 1.64 1.51 1.5

−5.19/−3.46 −5.25/−3.55 −5.16/−3.60 −5.2/−3.60

− − 5.3 × 10−5 b 6.0 × 10−3 b

DTS5 DTS6 DTS7

715 663 580

1.52 1.58 1.55

−5.30/−3.78 −5.15/−3.57 −5.12/−3.34

DTS8 DTS9 DTS10 DTS11 DTS12 DTS13 DTS14 DTS15 DTS16 DTS17 DTS18 DTS19 DTS20 DTS21 DTS22 DTS23 DTS24 DTS25 DTS26 DTS27 DTS28 DTS29 DTS30 DTS31 DTS32 DTS33 DTS34 DTS35 DTS36

620 680 − 758 714 751 749 781 742 726 746 − − − − 521 520 575 604 664 670 625 586 598 566 558 572 496 432, 544 610 592

1.63 1.51 1.44 1.41 1.55 1.40 1.43 1.36 1.47 1.50 1.45 1.49 1.54 1.46 1.55 1.92 1.87 1.9 1.64 1.63 1.62 1.59 1.76 1.75 1.82 1.85 1.81 1.92 − 1.66 1.75

DTS37 DTS38

HOMO/LUMO µe JSC (eV) (cm2V−1s−1) µh (cm2V−1s−1) (mA cm−2)

VOC (V)

FF (%)

PCE (%)

Refs.

− − 10.90 14.40

− − 0.70 0.78

− − 42 59.3

− − 3.20 6.70

[44] [44] [44] [45, 46]

−5.20/−3.48 −5.07/−3.55 −5.17/−3.73 −5.04/−3.71 −5.29/−3.30 −5.17/−3.74 −5.15/−3.71 −5.11/−3.79 −5.20/−3.73 −5.26–3.75 −5.08/−3.69 −5.2/−3.7 −5.3/−3.7 −5.1/−3.7 −5.3/−3.7 −5.55/−3.44 −5.52/−3.57 −5.29/−3.14 −5.20/−3.47 −5.26/−3.55 −5.34/−3.60 −5.42/−3.75 −5.28/−3.52 −5.27/−3.52 −5.47/−3.65 −5.35/−3.50 −5.34/−3.53 −5.05/−3.10 −5.85/−4.31

− − − − − − − − − − − 0.003 c − − − − − − − − − − − − − − − − − − − − 5.75 × 10−4 b

− − 1.5 × 10−6 b 2.0 × 10−3 b 0.20 c 0.05 c 0.01 c − − − − 0.14 c 0.04 c 0.03 c 0.05 c 0.18 c 0.001 c 0.06 c 0.10 c 0.06 c 0.10 c 0.06 c 0.07 c 0.15 c − − 2.55 × 10−4 b − − − − 1.375 × 10− c 2.825 × 10−2 c 2.705 × 10−4 c 2.185 × 10−3 c 2.995 × 10−3 c 6.545 × 10−6 c 3.65 × 10−5 b

9.80 0.90 12.80 15.47 − − 13.60 15.20 8.60 12.60 12.70 12.60 12.90 13.50 14.20 − − − − 2.60 2.59 6.39 − − − − 9.79 9.30 7.75 7.94 8.38 1.25 1.75

0.72 0.83 0.81 0.78 − − 0.71 0.66 0.83 0.75 0.72 0.75 0.75 0.76 0.71 − − − − 0.97 0.88 0.75 − − − − 0.82 0.82 0.94 0.87 0.86 0.64 0.79

45 25.8 68 75.1 − − 60 65 56 61 60 62 65 59 65 − − − − 47.58 32.90 27 − − − − 54 57 41 47 52 26 27

3.16 0.19 7.00 9.02 − − 5.80 6.50 4.00 5.70 5.50 5.80 0.30 6.10 6.50 − − − − 1.20 0.75 1.30 − − − − 4.31 4.31 3.00 3.22 3.75 0.21 0.40

[46] [46] [49] [50] [51] [51] [52, 53] [52, 53] [53] [53] [53] [53] [53] [53] [53] [54] [54] [54] [54] [55] [55] [56] [51] [51] [51] [51] [57] [57] [57] [57] [57] [57] [58]

−4.94/−3.28 −4.84/−3.09

4.11 × 10−5 b 2.86 × 10−4

4.955 × 10−4 b 8.715 × 10−5 b

13.97 4.58

0.83 0.84

70 34.19

8.02 1.29

[59] [60]

3.465 × 10−5 b 1.085 × 10−4 b

5.71 12.69

0.83 0.89

43.80 68.54

1.96 7.56

[60] [60]

b

DTS39 DTS40

563 622, 671

1.76 1.67

−4.86/−3.10 −4.92/−3.25

1.98 × 10−4 b 2.95 × 10−4 b

  Thin film.   Blend film mobilities measured via SCLC method. c  Neat film mobilities measured via OFET method. a

b

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TABLE 4.4  Optical, Electronic Properties, Mobilities, and Photovoltaic Properties of DTG-based Oligomers

DTG1 DTG2 DTG3 DTG4 DTG5 DTG6

λmax (nm)a

Eg (eV)a

HOMO/LUMO (eV)

JSC (mA cm−2)

VOC (V)

FF (%)

PCE (%)

Refs.

694 583,635 609,630 589,411 583,410 609,411

1.64 1.77 1.62 1.77 1.78 1.71

−5.15/−3.65 −5.25/−3.68 −5.17/−3.40 −5.03/−3.26 −5.04/−3.26 −5.00/−3.29

14.60 9.10 15.90 9.52 7.12 6.11

0.76 0.80 0.79 0.43 0.48 0.37

65.3 36.3 73.2 67 66 60

7.30 2.66 9.10 2.76 b 2.29 b 1.36 b

[63] [63] [64] [65] [65] [65]

  Thin film.   Blend films for dye-sensitized solar cells.

a

b

4.4 Benzodithiophene Derivatives Beimling et al. [66, 67] first reported BDT, which exhibits a symmetrical and rigid planar structure, in the 1980s. Subsequent developments in the synthesis of this core block have provided easy accessibility to p-type materials. The initial methods used to synthesize unsubstituted and substituted BDT are shown in Scheme 4.5. In the first step of the synthesis of the unsubstituted BDT unit, its structurally coplanarity, which was expected to result in dense π–π stacking, made it applicable to OFETs, leading to a high hole mobility of 0.04 cm2·V−1·s−1 [68]. To further utilize BDT core unit for various applications, substitutions into the central backbone were required for solution-processability, with the objective of exploiting the advantages of two-dimensional (2D)-conjugated structures. The first devised synthesis routes to the methyl-substituted BDT from 2,5-dimethylaniline were complicated, involving several steps; synthesis routes to methoxy-substituted BDT from benzo[1,2-b:4,5-b′]dithiophene-4,8-dione had low yields. However, in 2008, Hou et al. [69] modified the reported method and achieved high yields of the didodecyloxy-BDT monomer; a series of BDT units have since been developed with various substituents at the 4- and 8-positions, such as alkyl and conjugated groups (Scheme 4.4).

4.4.1 Conjugated Length Extension A symmetric design of small molecules comprises a BDT core and flanked second donor/acceptor blocks in sequence. The main purpose of flanking units with electron-donating and -withdrawing properties is to extend the conjugation lengths with energy-level tunings. In the molecules BDT1 and BDT2,[70] the number of 3-hexylthiophene groups as a π-spacer was varied from one to two with different conjugation lengths (Figure 4.10). With the extended lengths, BDT2 exhibited greater hole mobility and broader absorption, resulting in a higher PCE (5.11%) compared with that of BDT1 (4.15%). On the basis of the same concept of extending conjugated backbones with thiophene and bithiophene, BDT3 and BDT4 were also synthesized as donor materials in OPVs [71]. However, in this case of increasing the thiophene bridges, a lower VOC of 0.83 V in the BDT4-based device led to PCE (1.62%) lower than that of the BDT3based device because of largely separated nanostructures in the film. To extend the conjugation lengths of a BDT block, a thiophene-annulated benzotrithiophene unit was developed for BDT7 and BDT8 with various numbers of thiophene spacers [72]. In an analogous strategy for conjugated lengths, the incorporation of naphthodithophene and its isomeric analog in BDT5 and BDT6 molecules was successful for broad absorption and well-ordered molecular packing [73, 74]. Interestingly, BDT6 had high-lying HOMO/LUMO levels with an identical bandgap. Therefore, a substantial reduction in the VOC values of BDT6 from BDT5 was observed; however, enhanced hole mobility led to beneficial charge separation and to less recombination, with a higher FF and higher JSC. A dithienobenzodithiophene building block was used to broaden conjugated lengths horizontally,

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SCHEME 4.4  Synthetic routes to the substituted and unsubstituted BDT units.

affording BDT9 and BDT10 with different molecular structures [75, 76]. Because of their large areas for π–π stacking and weak donating properties, both BDT9 and BDT10 showed excellent thermal stability, strong light-harvesting ability, and superior charge transport characteristics, leading to PCEs of 4.98% and 5.42%, respectively. BDT blocks have been widely used as donor materials blended with acceptors for OPVs, affording excellent device performance. However, the synthesis of a benzodi(cyclopentadithiophene) core unit in molecule BDT11 made it available as an accepting material with donor polymers as a promising candidate of fullerene acceptors [77]. Covalent rigidification at the 3- and 7-positions with fused thiophene units by sp3-carbon bridges on a BDT unit resulted in a planar structure with strong aggregation and lowered energy levels, affording broad absorption regions, high mobility, and optimized film morphology. In combination with the wide-bandgap polymer donor PBDB-T, an extremely high PCE of 10.42% was achieved with a JSC of 17.85 mA·cm−2 because of efficient hole transfer from the acceptor to the donor material. Substitution of octyl groups onto a benzene center of the benzodi(cyclopentadithiophene) core unit with fluoride atoms in terminal accepting units was performed in BDT12 [78]. Compared with BDT11, BDT12, exhibited a narrower optical bandgap of 1.45 eV in the red-shifted absorption range, which led to an increase in the device photocurrent. The BDT12-based film delivered an outstanding PCE of 12.12% with a high JSC of 20 mA·cm−2 by fine tuning of energy levels through slight chemical modification. In contrast to fullerene-based acceptors, the nonfullerene systems based on BDT11 and BDT12 showed an ultimately efficient charge generation, followed by effective charge extraction and low recombination loss, demonstrating the versatility of BDT derivatives in photovoltaic applications.

4.4.2 Core Unit Modification Encouraged by the strong potential based on the photovoltaic performances of BDT-based small conjugated molecules, researchers have developed BDT cores with substituents at the 4- and 8-positions, forming one-dimensional (1-D) or 2-D structures with nonconjugated and conjugated solubilizing groups (Figure 4.11). First, to investigate the effect of alkyl chains in 1-D conjugated small molecules for OPVs, three calamitic molecules were synthesized with the same conjugated system but with different

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FIGURE 4.10  Conjugation length extension for BDT-based oligomers.

substitutions of n-octyl, 2-ethylhexyloxy, and 2-ethylhexylthio groups for BDT13,[79] BDT14,[80] and BDT15,[81] respectively. Because of the weak electron-donating property of the octyl group on the BDT core, a lower-lying HOMO energy level in BDT13 was achieved than in BDT14, leading to a higher VOC value of 0.98 V (vs. 0.93 V) before post treatments. After thermal and solvent-vapor annealing, an optimized morphology of blending samples with BDT14 and PC71BM led to a high PCE of 8.26% with an FF value greater than 0.70. Furthermore, an alkylthio side chain was adopted in a BDT15 molecule, resulting in extensive changes in the molecular packing and in the optoelectronic properties. After post treatments, ordered and fibril nanostructures with beneficial interpenetrating networks in the blending film substantially enhanced its photovoltaic performance, resulting in a PCE of 9.95% with a JSC of 14.61 mA·cm−2, a VOC of 0.92 V, and a FF of 0.74.

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FIGURE 4.11  Core unit modification for BDT-based oligomers.

To compare nonconjugated chains with a thienyl side group, BDT13 to BDT15 analogs and BDT16based derivatives were synthesized as 2-D BDT conjugated molecules. Thiophene-substituted BDT16 exhibited a redshifted absorption spectrum relative to that of the BDT13–15 molecules bearing nonconjugated side chains, leading to a PCE of 8.12% [82]. By exploiting the advantage of a thiol side-chain in a BDT15 molecule, researchers introduced an alkylthio-thienyl group in BDT17, which resulted in a downshifted HOMO level, redshifted absorption spectrum, and enhanced hole mobility; BDT17 exhibited a PCE of 9.20% without any additives or post treatments [83]. Molecule BDT18, incorporating a bulky and long alkyl chain, led to unfavorable film morphology and low mobility because of large steric hindrance, limiting the PCE to 6.79% [82]. Instead of substituting a thienyl side group onto the 4and 8-positions of a BDT core, thienothiophene-substitution in BDT19 resulted in a film with ordered bicontinuous networks with preferable morphologies, which achieved an impressive FF of 0.72 with a PCE of 8.70% [84]. Replacement of thiophene with bithiophene in BDT20 facilitated the delocalization of π-electrons, leading to a high FF, yet a sacrificed VOC caused by strong electron-donating ability [82]. To maintain rigid coplanarity in a conjugated backbone, the insertion of triisopropylsilylethynyl groups into BDT21 broadened its absorption region with strong intramolecular charge transfer bands and lowered the HOMO energy level with increasing VOC [85]. However, the conventional device based on BDT21 with additives exhibited poor photovoltaic performance because of high phase segregation in the film. The effect of 1-D and 2-D conjugated systems was investigated in BDT22 and BDT23, whose

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Conjugated Polymers: Perspective, Theory, and New Materials

structures differ from those of BDT13-to-21-based molecules, using alkoxy and alkyl thienyl groups [77]. Consistent with the trends in the aforementioned research, BDT23 with thiophene conjugated side chains exhibited intense absorption, good charge transport, and finely ordered packing in the blend films, resulting in a remarkable PCE of 9.73%. A selenophene group was used in BDT24[86] to narrow the optical bandgap to 1.86 eV with an improved JSC of 10.50 mA·cm−2 compared with that of thiophenesubstituted BDT25 [87]. In attempts to modify the BDT core units in small discrete molecules, sulfur atom in the backbone of the BDT core were replaced with furan groups and a photovoltaic device based on BDT26 exhibited properties similar to those of BDT27, with a PCE of 3.70% (vs. 3.60%) [88]. Harschneck et al. [89] reported BDT28 and BDT29 using a BDT isomer, benzo[1,2-b:6,5-b′]dithiophene unit for self-organization properties. Solubilizing side-chain of the 3,7-dimethyloctyl group in BDT29 provided the optimum balance between solid-state aggregation and a miscible phase with PC71BM, resulting in a PCE (5.53%) greater than that of the BDT28-based device (1.39%) with detrimental segregated phase in its film state.

4.4.3 End-Group Modification On the basis of alternating donor–acceptor moieties, calamitic molecules with discrete donor and acceptor sections based on a BDT core unit have been developed for efficient OPVs to harvest light-absorption efficiently with beneficial charge transport (Figure 4.12). Among the various accepting units, isoindigo with a highly planar backbone was used as the accepting unit in BDT30, which was used with PC71BM in fabricated devices that exhibited a JSC of 4.89 mA·cm−2, a VOC of 0.72 V, and a FF of 0.43, leading to 1.51% PCE [90]. In the molecule BDT31, a thienopyrroledione-based moiety, which possesses strong electronwithdrawing character, with an alkyl-substituted imide fused onto thiophene was incorporated as the flanking groups beside the BDT core [91]. Because of the strong accepting effect, the BDT31-based device exhibited a high VOC of 0.92 V, with low-lying HOMO levels and a PCE of 2.40%. Moreover, Li et al. [92] synthesized a set of 2,1,3-benzothiadiazole (BT)-based BDT molecules, BDT32 and BDT33, with different sequences of thiophene and BT units. Compared with BDT33, the film based on BDT32 with a finely ordered crystalline structure in a blend with PC61BM exhibited a substantially higher JSC of 9.33 and an FF of 0.55, resulting in a PCE of 4.53%. Yang et al. [93] synthesized BDT34 comprising three BDT units terminated with 3-ethylrhodanine. All-small-molecule OPVs based on BDT34 with wide bandgap exhibited a favorable interconnected nanostructure with a VOC of 0.98 V and a PCE of 9.08%. In this context, for fair comparison, the photovoltaic properties of a blended film of BDT34 with PC71BM are presented in Table 4.5; this film showed a JSC of 8.87 mA·cm−2, a VOC of 1.01 V, and a FF of 0.65, with a PCE of 5.82%. In BDT35, a DTS unit was introduced as a second donor component, which resulted in a much smaller bandgap than that of BDT34 because of the higher-lying HOMO of BDT35 [94]. Despite a high JSC of 10.08 mA·cm−2, a substantially low FF of 0.51 led to a PCE of 5.05%. Replacing a DTS unit with a terthiophene unit bearing bis-octyl side chains led to BDT16, which exhibited the most redshifted absorption spectrum, indicating a narrower optical bandgap compared with those of BDT34 and BDT35 [82]. Despite a relatively lower VOC of 0.93 V because of the high-lying HOMO level in BDT16, a preferred film morphology and high and balanced charge mobilities in a BDT16-based device led to a PCE of 8.12% with a largely high FF of 0.663 and a JSC of 13.17 mA·cm−2. Compared with BDT16 with 3-ethylrhodamine as a terminal unit, BDT36 was synthesized with oxo-alkylated nitrile as a terminal unit [95]. Since sacrificing good lightharvesting ability of the 3-ethylrhodanine dye unit, the BDT36-based device showed a decreased JSC of 8.00 mA·cm−2 but maintained an FF of 0.70, leading to a PCE of 5.26% (Table 4.5).

4.5 Indacenodithiophene Derivatives As shown in Scheme 4.5, the ladder-type IDT units can be synthesized by the cyclization of two adjacent thiophenes through intramolecular annulation via sp3-hybridizied carbon atoms, where they can bear

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FIGURE 4.12  End-group modification for BDT-based oligomers.

either aromatic side chains or alkyl side chains. Compared with the aforementioned building blocks, the pentacyclic IDT unit possesses many desirable features for incorporation into semiconductors: (i) its extended ladder-type framework can elongate the effective π-conjugation length as well as reduce conformational energetic disorder, contributing to stronger intermolecular interactions and greater charge carrier mobility; (ii) the side chains introduced at the cyclopentadienyl ring can ensure solubility and control the thin-film morphology; and (iii) the choice of bridging atom can influence both the degree of aromaticity of the repeat unit and its electron density [97]. The first ladder-type IDT donor unit was reported by Wong et al. in 2006 [98]. Thereafter, various IDT derivatives followed, shedding light on their numerous advantages, which include synthetic flexibility and desirable photophysical and electronic properties. Particularly, the recent development of acceptor– donor–acceptor structural n-type oligomers utilizing IDT or its derivatives as the core and different electron-withdrawing units as end-capping groups enabled substantial breakthroughs in terms of improving the efficiency of OPVs, yielding values greater than 13%, which is considered a milestone for OPVs [99].

4.5.1 Core Unit or π-Bridge Modification In 2015, Zhan et al. first reported the small-bandgap n-type oligomer IDT1 with strong absorption in the 500–750 nm region [100]. It was synthesized using n-hexylphenyl-substituted IDT as a central building

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Conjugated Polymers: Perspective, Theory, and New Materials

TABLE 4.5  Optical, Electronic Properties, Mobilities, and Photovoltaic Properties of BDT-based Oligomers λmax (nm)a

Eg (eV)a

HOMO/ LUMO (eV)

µe (cm2V−1s−1)

µh (cm2V−1s−1)

BDT1 BDT2 BDT3 BDT4 BDT5

780 773 629 620 610

1.59 1.60 − − 1.70

−5.18/−3.56 −5.16/−3.52 −5.15/−3.44 −5.20/−3.64 −5.11/−3.39

− − 1.6 × 10−2 c 5.7 × 10−3 c −

BDT6

610

1.70

−5.21/−3.60

−

BDT7

396, 569, 618 561 532, 577 585, 636 731 760 583 583 586 591 588, 632 591, 640 589, 635 591, 639 578 − − 524, 570 528, 573 − − 724 725 655 523, 566

1.75

−5.48/−3.61

3.7 × 10−4 b

1.71 × 10 2.63 × 10−2 b 2.2 × 10−2 c 1.1 × 10−3 c 1.0 × 10−4 b 5.7 × 10−2 d 2.5 × 10−7 b 4.6 × 10−2 d 2.5 × 10−4 b

1.72 1.88 1.97 1.56 1.45 1.74 1.79 1.74 1.72 1.73 1.76 1.78 1.76 2.09 1.87 1.87 1.97 1.86 1.70 1.71 1.72 1.71 1.68 1.93

−5.41/−3.57 −5.61/−3.55 −5.11/−3.14 −5.40/−3.83 −5.36/−3.89 −5.02/−3.27 −5.08/−3.27 −5.07/− −5.02/−3.27 −5.18/−3.25 −5.06/3.29 −5.13/−3.33 −5.07/3.29 −5.15/−3.06 −5.28/−3.01 −5.31/−3.03 −5.36/−3.39 −5.34/−3.48 −5.16/−3.40 −5.19/−3.45 −5.20–3.40 −5.20/−3.56 −5.18/−3.45 −5.27/−3.34

6.61 × 10−4 b 3.24 × 10−5 c 2.78 × 10−4 c 1.38 × 10−4 c 1.58 × 10−4 c − 3.13 × 10−4 c 4.84 × 10−4 c 4.19 × 10−4 c 7.63 × 10−4 c 3.42 × 10−4 c 5.27 × 10−4 c 4.19 × 10−4 c − 1.18 × 10−5 c 5.13 × 10−4 c − − − − − − − −

1.77 1.77

−5.17/−3.34 −5.11/−3.34

2.0 1.82 1.76

−5.51/−3.34 −5.20/−3.33 −5.19/−3.46

BDT8 BDT9 BDT10 BDT11 BDT12 BDT13 BDT14 BDT15 BDT16 BDT17 BDT18 BDT19 BDT20 BDT21 BDT22 BDT23 BDT24 BDT25 BDT26 BDT27 BDT28 BDT29 BDT30 BDT31 BDT32 BDT33 BDT34 BDT35 BDT36

406, 596 372, 573, 615 545 582 585

b

FF (%)

PCE (%)

Refs.

9.47 8.58 4.66 4.60 11.70

0.91 0.92 0.92 0.83 0.76

48.2 64.8 47.0 43.0 50.1

4.15 5.11 2.01 1.62 4.40

[70] [70] [71] [71] [74]

11.20

0.84

42.7

4.00

[73, 74]

6.32

0.88

53.6

2.98

[72]

9.94 10.60 15.60 17.85 20.33 12.21 12.56 14.61 13.17 13.45 11.92 12.93 12.09 8.67 10.5 15.21 9.10 10.50 8.67 8.97 4.57 11.40 4.89 4.70

0.86 0.85 0.81 0.87 0.84 0.93 0.94 0.92 0.93 0.97 0.96 0.91 0.92 0.97 0.96 0.98 0.97 0.90 0.85 0.83 0.78 0.77 0.72 0.92

59.1 56.0 42.6 67.2 71.0 65.0 70.0 74.0 66.3 70.5 59.4 71.0 72.1 60.0 55.0 65.0 52.0 46.3 49.0 50.0 38.9 63.2 43.0 54.4

5.05 4.98 5.42 10.42 12.12 7.38 8.26 9.95 8.12 9.20 6.79 8.70 8.02 5.03 5.51 9.73 4.62 4.37 3.60 3.70 1.39 5.53 1.51 2.40

[72] [75] [76] [77] [78] [79] [80] [81] [82] [83] [82] [84] [82] [85] [96] [96] [87] [86] [88] [88] [89] [89] [90] [91]

− −

1.0 × 10−3 b 8.29 × 10−5 c 5.24 × 10−4 c 3.68 × 10−4 c 3.92 × 10−4 c 4.08 × 10−4 c 2.47 × 10−4 c 6.13 × 10−4 c 3.29 × 10−4 c 6.57 × 10−4 c 1.52 × 10−4 c 5.41 × 10−4 c 3.29 × 10−4 c 2.99 × 10−6 c 7.9 × 10−5 c 7.7 × 10−5 c 8.87 × 10−8 c 3.04 × 10−6 c 5.0 × 10−2 d 3.2 × 10−2 d 6.39 × 10−3 d 3.71 × 10−3 d 3.8 × 10−5 c 8.76 × 10−5 b 6.52 × 10−6 c 4.7 × 10−4 c 0.86 × 10−4 c

9.33 4.74

0.89 0.82

54.5 40.5

4.53 1.58

[92] [92]

2.89 × 10−4 c − −

2.64 × 10−4 c − 1.4 × 10−4 c

14.25 10.08 8.00

0.98 0.98 0.94

65.0 51.3 70.0

9.08 5.05 5.26

[93] [94] [95]

  Thin film.   Neat film mobilities measured via SCLC method. c  Blend film mobilities measured via SCLC method. d  Neat film mobilities measured via OFET method. a

JSC (mA cm−2) VOC (V)

−4 b

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169

SCHEME 4.5  Synthetic routes to the IDT unit.

FIGURE 4.13  Core or pi-bridge modification for IDT-based oligomers.

block, 1,1-dicyanomethylene-3-indanone (DC) as electron-withdrawing end groups, and thiophene as π-bridges (Figure 4.13). The IDT1 thin film exhibited good electron mobility, as high as 3.3 × 10−4 cm2·V−1·s−1. In addition, its cyclic voltammetry-derived HOMO and LUMO levels were −5.43 eV and −3.85 eV, respectively, matching the HOMO and LUMO of many classical donor materials. For example, with the PBDTT-C-T:IDT1 combination, a PCE of 3.93% was achieved without any post-treatment. On the basis of IDT1, alkyl side chains were further introduced onto the thiophene bridges, resulting in IDT2, without substantially influencing its UV-absorption and frontier energy orbitals [101]. When IDT2 was blended with low-bandgap-donor polymer PTB7-TH, the resulting nanoscale interpenetrating microstructure offered a higher PCE of 6.31% with a greatly enhanced JSC in the device. When the alkyl group was replaced with an electron-donating alkoxy group on the thiophene bridges, the solid film of IDT4 oligomer exhibited a narrow bandgap of 1.34 eV with a λmax at 785 nm, which was redshifted by ~90 nm relative to that of IDT2 [102]. As a result, the IDT4-based nonfullerene OPVs with

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PBDTTT-E-T as a donor showed a maximum PCE of 8.4% with a much higher JSC of 17.7 mA·cm−2 but a reduced VOC of 0.82 V. To avoid the VOC loss, a fused five-heterocyclic ring containing selenium atoms (IDSe) was synthesized to replace the IDT moiety in IDT2, leading to a lower bandgap of 1.52 eV as well as an upshifted LUMO level [103]. When the medium-bandgap polymer J51 (Eg = 1.91 eV) was used as a donor, the reduced energy loss (Eloss), together with complementary absorption in the vis–NIR region (350–850 nm), contributed to simultaneously high JSC and VOC, giving a PCE as high as 8.6%. Moreover, the acceptor (IDT5) with noncovalently conformational locking exhibited good planarity and rigidity in the solid state, with a delocalized LUMO level, which can also effectively suppress the nonradiative energy loss and afford higher VOC values in devices [104]. Later, a bulky seven-ring fused core (indacenodithieno[3,2-b]thiophene, IDTT) with stronger electron-donating ability, end-capped by 2-(3-​oxo-2​,3-di​hydro​inden​-1-yl​idene​)malo​nonit​rile (INCN) groups, namely IDT6, was designed and synthesized [105]. The push–pull structure in IDT6 can induce intramolecular charge transfer and extend absorption to 780 nm. In addition, the appropriate energy levels (−5.48 and −3.83 eV for the HOMO and LUMO, respectively) and good miscibility of IDT6 make it compatible with widely used donor materials. For example, the OPVs based on a PBDB-T:IDT6 blend film showed OPVs as high as 11.21%, together with excellent thermal stability, indicating its great potential in the practical application of highly efficient OPVs [106]. Because of the σ-inductive effect of thienyl side chains, the IDT7 exhibits deeper energy levels (HOMO = −5.66 eV, LUMO = −3.93 eV) than IDT6, enabling it to properly match with medium- and even wide-bandgap donor materials [107]. Moreover, the enhanced intermolecular interaction induced by sulfur–sulfur interaction offers a substantially higher electron mobility (6.1 × 10−4 cm2·V−1·s−1). The side-chain isomerization on IDT6 by alternating para-alkyl-phenyl and meta-alkyl-phenyl groups makes the IDT8 highly intermolecular self-assembling without obviously affecting either its energy levels or its absorption spectrum, contributing to an outstanding PCE of 11.77% when IDT8 is combined with J61 as a donor in devices [108].

4.5.2 Conjugation Length Extension A series of IC-nIDT-IC (1 ≤ n ≤ 3) oligomers were synthesized to investigate the effect of the fused-ring number (Figure 4.14) [109]. As the number of IDT units increases, the LUMO energy level remains similar (approximately −3.8 eV), whereas the HOMO level increases from −5.91 to −5.42 to −5.29 eV, and the absorption spectrum gradually redshifts. Compared with IDT11 and IDT12, IDT10 has more planar main-chains, stronger crystallinity, and, thus, relatively higher electron mobility. Further changing the side chains in IDT10 from alkylphenyl to alkyl leads to larger order ranges, a redshifted absorption spectrum, and improved electron transport, which helps improve the JSC and FF of OPVs. As-cast OPVs based on PM6:IDT9 displayed a high PCE of 11.9% with a JSC of 17.8 mA·cm−2 and an FF of 0.69 [110]. Notably, the device performances are insensitive to the active layer thickness (~95–255 nm) and device area (0.20–0.81 cm2), with PCEs exceeding 11%, which facilities large-scale roll-to-roll fabrication of high-performance OPVs. Using a fused 10-heterocyclic ring (inda​cenod​ithio​pheno​-inda​cenod​ithio​ phene​) as the core with DC as end groups (IDT13), the extended effective conjugation length contributes to both broadening of the absorption and the upshift of the LUMO energy level, resulting in a low Eloss of 0.59 eV and an external quantum efficiency as high as 63% in OPVs [111]. Besides, p-type oligomers based on an IDT unit as the building block also attracted considerable attention. For example, IDT14, consisting of an IDT unit as a core flanked by electron-deficient BT units and end-capped with hexyl-substituted bithiophene, was synthesized for photovoltaic applications [112]. This molecule shows broad absorption in the visible range and a low-lying HOMO energy level of −5.21 eV. The OPVs based on the blend of IDT14:PC71BM showed a PCE of 4.25%, with a high VOC of 0.93 V. To elucidate the synergism effect of conjugation length and fluorination on both the backbone conformation and physical properties of the desired materials, a series of multifluorine-substituted oligomers with an IDT unit as the electron-donating moiety and a difluorobenzothiadiazole unit as the electron-withdrawing moiety were synthesized [113]. In solution, longer oligomers exhibited a progressive

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FIGURE 4.14  Conjugation length extension for IDT-based oligomers.

redshift of the absorption maximum with an enhanced molecular absorption coefficient. As estimated by the absorption onsets of thin films, the corresponding optical bandgaps of IDT15, IDT16, IDT17, and IDT18 are 1.81, 1.79, 1.78, and 1.77 eV, respectively. The HOMO energy level was slightly upshifted with increasing chain length because of the extended conjugation length. Notably, the best PCE of 9.1% with a very high FF of 0.76 was achieved in the medium-sized oligomer IDT16-based devices with PC71BM as the acceptor.

4.5.3 End-Group Modification Recently, four new fused-ring electron oligomers based on a fused-nonacyclic IBDT (6,6,​12,12​-tetr​a kis (​4-hex​ylphe​nyl)-​indac​enobi​s(dit​hieno​[3,2-​b;2ʹ,3ʹ-d]thiophene)) core end-capped with nonfluorinated or fluorinated DC were designed and synthesized for application in nonfullerene OPVs (Figure 4.15) [114]. The larger rigid and coplanar conformation of IBDT has stronger electron-donating ability than both IDT and IDTT, which is expected to further enhance absorption range and charge transport. Substituting electron-deficient F into the end-capping DC groups results in INICn, whose LUMO energy level was gradually downshifted with increasing F-substituent number and whose absorption spectrum was redshifted to 850 nm in IDT22. In addition, the fluorine-induced intermolecular interactions lead to compact semi-crystalline stacking and thus higher electron mobilities in INICn-based blend films. The combination of the fluorinated IDT22 acceptor and the wide-bandgap donor polymer FTAZ resulted in an excellent PCE of 11.5% in OPVs; this PCE is substantially greater than that of the nonfluorinated INIC one (7.7%). The fluorinated ITIC acceptor, denoted as IDT23, was also prepared [99]. IDT23exhibited a more redshifted absorption and lower-lying HOMO/LUMO energy levels than

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FIGURE 4.15  End-group modification for IDT-based oligomers.

IDT6. To avoid the decrease in device VOC resulting from the downshifted LUMO energy level of the acceptor, a polymer named PBDB-T-SF with a low LUMO energy level was used as the donor material, resulting in a record certified PCE of 13% in devices. By contrast, the LUMO energy level of ITIC derivatives can be upshifted by incorporating electron-rich substitutions such as methyl groups into the endcapping DC groups, resulting in oligomers IDT24 and IDT25 [115]. When PBDB-T was used as a donor polymer, the IDT24-based device achieved a higher VOC value than that of the ITIC-based device. The combination of efficient exciton dissociation and reduced bimolecular recombination resulted in a high PCE of 12.05% in PBDB-T:IDT25 devices (Table 4.6).

4.6 Rylene Diimide Derivatives Rylenes are hydrocarbon compounds that can be regarded as naphthalene oligomers with bonds between the 1 and 1ʹ positions and between the 8 and 8ʹ positions of adjacent naphthalene units. When two sixmembered dicarboxylic imide rings are fused to the terminal naphthalene units, rylene diimides are formed. Interest in rylene diimides stems from early observations of electron transport behavior and the ability to fine-tune molecular electronic properties through either variation of the substituents on the rylene core or the substituents on the imide nitrogen atoms [117]. The high electron affinities and excellent chemical, thermal, and photochemical stabilities of rylene diimides makes them attractive candidates for application in OFET and OPV devices [118].

4.6.1 Conjugation Length Extension The simplest rylene diimide, naphthalene-1,8:4,5-tetracarboxylic diimide (NDI), is extensively utilized as an n-type material, providing a unique variability in structure modification and a widely tunable absorption. For example, Jenekhe et al. processed a homologous series of oligothiophene-functionalized NDI molecules (NDI1–NDI6) to fine-tune their optical and electrochemical properties and their charge carrier mobilities by virtue of intramolecular charge transfer interactions, as shown in Figure 4.16 [119]. Thin films of the oligomers had optical band gaps that varied from 2.06 eV (NDI1) to 1.57 eV (NDI3) to 1.39 eV (NDI4). The LUMO levels of the oligomers were similar at ~4.0 eV, whereas the HOMO levels acutely varied from −6.1 eV in NDI1 to −5.5 eV in NDI3 and NDI4. Spin-coated thin films were mostly crystalline and had field-effect electron mobilities as high as 2 × 10−4 to 9 × 10−4 cm2·V−1·s−1. In particular, the OPVs incorporating the NDI3 and P3HT donor polymer showed a bicontinuous nanoscale

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Advances in Discrete Length and Fused Conjugated Oligomers TABLE 4.6  Optical, Electronic Properties, Mobilities, and Photovoltaic Properties of IDT-based Oligomers λmax (nm)a IDT1 IDT2 IDT3 IDT4 IDT5 IDT6 IDT7 IDT8 IDT9 IDT10 IDT11 IDT12 IDT13 IDT14 IDT15 IDT16 IDT17 IDT18 IDT19 IDT20 IDT21 IDT22 IDT23 IDT24 IDT25

735 730 688

716 656 690 690 721 610 618 627 635 635 706 728 720 744 717 700 692

Eg (eV)a

HOMO/ LUMO (eV)

µe (cm2V−1s−1)b

µh (cm2V−1s−1)b

JSC (mA cm−2)

VOC (V)

FF (%)

PCE (%)

Refs.

1.55 1.57 1.52 1.34 1.63 1.59 1.60 1.58 1.62 1.70 1.57 1.53 1.53 1.80 1.81 1.79 1.78 1.77 1.57 1.52 1.56 1.48 1.52 1.60 1.63

−5.43/−3.85 −5.42/−3.82 −5.45/−3.79 −5.32/−3.95 −5.51/−3.78 −5.48/−3.83 −5.66/−3.93 −5.52/−3.82 −5.69/−3.91 −5.91/−3.83 −5.42/−3.80 −5.29/−3.79 −5.42/−3.82 −5.21/−3.58 −5.33/−3.17 −5.32/−3.17 −5.29/−3.18 −5.28/−3.19 −5.45/−3.88 −5.54/−3.97 −5.52/−3.98 −5.52/−4.02 −5.66/−4.14 −5.58/−3.98 −5.56/−3.93

1.5 × 10 1.0 × 10−4 7.72 × 10−5 4.6 × 10−4 4.99 × 10−4 1.1 × 10−4 2.7 × 10−4 1.30 × 10−4 2.9 × 10−4 6.1 × 10−6 9.5 × 10−5 1.1 × 10−4 4.54 × 10−5 − 1.81 × 10−4 1.06 × 10−4 1.16 × 10−4 1.03 × 10−4 2.3 × 10−5 1.1 × 10−4 9.1 × 10−5 1.4 × 10−4 4.32 × 10−4 1.10 × 10−4 4.70 × 10−5

2.0 × 10 4.5 × 10−4 8.25 × 10−5 1.5 × 10−3 5.31 × 10−4 4.3 × 10−5 2.3 × 10−4 1.54 × 10−4 5.1 × 10−5 4.2 × 10−4 2.7 × 10−5 2.2 × 10−5 − − 8.50 × 10−4 1.04 × 10−4 2.27 × 10−4 2.81 × 10−5 3.0 × 10−4 1.8 × 10−4 2.2 × 10−4 2.0 × 10−4 3.25 × 10−4 3.33 × 10−4 2.29 × 10−4

8.33 13.55 15.20 17.70 17.52 14.21 16.24 18.31 15.05 13.01 12.75 3.51 14.49 9.42 12.25 13.39 10.90 11.72 13.51 17.56 16.63 19.44 20.88 17.44 16.48

0.90 0.97 0.91 0.82 1.01 0.81 0.88 0.91 0.89 0.99 0.93 0.93 0.94 0.93 0.89 0.89 0.89 0.91 0.96 0.90 0.93 0.86 0.88 0.94 0.97

52.3 48.0 62.0 58.0 54.0 59.1 67.1 70.6 65.0 57.0 37.0 32.0 47.5 48.5 76.3 75.2 69.1 65.8 57.9 66.8 64.3 67.4 71.3 73.5 70.6

3.93 6.31 8.58 8.40 9.60 6.80 9.6 11.77 8.71 7.40 4.38 1.05 6.48 4.25 8.35 9.09 6.78 7.10 7.70 10.80 10.10 11.50 13.10 12.05 11.29

[100] [101] [103] [102] [104] [106] [107] [108] [116] [109] [109] [109] [111] [112] [113] [113] [113] [113] [114] [114] [114] [114] [99] [115] [115]

−3

−3

  Thin film.   Blend film mobilities measured via SCLC method.

a

b

morphology, resulting in a PCE of 1.5%, with a VOC of 0.82 V. To extend the planar and conjugation length, an NDI dimer (NDI7) using a vinyl linker was synthesized. Thin-film transistors of NDI7 exhibited excellent electron mobility of 0.365 cm2·V−1·s−1 in ambient atmosphere [120]. A highest PCE of 2.41% was achieved for the OPVs based on PTB7:NDI7 blends. The trimer NDI derivative NDI8 is a strained π-conjugated system that contains three naphthalene monoimide rings oriented in a three-fold symmetric pattern around a central benzene ring. The fusion of three imide groups to decacyclene lowers the LUMO level of NDI8 to −3.9 eV and results in a material that absorbs strongly in the wavelength region from 300 to 550 nm [121]. When blended with P3HT, a 1.6% PCE was achieved with a high FF of 0.57. Heterocyclic diimides, tetraazabenzodifluoranthenes (BFIs), were synthesized by the fusion of the heterocyclic ring tetraazaanthracene and two naphthalene imide units. Linking two BFI building blocks in the central tetraazaanthracene position with thiophene produces a nonplanar 3D oligomer, NDI10, which shows an overall improved performance over the monomeric analog (NDI9) in OPVs with the donor polymer PSEHTT, demonstrating a maximum PCE of 5.04% with higher and more balanced charge mobilities in devices [122]. When the arylene (Ar) linker between the bulky BFI units is varied, the twisting angle between two BFI units increases from 40° (NDI11) to 53° (NDI13) to 62° (NDI12), leading to obvious changes in electron mobility with negligible optical and electronic differences [123]. When paired with PSEHTT as a donor polymer, the NDI12-based device shows a PCE of 6.4%, with an outstanding external quantum efficiency of 80%.

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FIGURE 4.16  Conjugation length extension for NDI-containing oligomers.

Another classic rylene diimide analog is perylene-3,4:9,10-tetracarboxylic diimide (PDI), which has been frequently used in OPVs as a small molecular acceptor. However, monomeric PDI possesses a highly planar conformation and, thus, strong intermolecular π–π stackings, leading to large crystalline domain formation and severe phase separation in BHJ blends. One approach to alleviating this problem is to construct nonplanar PDI derivatives to reduce the self-aggregation effect of PDIs [13]. The dimerization of PDI moieties with (or without) functional bridges at a bay position is quite effective for developing nonplanar PDI acceptors (Figure 4.17). The different degrees of twisting and flexibility by bay-linkages results in different conjugation conformations over the whole π-system. For instance, Wang et al. reported three perylene bisimide (PBI) dyes with single linkages (PDI1–PDI3), chiral double linkages (PDI4), and graphene-like triple linkages (PDI5) [124]. The flexibly twisted oligomers (PDI1– PDI3) with an angle between two PDI units of approximately 70° demonstrated strong absorption in the range from 400 to 600 nm and appropriate energy levels that matched the low-bandgap donor polymer PBDTTT-C-T, affording a relatively higher PCE value of 3.63%, compared with those of the other two locked twisted acceptors. Encouraged by the effective dimerization at a bay position of PDI units to avoid agglomerated properties, annulation of heteroatoms in opposite bay positions of a bis-linked PDI unit was reported in PDI6 and PDI7 by S and Se, respectively [125, 126]. Owing to heteroatomic steric repulsion, dihedral angles of PDI subunits in annulated PDI derivatives are significantly enlarged up to 80°, leading to more twisted configurations compared with non-annulated ones. In conventional OPVs based on PDBT-T1 as a donor polymer, PDI7 exhibits a PCE of 8.42%, higher than 7.16% from PDI6. Nonplanar PDI conformations can also be achieved by inserting a heteroaromatic bridge between the two PDI units. With variation of the chalcogen atoms from S to O to Se, a series of semiflexible PDI dimers PDI8–PDI10 were synthesized via Stille coupling reaction between monobromo PDI and

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FIGURE 4.17  Bay-position conjugation extension for PDI-containing oligomers.

aromatic heterocyclic stannyl reagents; subsequent oxidative cyclization afforded ring-fused dimers. With increasing chalcogen atomic size (O  ​t hiop​hene ​>  sel​enoph​ene  >​  tell​uroph​ene  >​  fura​n.8 The relative reaction rates for electrophilic substitution of these heterocycles occur in the reverse order of aromaticity. For example, furan can act as a diene in Diels–Alder reactions whilst thiophene cannot.9,10 These differences in aromaticity dictate that the chemistries employed to synthesize these important precursors are varied and rich. The aromaticity of these heterocycles not only influences their reactivity, but also plays an important role in the stability and properties of the resulting conjugated polymers (see 7.4). One of the advantages of conjugated polymers is their potential for flexibility and solution or melt processability relative to inorganic semiconductors. To provide solubility, side chains are often installed on the 3-positions of chalcogenophenes. For this purpose, 3-halochalcogenophenes are attractive intermediates. There are a very limited number of reports of direct synthesis of 3-halochalcogenophenes from readily available precursors. For example, 3-bromofuran can be directly synthesized by oxidative cyclization of monobrominated diol in the presence of chromic acid in 10% yield (Scheme 7.1).11,12 3-Iodoselenophene can be synthesized via electrophilic cyclization of accessible alkynyl selenides followed by a dehydrogenation reaction using 2,3-dichloro-5,6-dicyanobenzoquinone (DDQ) to give 3-iodoselenophenes (33% yield, Scheme 7.2).13 Similarly, the iodocyclization of (Z)-tellurobutenynes by reaction with I2/petroleum ether affords the desired 3-iodotellurophenes (40% yield, Scheme 7.2).14 In this case, it has been proposed that the mechanism involves the activation of triple bonds by coordination of I2 to generate an iodonium intermediate. Aside from the low to moderate yields of these direct syntheses, another drawback is the tedious synthesis of the ringclosing precursors.

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SCHEME 7.2  Synthesis of 3-halogenated selenophene or tellurophene.

Alternatively, the synthesis of polyhalogenated derivatives followed by dehalogenation can be adopted to afford 3-halochalcogenophenes. The synthesis of 3-bromofuran can be accomplished by monolithiation of 3,4-dibromofuran and quenching by a proton source (Scheme 7.1).15 The ring-closing precursor for 3,4-dibromofuran is readily accessible due to the symmetry of functional groups compared to the aforementioned monobrominated diol. 3-Bromofuran can also be prepared by a simultaneous dehydrobromination/retro-Diels–Alder (D–A) reaction by heating in quinoline the brominated D–A adduct of furan with maleic anhydride (50% yield, Scheme 7.1).16–18 Similarly, 3-bromoselenophene can be synthesized by the reduction of 2,3,5-tribromoselenophene, in a method akin to the synthesis of 3-bromothiophene (Scheme 7.2).19,20 This synthetic route has since been optimized as a one-pot reaction.21 3-Iodoselenophene can be prepared by a tetraiodination-reduction protocol (Scheme 7.2).22,23 Unlike the lighter, group 16 analogues, there are few reports of 3-halogenated tellurophenes synthesized by this method, likely due to the instability of the heterocycle under the harsh reaction conditions necessary for halogenation. Until recently, the only reported synthesis of 3-bromotellurophene was by the selective protodeboronation of perborylated tellurophene followed by BPin/Br exchange reaction with excess CuBr2.24 It should be noted that all of the aforementioned processes suffer from low atom efficiency or the use of undesirable reagents such as mercury compounds. After obtaining 3-halochalcogenophenes, there are a variety of coupling or substitution pathways available for further derivatization at the 3-position. 3-Alkylfurans18 and 3-alkylselenophenes23 are often synthesized using this method, similar to 3-alkylthiophenes25 (Scheme 7.3). Recently, a convenient Co-catalyzed alkylation method was reported for the facile synthesis of 3-substituted furans, thiophenes and selenophenes.26 Instead of using moisture-sensitive organometallic reagents (RMgX or RZnX), this type of functionalization involves the use of various alkyl iodides as reactants. In recent years, the synthesis of 3-alkylselenophenes27,28 and 3-alkyltellurophenes29 by a ring-closing reaction has been developed, which offers improved yields. Moreover, both 3-alkylselenophenes and

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SCHEME 7.3  Synthesis and halogenation of 3-alkylchalcogenophenes.

3-alkyltellurophenes can be conveniently prepared from the same precursor (Scheme 7.3). The synthesis begins with the reaction of a commercially available Weinreb amide with Grignard reagents bearing the desired alkyl chain. A key benefit of this synthetic route is that the amide can be treated with an excess of the Grignard reagent without over substitution occurring at the resulting ketone. Thus, the Weinreb amide route29 leads to a higher yield with more facile purification. The ring-closing precursor is synthesized by treatment of the resulting α-chloroketone with ethynylmagnesium bromide. The ring-closing reaction involves the addition of the ring-closing precursor to sodium selenide or sodium telluride generated in situ,30 followed by dehydration to afford 3-alkylselenophene and 3-alkyltellurophene, respectively. Notably, the synthesis of 3-alkyltellurophenes requires an additional aromatization step using p-toluenesulfonic acid at elevated temperature while 3-alkylselenophene is afforded during the work-up by washing the organic layer with dilute hydrochloric acid. This is presumably driven by the higher aromaticity of selenophene heterocycles. Unfortunately, the synthesis of 3-alkylfurans cannot be accomplished using this route. Alternative synthetic routes for 3-alkylfurans can be found in a 1998 review by Wong et al.31 In general, 3-alkylchalcogenophenes need to be functionalized at the 2- and 5-positions for cross-­ coupling polymerizations (Scheme 7.3). Bromination by electrophilic substitution using molecular bromine or N-bromosuccinimide (NBS) is commonly adopted for furans,18,32 thiophenes,33 and selenophenes.34 Compared to thiophene, dibromination of selenophene tends to be faster, and can be accomplished using milder conditions, which is indicative of the lower aromaticity of selenophene.8 However,

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when applied to tellurophene, this method has proven to be problematic,35 probably due to the unique coordination chemistry of the tellurium heteroatom.36 This can be overcome by dilithiation of 3-alkyltellurophene followed by the addition of 1,2-dibromotetrachloroethane to afford the desired dibrominated compound.35 Similarly, when the dilithiated mixture is quenched with I2, diiodotellurophene is afforded.29 An alternative is electrophilic iodination using N-iodosuccinimide (NIS).37 Rivard and coworkers have developed a tandem synthetic strategy that combines Zr-mediated metallacycle transfer chemistry and conversion of a carbon-boron bond into a carbon-bromine bond through boronic ester intermediates (Scheme 7.3).24,38 Moreover, the direct reaction of zirconacyclopentadienes with element halides, which is also called Fagan−Nugent reaction, represents a straightforward and versatile synthetic route available for various thiophene, selenophene, and tellurophene derivatives from simple starting materials.39 The regioregularity of conjugated polymers greatly impacts properties. Due to the asymmetry of 3-alkylchalcogenophenes, specific strategies need to be employed to achieve high regioregularity. Two strategies have been developed towards this end. The first is associated closely with a specific polymerization technique, catalyst transfer polycondensation (CTP), which we will discuss in detail later. The metathesis reaction of 3-alkyl-2,5-dibromo or -diiodochalcogenophene monomers with Grignard reagents leads to a mixture of two regioisomers. However, the selective consumption of one monomer over the other leads to high regioregularity of the resulting polymers. Alternatively, one can use an asymmetrical monomer or monomer precursor. For instance, 3-alkyl-2-bromo-5-iodochalcogenophenes are commonly adopted.40 The selective activation at the 5-position leads exclusively to one monomer for polymerization. It is worthwhile to point out that the concerns regarding regioregularity are no longer relevant when symmetric monomers are used. This is exemplified by the use of fused five-membered (such as 3,4-cyclopentane,41 3,4-diimide42) or fused six-membered rings (such as 3,4-cyclohexane, Scheme 7.3). Despite synthetic challenges, most of these heterocycles are stable at low or room temperature. Furan derivatives are not stable when both light and oxygen are present, and they are very sensitive to acid. Brominated furans are especially sensitive to acidic conditions, including standard silica gel.32 For tellurophene, significant decomposition is found in reactions using acetic acid as solvent, likely due to acid-assisted ring-opening reaction.

7.3 Furan, Selenophene, and Tellurophene Homopolymers Polymerization of chalcogenophenes can be classified into two categories: electrochemical/chemical oxidative and cross-coupling. Electrochemical oxidative polymerization was historically used to synthesize polychalcogenophenes. The electrochemical syntheses of polyfurans,43 polyselenophenes,44 polytellurophenes,45 and their respective derivatives have been thoroughly reviewed elsewhere. One of the concerns, especially for polyfurans, is the high oxidation potential required for polymerization, which is not only practically inaccessible but also causes irreversible degradation.46–48 The first polyfuran was prepared via electrochemically polymerizing a wisely chosen monomer, terfuran, to lower the applied potential.49 The oxidation of chalcogenophenes can also be accomplished using oxidative reagents, such as FeCl 3. Due to the lack of solubilizing groups, these polymers are generally insoluble in common solvents (with the exception of oligofurans where the C–O dipole and weak interchain interactions enhance solubility). However, these electrochemical polymerization methods are uncontrolled. Furthermore, poor solubility inhibits the polymerization of these unsubstituted monomers or oligomers. Thus, solubilizing side chains are imperative for successful polymerization as well as solution processability. The introduction of side chains removes the symmetry of the fivemembered rings. Consequently, the repeat units must be arranged in a head-to-tail or regioregular fashion to preserve the conjugation pathway along the polymer backbone.1,50–53 Despite the progress that has been made in optimizing polymerization conditions, neither electrochemical nor chemical oxidative polymerization offers the ability to control the regioregularity or polymeric architecture of the resulting polymers.

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Cross-coupling polymerizations are based on reactions of various organometallic reagents (magnesium, zinc, tin, or boron, for example) and aryl halides (Cl, Br, I) to form new carbon-carbon bonds. This polymerization technique is usually catalyzed by transition metals such as nickel or palladium. The first application of cross-coupling polymerization was reported by Yamamoto et al. to afford poly(pphenylene).54 Since then, a deeper understanding of polymerization mechanisms and improved catalyst design has expanded the scope of applicable monomers to include the entire chalcogenophene series. Of all the organometallic polymerization techniques, the discovery of CTP is particularly important for the preparation of well-defined, conjugated polymers. CTP typically involves three steps: 1) preparing a dihaloheterocycle precursor; 2) generating an organometallic active monomer and; 3) initiating the polymerization with a catalyst.55 In 2004, the Yokozawa 56,57 and McCullough58 groups independently identified the quasi-living, chain-growth mechanism behind CTP. Evidence for the quasi-living nature of this polymerization technique includes: 1) linear increase of number average molecular weight (Mn) as a function of monomer consumption; 2) low dispersity; 3) one combination of polymer end groups (H/ Br); 4) linear correlation between molecular weights with monomer-to-catalyst ratios and; 5) the ability to perform chain extension yielding block copolymers by sequential monomer addition. The developments in CTP not only enable the synthesis of well-defined polymers with reproducible Mn, low dispersities, and end-group control but also enable access to more complex polymeric architectures, including gradient and block copolymers.37,59–61 As previously alluded to, monomer reactivity differs greatly when traversing group 16, thus the application of CTP to each of these polymers has been divided into short sections below.

7.3.1 Preparation of Polyfurans Polyfurans are often cited as greener alternatives to polythiophenes, as furan is an organic, biodegradable monomer which can be obtained from renewable sources.62 Moreover, incorporation of O in oligofurans imparts greater solubility, less weight, increased rigidity, and enhanced fluorescence relative to S, Se and Te-containing oligomers.10,63,64 The main challenge associated with accessing high molecular weight polyfurans is their susceptibility to decomposition,65 which makes monomer synthesis difficult and destabilizes the resulting polymer.32,43 Bendikov and coworkers have reported their heroic efforts to synthesize stable and highly fluorescent, monodisperse oligo(3-alkylfuran) up to the 16-mer via sequential cross-coupling reactions.64,65 These oligomers were found to be stable when stored for several weeks under ambient conditions and maintained stability up to 250 °C in the case of the 9-mer and 180°C for the 16-mer by thermogravimetric analysis (TGA). These results contradict the implication that oligofurans above the 4-mer are unstable66 and demonstrate that derivatives of these important materials can be stable. In 2016, Noonan and coworkers reported the first synthesis of regioregular poly(3-hexylfuran) (HT-P3HF) alongside the regiosymmetric analogue (HH-P3HF) with low dispersities (Ð = 1.2–1.3) and, albeit low, molecular weights (Mn = 2.9–4.1 kDa) via Kumada Catalyst Transfer Polycondensation (KCTP) of HT- and HH-coupled furan dimers (Scheme 7.4).32 Notably, P3HF was shown to adopt the same conformation in solution as the solid-state, irrespective of the degree of regioregularity, thus demonstrating the enhanced rigidity of the polymer backbone relative to polythiophenes (see Section 7.4). This planarity arises from the greater quinoid character of polyfurans relative to polythiophenes (see Section 7.4). To date, it has not been possible to synthesize high molecular weight poly(3-hexylfuran) homopolymers due to aggregation of the propagating chains during polymerization, which leads to premature chain termination, as well as strong association of nickel with the furan monomer, which hinders chain propagation.32 To overcome these difficulties, incorporation of furan into alternating furan-thiophene copolymers (P3HF-alt-P3HT) was accomplished by KCTP of HT furan-thiophene dimers, which yielded polymers with higher molecular weights (Mn up to 12 kDa) (Scheme 7.4).32 Interestingly, these polymers all exhibit similar nanofibrillar morphologies in thin films, suggesting that furan incorporation has little effect on solid-state packing in these materials. Unfortunately, it was noted by the authors that all P3HF-containing polymers were sensitive to decomposition by light and oxygen exposure.

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SCHEME 7.4  Synthesis of HT-P3HF, HH-P3HF, and P3HF-alt-P3HT. Ni(dppp)Cl2: [1,3-bis(diphenylphosphino) propane] dichloronickel(II). Reproduced with permission from reference 32. Copyright 2016 American Chemical Society.

Despite possessing many advantageous properties, polyfurans has been hindered in applications to organic electronics by their instability and wide HOMO-LUMO (the highest occupied molecular orbital-the lowest unoccupied molecular orbital) gaps. Conjugated polymers with narrow HOMOLUMO gaps are desirable for applications such as organic photovoltaics as more energy can be absorbed from the solar spectrum with more red-absorbing chromophores. While there are several ways to narrow the HOMO-LUMO gap of conjugated polymer, many come at the expense of increasing the HOMO level energy. This increases the oxidation potential and decreases the possible open circuit potential, which ultimately decreases the ambient stability and efficiency of the device. Thus, any modifications to a conjugated polymer should not come at the expense of raising the HOMO level. Heavier group 16 atom heterocycles have narrower HOMO-LUMO gaps than their light analogues. This, combined with the observation that the S atom in thiophene contributes very little to the electron density at the HOMO compared to the LUMO,67,68 has made heavy atom substitution of the S atom widely pursued as a strategy to improve performance of organic electronics by selectively lowering the LUMO position. Low LUMO levels also facilitate n-channel transport in organic transistors. For these and other reasons, the controlled synthesis of Se- and Te-containing polymers has been developed.29,34,37,38

7.3.2 Preparation of Polyselenophenes Of all the group 16 heterocycles, the requirements for the polymerization of thiophene and selenophene are the most similar. The synthesis of high molecular weight (Mn up to 200 kDa), regioregular poly(3-hexylselenophene) (HT-P3HS) by CTP can be accomplished in an analogous fashion to that of poly(3-alkylthiophenes)56,58 and was first reported by Heeney and coworkers in 2007 (Scheme 7.5).34 The authors found that the regioregularity of P3HS obtained by this method was greater than 97% by 1 H-NMR integration of the methylene proton peak associated with HH coupling to that associated with HT coupling. Moreover, as the selenium-77 isotope is abundant (ca. 8%) and has a spin of 1/2,

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SCHEME 7.5 Synthesis of poly(3-hexylselenophene). Ni(dppe)Cl2: [1,2-bis(diphenylphosphino)ethane] dichloronickel(II). Reproduced from reference 34 with permission from The Royal Society of Chemistry.

the authors demonstrated that regioregularity could be characterized by 77Se-NMR, which contained a single resonance indicative of high regioregularity. Significantly, in contrast to HT-P3HF, HT-P3HS was shown to be more stable to UV and temperature than HT-P3HT when prepared using analogous conditions. However, these initial studies reported polydispersities of around 2, which indicates a lack of control of the polymerization. This is likely due to the reduced solubility of polyselenophenes relative to their lighter heteroatom analogues, as the increased polarizability of Se leads to strong interchain interactions. A similar trend has been noted for polytellurophenes (see below). In recent years, subtle modifications have made the polymerization of well-defined, high molecular weight, regioregular poly(3-alkylselenophenes) (HT-P3AS) possible.69 These include the use of less sterically bulky Grignard metathesis reagents;70 monomers with longer and branched side chains to improve solubility and purification;28,71 and THF-soluble nickel initiators which block one end of the propagating polymer chain to reinsertion.59,72,73 These approaches have yielded HT-P3AS with dispersities as low as 1.2. The improved control of this polymerization has paved the way for the synthesis of thiophene-selenophene copolymers with complex architectures and unprecedented control over properties (see 7.5).28,59–61,70,71,74,75

7.3.3 Preparation of Polytellurophenes Oxidative polymerization is the most common way in which polytellurophenes have been synthesized. The most significant hurdle to accessing polytellurophenes is a lack of synthetic routes to prepare monomers suitable for mild and efficient polymerizations. In 2013, the Seferos group reported the first series of poly(3-alkyltellurophene)s prepared by both electrochemical and chemical polymerization methods.29 These polymers have moderate molecular weights (Mn = 5.4–11.3 kDa) and broad distributions (Đ = 1.9–2.2) but, critically, were solution processable. High degrees of regioregularity were confirmed by integration of the two methylene protons (>93%) corresponding to HH and HT couplings in the 1HNMR spectra. The optical HOMO-LUMO gaps determined by onset of absorption of poly(3-hexyltellurophene) (P3HTe), poly(3-dodecyltellurophene) (P3DDTe), and poly(3-(2ʹ-ethylhexyl)tellurophene) (P3EHTe) are 1.44 eV, 1.44 eV, and 1.57 eV, respectively, which is consistent with the theoretical prediction that polytellurophenes possess a narrower HOMO-LUMO gap than polythiophenes and polyselenophenes (Figure 7.2). When a branched side chain is adopted, a blue-shift absorption is observed, likely due to the increased degree of twisting of the polymer backbone resulting from the bulky side chains. P3HTe, P3DDTe, and P3EHTe have molar absorptivities of 3900, 5100, and 6400 M−1⋅cm−1 (calculated per repeat unit), respectively, revealing that all three polymers are strong light-absorbers. However, these CTP conditions did not lead to well-defined polytellurophenes. For example, Mn of the polymers were relatively low and the broad dispersities are indicative of uncontrolled polymerizations. According to the proposed mechanism of CTP, the catalytic cycle proceeds by a Ni-π complex.56,58 The association complex between the catalyst and π-system of the growing polymer chain is critical for controlled CTP. According to the prediction of gas-phase, single-point energy (SPE) DFT calculations,

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FIGURE 7.2  Calculated [PBC/B3LYP/6-31G(d)] HOCO (Highest Occupied Crystal Orbital) and LUCO (Lowest Unoccupied Crystal Orbital) energy levels of different polymers. (Reproduced from Patra, A. and Bendikov, M. Polyselenophenes. J. Mater. Chem. 2010, 20(3), 422–433. With permission from The Royal Society of Chemistry.)

the tellurophene-catalyst complex has similar stability upon complexation to the thiophene-catalyst complex,37 which implies that comparable control should be possible for 3-alkyltellurophenes. Seferos and coworkers concluded that solubility, rather than any other factor, was limiting the polymerization. Moving down group 16, the solubility of the resulting polymers decreases significantly when an identical side chain is adopted because of increasing polarizability of the heteroatoms on increasing size. Since both short and long linear side chains are not able to afford sufficient solubility, a branched side chain is imperative to successful control of the tellurophene polymerization. However, when the widely used 2-ethylhexyl side-chain was adopted, slow kinetics and early termination were observed, likely due to the increased steric hindrance. To tackle this problem, Seferos and coworkers adopted a tellurophene monomer with a 3-ethylheptyl side chain. With the branching point one carbon further away from the tellurophene, the effect of steric crowding was decreased while maintaining sufficient solubility. Detailed kinetic studies show this subtle change leads to a 3.4-fold improvement in the kinetics and controlled polymerization of 3-alkyltellurophenes. The Mn increases linearly as a function of monomer conversion without deviation, and dispersity remains under 1.2 throughout the polymerization (Figure 7.3). Further evidence for successful CTP can be seen in the molecular weights control experiments. Specifically, the Mn of resulting polymers increase from 5.2 kDa to 24.9 kDa, which is proportional to monomer-to-initiator ratio (Figure 7.4). Additionally, when the polymerization starts with an o-tolyl-functionalized Ni complex as the external initiator73 and is quenched with ethynylmagnesium bromide,76 one main population of end groups (o-tolyl and ethynyl) was observed in the matrix-assisted laser desorption/ionization time-of-flight (MALDI-ToF) spectrum. Results from kinetics studies, molecular weight control experiments, end-group analysis, self-extension, and block copolymerizations all indicate that 3-alkyltellurophenes can undergo CTP with significant living character. It should be noted that Ni(dppe)Cl2, which is similar in structure to Ni(dppp)Cl2, was used to catalyze the polymerizations of alkyltellurophenes. The two catalysts are interchangeable and give comparable polymerization results in most synthesis of poly(3-alkylchalcogenophenes). However, Ni(dppe)Cl2 tends to be more active for monomers bearing branched side chains.28 The branch in the side chain increases the monomer bulkiness and slows the polymerization kinetics. The higher activity of Ni(dppe)Cl2 may originate from the smaller bite angle and smaller chelate ring of the dppe ligand. To date, this is the only

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FIGURE 7.3  (a) Semilogarithmic kinetic plots of 3-ethylheptyltellurophene monomer consumption as a function of polymerization time. (b) SEC elution profiles for each aliquot. (c) Number average molecular weight (black) and dispersity (grey) as a function of monomer conversion for 3-ethylheptyltellurophene carried out with a 1:50 catalyst:monomer ratio at 0.1 M monomer concentration. (Reprinted from Ye, S., et al. What limits the molecular weight and controlled synthesis of poly(3-alkyltellurophene)s? Macromolecules 2016, 49(5), 1704–1711. With permission. Copyright 2016 American Chemical Society.)

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FIGURE 7.4  (a) SEC elution profiles of poly(3-ethylheptyl)tellurophene prepared with different monomer:catalyst ratios. (b) Number average molecular weight (black) and dispersity (grey) as a function of 3-ethylheptyltellurophene monomer:catalyst ratio. (Reprinted from Ye, S., et al. What limits the molecular weight and controlled synthesis of poly(3-alkyltellurophene)s? Macromolecules 2016, 49(5), 1704–1711. With permission. Copyright 2016 American Chemical Society.)

SCHEME 7.6  Aromatic and quinoidal resonance structure of polyselenophene.

report of controlled synthesis of polytellurophenes and the first example to show that the CTP kinetics can be adjusted by side chain branching point engineering. The controlled polymerization of tellurophene will allow more complex polymeric architectures involving other monomers to be synthesized. As alluded to in the previous sections, heteroatom substitution has a drastic effect on the optoelectronic properties of conjugated polymers. Experimentally demonstrating this has been facilitated by the recent advent of synthetic protocols to control the polymerization of these heterocycles as regioregularity and molecular weight also have significant effects on the resulting optoelectronic properties.50,53,77–80

7.4 Properties and Applications of O, Se-, and Te- Polymers 7.4.1 Structure and Rigidity Polyfurans, polyselenophenes, and polytellurophenes are more rigid than polythiophenes owing to their reduced aromaticity. As a consequence, these polymers exhibit more quinoid character (Scheme 7.6) than polythiophene, which can be seen with shorter interring C-C distances and greater energy penalties for twisting the polymer backbone (Table 7.1). Consequently, polyfurans, polyselenophenes, and polytellurophenes are more planar than polythiophenes, which has a marked impact on properties such as their solubility, packing behavior, and charge transport.2,35,68,81,82 For example, alkyl substituents and HH defects, which have a considerable effect on twisting in oligo- and polythiophenes, have minimal effect on the planarity in a 16-mer of oligofuran.64 This has also been observed directly by scanning tunneling microscopy of poly(3-dodecylthiophene) (P3DDT) and poly(3-dodecylselenophene) (P3DDS) monolayers, which has revealed very little chain-folding in P3DDS even at high degrees of polymerization (N ≈ 50) (Figure 7.5),71 and by Raman spectroscopy for P3HS83 and P3HTe,84 which showed higher frequency vibrational peaks corresponding to shorter C=C bond lengths as the size of the heteroatom increases. Compared to analogous polythiophenes, polyselenophenes and polytellurophenes are less soluble and melt at higher temperatures. This is largely due to the increased size of Se and Te, which leads to

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TABLE 7.1  Structural Characterization of Poly(3-hexylchalcogenophene)s

HT-P3HF HT-P3HT HT-P3HS HT-P3HTe

C-C Bond Lengtha (Å)

(100) distance (Å)

(010) distance (Å)

χb of Heteroatom

Twisting Potential (kcal/mol)

1.4322,65 1.44168 1.43468 unknown

unknown 15.7–17.080,86 15.2–15.534,69 12.429

unknown 3.8–3.980,86 ~4.234 unknown

3.44 2.58 2.55 2.10

12.510 2.310–2.5c,81 3.4c,68 unknown

  From X-ray data.   χ: electronegativity. c  At 35° in the unsubstituted 6-mer except values marked by an asterisk which were calculated for a 36° twist. a

b

FIGURE 7.5  TEM and STM images of P3DDT50 and P3DDS50 (where subscripts denote degree of polymerization) homopolymer fibers demonstrating a lack of chain-folding in the latter. Adapted from Reference 71 with permissions from Wiley.

larger π-stacking (010) plane (Table 7.1) distances. Because the π-stacking plane is pushed further apart, polyselenophene and polytellurophene pack more closely in the (100) plane. These properties contribute to the distinct crystallization behavior of polyselenophenes and polytellurophenes in contrast to polythiophenes.69,71 P3ATs with longer alkyl chains are known to adopt a second crystal polymorph (type-II) that is characterized by shorter interchain distances arising from alkyl chain interdigitation and tilting of the polymer main chains relative to one another and an accompanying hypsochromic shift in the absorption spectrum.69,85,86 The type-II polymorph is kinetically trapped and equilibrates to the type-I on annealing and has been observed experimentally in various different P3ASs (Figure 7.6).69–71 Contrary to P3HS, only type-I phases are observed in both high (19 kDa) and low (3.5 kDa) molecular weight P3HT under analogous experimental conditions. One would expect similar crystallization behavior in polytellurophenes where larger π-stacking occurs.

7.4.2 Optoelectronic Properties Polyfurans are more planar, more soluble, and exhibit higher fluorescence than polythiophenes.2,32,65 In general, the heavier the heteroatom, the higher the propensity for intersystem crossing5,87,88 and, thus, polyfurans have the highest quantum yield among the group 16 heterocyclic polymer series. Polyfurans have a wider band gap than analogous polythiophenes and a lower ionization potential, which arises from a higher HOMO level (Table 7.2, Figure 7.2). In polythiophene, extension of the

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FIGURE 7.6  Wide-angle X-ray scattering of P3HS films prepared from different molecular weight samples before (a) and after (c) annealing at 140°C. The insert in c shows the magnification of Mn = 5.9 kg mol−1 P3HS from 3° to 10°. Scattering of an un-annealed P3HT film (d). Proposed supramolecular structure of the type-I and type-II polymorphs (b). (Reprinted from Li, L., et al. Polyselenophenes with distinct crystallization properties. Chem. Sci. 2011, 2(12), 2306–2310. With permission from The Royal Society of Chemistry.) TABLE 7.2  Optoelectronic Properties of Poly(3-hexylchalcogenophene)s

HT-P3HF HT-P3HT HT-P3HS HT-P3HTe

λmax solution (nm)

λmax film (nm)

Eg Optical a (eV)

Eg CV (eV)

46532 45077 49934 55829

48932 61577 63034 612–69029,84

2.1932 1.934 1.634 1.4429

unknown 2.234 1.934 1.3729

  The optical band gaps are calculated from the onsets of the thin-film absorption.

a

conjugation pathway caused by ordering in the solid-state causes a significant bathochromic shift from the solution absorption profile, as well as the appearance of vibronic bands corresponding to planarization and π-π stacking in the solid-state.77,89 As polyfurans are more planar than polythiophenes, they adopt a similar conformation in solution to the solid-state which is reflected in the similarity in absorption and vibronic features between solution and thin film.32 The relative planarity of polyfurans has been confirmed by computationally calculating the twisting enthalpy of the sexifuran

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(6F) backbone at 35°, which is an order of magnitude higher than that of sexithiophene (6T) and sexiselenophene (6Se) (Table 7.1), and has been confirmed from the C-C bond distance from Raman spectroscopy and X-ray data for oligomers.10 In contrast, polyselenophenes and polytellurophenes have lower band gaps than polyfurans and polythiophenes (Table 7.2), as the majority of the LUMO coefficient exists on the heteroatom.67,68 Therefore, the absorption spectra of P3AS and P3ATe are generally bathochromically shifted with respect to the lighter group 16 analogues, with P3ATes being the most red-shifted. The exception to this rule is the aforementioned type-II polymorph that is prevalent in P3AS. As previously mentioned, these heavier group 16 polymers have reduced fluorescence than the lighter analogues due to increased intersystem crossing.

7.5 Furan, Selenophene, and Tellurophene Copolymers and Self-Assembly Behavior The recent development of controlled synthetic routes to poly(3-alkylchalcogenophene) block copolymers makes them viable candidates for self-assembly studies of double-crystalline polymers. To date, such studies on double-crystalline block copolymers have been limited. However, pioneering work in this area has indicated that unique nanostructures can be fabricated with exquisite control in solution or the solid-state from the crystallization-driven self-assembly of copolymers with a crystallizable component.90–96 Differences in electron density and phase contrast between P3ATs, P3ASs, and P3ATes facilitates the use of imaging techniques such as transmission electron microscopy (TEM) and atomic force microscopy (AFM) to characterize morphology. Complementary X-ray scattering studies can also be utilized to provide additional molecular-level, structural insight. The pseudo-living nature of the polymerization of group 16 heterocycles also allows quenching of polymer chains with Grignard reagents to access polymers with one set of functional end groups which can be used to synthesize block copolymers via various click chemistries.37,76,97 The similarity in reaction conditions for KCTP of furans, selenophenes, and tellurophenes also allows access to copolymers of these monomers. Despite their structural similarities, poly(3-alkylthiophene)-block-poly(3-alkylselenophene) (P3AT-b-P3AS) copolymers undergo phase separation in thin films at comparatively low molecular weights, a phenomenon termed “heterocycle-induced phase separation”.23,28,70 This behavior is presumably crystallization-driven and is facilitated by differences in rigidity between the thiophene and selenophene blocks as well as the enthalpy gain from forming pure thiophene or selenophene domains. Interestingly, the relative block ratios were found to have little impact on morphology in P3AT-b-P3AS thin films; however, some differences in self-assembly were noted when the side chains were altered.28,70 The architecture of copolymer backbones can also have a significant impact on solid-state ordering, which will affect both mechanical and charge transport properties. To probe this, McNeil and coworkers have prepared a series of compositionally similar P3HT and P3HS copolymers with block, gradient, and random architectures (Figure 7.7).98 Despite morphological similarities, phase separation was found to be greatest in the block copolymer thin films and weakest in the random copolymer films with the gradient copolymer exhibiting intermediate absorbance, thermal behavior, and crystallinity. Significantly, in photovoltaic devices, P3HT-grad-P3HS was shown to have an improved stability and formed large, interconnected fibers with greater donor:acceptor interfacial area, as well as enhanced initial carrier density relative to the block copolymer analogue, 59 thus demonstrating the importance of composition and solid-state structure for application in such devices. To date, there have been no systematic studies of polyfuran block copolymer self-assembly because of their limited stability (see previous discussion). However, in P3HF-alt-P3HT thin films, a nanofibrillar morphology, similar to that found in P3HT films, was observed, indicating that the incorporation of furan has little effect on morphology.32,75 It should be noted that phase separation was not characterized in these samples.

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FIGURE 7.7  Block, gradient, and random structures of polyselenophene-block-polythiophene copolymers. (Reprinted from Wang, J., et al. A facile way to prepare crystalline platelets of block copolymers by crystallization-driven self-assembly. Polymer. 2013, 54(25), 6760–6767. With permission. Copyright 2012 American Chemical Society.)

7.6 Summary and Outlook Despite their structural similarities, the chemistry and properties of polychalcogenophenes differ greatly by navigating group 16. The development of methodologies which enable the controlled synthesis of furan, selenophene, and tellurophene homopolymers and copolymers has led to a surge in the field of solution processable, polythiophene analogues. Controlled CTP allows one to access high quality samples of these previously unexplored polyheterocycles, fine tune properties by heteroatom substitution, and control polymer composition and sequence. While much work lies ahead in the study and application of these new linear polymers, one can only imagine that these synthetic advances will soon lead to group-16 polymers with even more complex architectures.

References 1. Jeffries-El, M.; McCullough, R. D. Regioregular polythiophenes. In Handbook of Conducting Polymers; Skotheim, T. A., Reynolds, J. R., Eds.; CRC Press: New York, 2006; pp 349–398. 2. Sharma, S.; Zamoshchik, N.; Bendikov, M. Polyfurans: A computational study. Isr. J. Chem. 2014, 54(5–6), 712–722. 3. Patra, A.; Bendikov, M. Polyselenophenes. J. Mater. Chem. 2010, 20(3), 422–433. 4. Hollinger, J.; Gao, D.; Seferos, D. S. Selenophene electronics. Isr. J. Chem. 2014, 54 (5–6), 440–453. 5. Carrera, E. I.; Seferos, D. S. Semiconducting polymers containing tellurium: Perspectives toward obtaining high-performance materials. Macromolecules 2015, 48(2), 297–308. 6. Jeffries-El, M.; Kobilka, B. M. B.; Hale, B. J. B. Optimizing the performance of conjugated polymers in organic photovoltaic cells by traversing group 16. Macromolecules 2014, 47(21), 7253–7271. 7. Parke, S. M.; Boone, M. P.; Rivard, E. Marriage of heavy main group elements with π-conjugated materials for optoelectronic applications. Chem. Commun. 2016, 52(61), 9485–9505. 8. Katritzky, A. R.; Ramsden, C. A.; Joule, J. A.; Zhdankin, V. V. Reactivity of five-membered rings with one heteroatom. In Handbook of Heterocyclic Chemistry; Elsevier Ltd.: Amsterdam, 2010; pp 383–472. 9. Cappe, O. C.; Murphree, S. S.; Padwa, A. synthetic applications of furan diels-alder chemistry. Tetrahedron 1997, 53(42), 14179–14233. 10. Gidron, O.; Bendikov, M. α-oligofurans: An emerging class of conjugated oligomers for organic electronics. Angew. Chem. Int. Ed. 2014, 53(10), 2546–2555. 11. Kraus, G. A.; Wang, X. An improved synthesis of 3-substituted Furans from substituted butene1,4-diols. Synth. Commun. 1998, 28(6), 1093–1096.

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8 Donor-Acceptor Polymers for Organic Photovoltaics 8.1 Introduction.......................................................................................283 8.2 Donor-Acceptor Conjugated Polymers..........................................284

Desta Gedefaw and Mats R. Andersson

Fluorene, Silafluorene, Carbazole, and CyclopentadithiopheneContaining Donor-Acceptor Polymers • Thiophene and Derivatives as a Donor Unit in Donor-Acceptor Polymers • Benzodithiophene as a Donor Unit for the Synthesis of Donor-Acceptor Polymers • Indacenodithiophene and Its Derivatives as a Donor Unit in the Construction of Donor-Acceptor Polymers • Summary and Outlook

Acknowledgments......................................................................................... 316 References........................................................................................................317

8.1 Introduction Conjugated polymers are materials that consist of alternating double and single bonds in their chain, which allows the movement of π-electrons in the backbone and gives a semiconductor behavior. Some of the most common classes of conjugated homo polymers include polyacetylene,1 poly(p-phenylene) (PPP),2 polypyrrole (PPy),3 poly(p-phenylenevinylene) (PPV),4 polythiophene (PT),5 polyaniline ((leucoemeraldin),6 and polyfluorene.7 One of the challenges to work with these conjugated polymers is the lack of easy processability as their solubility in organic solvents is limited. These kinds of polymers can only be prepared directly onto the substrate used for fabrication of the devices. The attachments of alkyl side chains on these structures by chemical bonding will impart solubility of the materials in organic solvents and enable an easy solution processing of the polymers into films. For instance, a solution of the polymers in organic solvents can be processed by different processing techniques such as spin coating, doctor-blading, and even printing on flexible substrates to produce the polymer films and subsequently complete the device fabrications. In addition to enabling solution processability, the side chains can also play a very important role in tuning the properties of the polymer, such as by modifying the polymer chain solid-state packing and crystallinity behaviors. Figure 8.1 shows few examples of thiophene-based homopolymers with different side chains attached to the backbone. The side chains affect the optical properties of the polymers as an example. Poly(3-octylphenyl)thiophene (POPT) is a purple-blue polymer at solid state, but replacing the side chain with cyclohexane ring changes the color of the polymer into yellow, and adding methyl group at carbon 4 of the thiophene ring in addition the cyclohexyl side group turns the polymer colorless.8 The color change arises from the change in the conformation of the polymer chains, such as the twisting of the backbones due to the effect of the side chains. Thus, careful selection of side chains is needed to make polymers with desired properties in addition to enhanced solution processability. 283

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FIGURE 8.1  Effect of solubilizing side chains on polythiophenes.

FIGURE 8.2  Structures of some electron donors (R = alkyl or aryl group).

8.2 Donor-Acceptor Conjugated Polymers Undoubtedly, the donor-acceptor (D–A) design strategy developed by Havinga et al.9 in the early 1990s is one of the most frequently used methods for the synthesis of conjugated polymers for optoelectronic applications. Properly functionalized electron-rich (donor) aromatic building blocks such as thiophene, bithiophene, fluorene, carbazole, dibenzosilole, benzodithiophene, etc. (Figure 8.2) are combined with functionalized electron-deficient (acceptor) aromatic units such as quinoxaline, benzothiadiazole (BTD), diketopyrrolopyrrole, isoindigo, and others (Figure 8.3) through metal-catalyzed cross-coupling reactions (Stille, Suzuki, Heteroarylation) to form D–A conjugated polymers.10 These types of polymers are sometimes called push-pull conjugated polymers. The D–A design motif promotes partial charge transfer between the donor and the acceptor moieties, enhancing the double bond character between the electron-rich and electron-poor units. This charge transfer decreases the bond length alternation in the repeating unit and reduces the bandgap11,12 of the polymer. For example, in the benzo-bis-(thiadiazole) system depicted in Figure 8.4, the bandgap of the material is reduced due to the partial charge transfer between the donor thiophene and the electron acceptor bis-thiadiazole units. Figure 8.5 shows a diagrammatic representation of the molecular energy levels and bandgap of the D–A polymer vis-à-vis the energy levels of the donor and the acceptor moieties.13 The strength of the electron-donating and electron-accepting nature of the donor and acceptor, respectively, has paramount importance in determining the energy levels and bandgap of the D–A polymer. The D–A polymer will have a highest occupied molecular orbital (HOMO) energy level closer to the HOMO of the donor while the lowest unoccupied molecular orbital (LUMO) of the polymer will be

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FIGURE 8.3  Structures of some electron acceptors (R = alkyl or aryl group).

FIGURE 8.4  Resonance structures in a benzo-bis-thiadiazole derivative.

closer to that of the acceptor. The use of a strong donor raises the HOMO of the D–A polymer, and the use of strong acceptor pushes the LUMO of the D–A down, narrowing the band gap of the polymer. In connection with bandgap engineering, Reynolds and coworkers studied materials that possess hetero bridging atoms and their effect on the bandgap of the polymer.14 Some of the materials included in the study are shown in Figure 8.6. The benzotriazole (BTz) unit is not a strong electron-withdrawing group because of the electrondonating effect of the central nitrogen to the ring through resonance.14 Hence, materials that possess BTz are known to have raised LUMO, giving a wider bandgap material. On the other hand, BTD containing polymers (the electron-withdrawing effect of the sulfur atom outweighs the electron-donating effect toward the ring) will give lowered LUMO.14 At the same time, BTz causes high bandgap materials, and

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LUMO LUMO

Bandgap (D)

Bandgap (D-A)

LUMO

Bandgap (A)

HOMO HOMO

Donor (D)

HOMO

D-A

Acceptor (A)

FIGURE 8.5  Representation of D–A orbital mixing during chemical linkage.

FIGURE 8.6  Examples of molecular structures with a heteroatom.

the fact that it possess an alkylatable nitrogen atom is an advantage to produce organic solvent soluble materials. On the other hand, BTD has no alkylatable atom and hence, in most cases, side-chaincontaining groups need to be attached to the BTD unit to promote solubility, which will partly complicate the synthetic process. A hybrid structure of BTD and BTz known as BTzTD,14,15 though, has an intermediate electron-withdrawing property for the synthesis of low bandgap polymers besides its alkylatable structural feature. The BBT aromatic ring possesses an even stronger electron-withdrawing property for the synthesis of low bandgap polymers, but it suffers from its non-alkylatable nature and hence it necessitates the use of alkyl substituted thiophenes for solubilizing the material in organic solvents. In order to understand the effect of the heteroatoms in materials that possess these molecular units, gas phase theoretical calculation was performed on model molecules that consist of thiophene sandwiched BTz, BTD, BTzTD, and BBT. As expected, the BTz-based molecule resulted in a wide bandgap with deeper HOMO and raised LUMO energy levels as compared to the rest of the model molecules. In fact, the LUMO of BTz was found to be raised by 1.1 eV as compared to that of the BTzTD-based model molecule, indicating the effect of the bis-thiadiazole unit in lowering the LUMO energy levels. The BBT-unit-containing molecule had the lowest LUMO energy level of all model molecules, while the HOMO energy level of this molecule was found to be comparable with that of the BTD- and BTz-based molecules. Meanwhile, compared to BTzTD, it was down shifted. When it comes to the synthesis of D–A polymers, donor (Figure 8.2) and acceptor (Figure 8.3) moieties are covalently linked together by using transition-metal-catalyzed cross-coupling reactions. The most frequently used metal-catalyzed reactions are Suzuki-Myaura, Stille, and direct arylation. The Suzuki coupling reaction requires a diborolane or boronic acid and dibromo functionalized monomeric units (Scheme 8.1a). On the other hand, Stille coupling reaction is a carbon-carbon bond-forming reaction

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SCHEME 8.1  Suzuki (a), Stille (b), and direct arylation (c) reactions.

between stannyl functionalized monomers and bromine functionalized aromatic units (Scheme 8.1b). An emerging, facile, and environmentally friendly polycondensation method for the synthesis of D–A polymers is the direct arylation method (Scheme 8.1c).16–21 The direct arylation method has attracted growing interest as it reduces synthetic steps and avoids toxic chemicals and wastes, which are often incorporated in the target molecules as impurities. Nevertheless, direct arylation suffers from the drawbacks of side reactions such as branching, cross-linking, and homo-couplings, which are responsible for structural defects. As discussed above, the versatility of the D–A approach led to the synthesis of a wide range of materials that have been subsequently used to fabricate photovoltaic devices. Through a continuous collaborative research effort, the current device power conversion efficiencies based on D–A polymer solar cells rose higher than 13% in single junction devices.22 Clearly, this is an impressive achievement of the field of organic solar cells and heralds its bright future toward commercialization for practical applications of energy production. However, in the literature, ubiquitous materials have been developed in the past three decades. This chapter will only focus on discussing some of these materials to showcase the advances made in the past few decades. The topics are divided into subsections based on the type of donor units used to construct the polymers.

8.2.1 Fluorene, Silafluorene, Carbazole, and CyclopentadithiopheneContaining Donor-Acceptor Polymers Polyfluorene-based D–A conjugated polymers have been extensively studied and incorporated in polymer solar cells and light emitting diodes due to their high fluorescence quantum yields, excellent holetransporting properties, good film-forming properties, and exceptional chemical stability.23,24 The fact that fluorene and its derivatives have large bandgaps and low-lying HOMO energy levels make them stable to photo degradation and thermal oxidation. Besides the stability, the synthetic steps towards functionalization of a fluorene unit is straightforward. The hydrogens at C-9 of the fluorene ring are significantly acidic and therefore can be removed with bases. This allows for the facile introduction of alkyl groups, which will impart high solubility of the corresponding polymers in organic solvents. In addition, fluorene can be easily brominated at the C-2 and C-7 positions. This will provide functional handles for the transformation of fluorene-containing monomers into boronate esters or boronic acids, both of which are suitable for metal-catalyzed cross-coupling reactions. The synthesis of the borolanefunctionalized fluorene is shown in Scheme 8.2,25 which then is combined with different acceptor moieties to prepare D–A polymers. Alternating polyfluorenes (APFOs) are materials that consist of fluorene as a donor unit coupled with acceptors such as quinoxaline, BTD, or pyrazinoquionxaline and a thiophene unit introduced as a π-bridge.26 A few examples of polymers of the APFO series are shown in Figure 8.7.

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SCHEME 8.2  Synthesis of borolane-functionalized fluorene.

FIGURE 8.7  Structures of some APFOs, silafluorenes, and carbazoles (P1–P7).

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Most of the APFO series polymers were prepared by the Suzuki coupling reaction due to the easy synthesis of the boronic ester-functionalized fluorene unit. The UV-visible spectroscopy study showed that the absorption edge of P1 APFO-3 and P2 APFO-1527,28 fall in the range of 640–655 nm due to the weaker BTD and quinoxaline acceptor moieties, respectively, used in the construction of the polymers. On the other hand, the absorption onset of the polymers that possess a pyrazinoquinoxaline acceptor (P3) unit showed a significantly red-shifted absorption edge. In fact, P3 APFO-Green 9 has an absorption edge of 900 nm,29 significantly red-shifted compared to that of both P1 APFO-3 and P2 APFO-15. While the HOMO energy levels of the polymers (P1-P3) as determined by square wave voltammetry is deep enough (below 5.7 eV), the LUMO of the polymers was found to be significantly affected by the electron-withdrawing strength of the acceptor groups. For instance, P2 APFO-15 was found to have a LUMO of −3.6 eV while the value for P3 APFO-Green 9 was measured to be −3.9 eV. Clearly, the LUMO energy level of P3 is close to the LUMO of PCBM-based acceptors (~-4.0 eV), which will lower the driving force for the charge transfer to the acceptor as compared to that of P2 APFO-15. The polymers were used for photovoltaic device fabrication by mixing with PCBM[60] (P1 APFO-3 and P2 APFO-15) and PCBM[70] (P3 APFO-Green 9) in a standard device configuration giving rise to 3.5%, 3.7%, and 2.3%, respectively. The lower performance of P3 APFO-Green 9 as compared to the other polymers is limited by two key factors: the charge transport (affecting the fill factor) and less efficient charge formation (affecting the short circuit current). The introduction of silicon and nitrogen bridging atoms at position 9 of the fluorene unit to yield the corresponding silafluorene and carbazole based D–A conjugated polymers, respectively, have been a focus of intense research to further tune the properties of the polymers. Some of the earlier polymers developed are shown in Figure 8.7. P4 PSI-DBT contains silicon as a bridging atom in the fluorene and BTD as an acceptor sandwiched between thiophene rings. This polymer is structurally analogous to APFO-3, except for the use of the silicon heteroatom as a bridge in the fluorene. The functionalized silicon-containing fluorene is prepared according to Scheme 8.3, which was further combined with the thiophene flanked acceptor units such as BTD in P4. The polymer was prepared by the Suzuki coupling reaction. P4 PSI-DBT showed red-shifted optical absorption with the onset of absorption measured to be 681 nm as compared to that of P1 APFO-3 (655 nm). P4 PSI-DBT gave an improved performance30 when blended with PCPBM[60], giving a PCE of 5.4% with a Voc of 0.90 V, a Jsc of 9.5 mA/ cm2, and a FF of 50.7% in single junction solar cells. The higher performance of P4 PSiF-DBT in solar cells is related to its higher hole mobility of 1×10−3 cm2/Vs, which was nearly ten times higher than that of P1 APFO-3 (3×10−4 cm2/Vs). The higher charge mobility of P4 PSIF-DBT is believed to originate from the improved crystallinity of the polymer due to the bigger atomic size of silicon that pushes the side chain away from the backbone as compared to the carbon counterpart. High mobility can ensure effective charge carrier transport to the electrode and reduce photocurrent loss in solar cells.31,32 Another structural modification strategy explored is replacing the bridging carbon atom of fluorene with nitrogen to prepare a carbazole based donor unit for the synthesis of D–A conjugated polymers. Scheme 8.4 shows the synthesis of the functionalized carbazole donor unit. One very popular polymer in this class is P5 PCDTBT33 (Figure 8.7), which possesses the same BTD acceptor unit sandwiched between two thiophenes as in P1 APFO-3 and P4 PSiF-DBT. The absorption onset of P5 PCDTBT was found to be 660 nm at solid state with corresponding optical bandgap of 1.88 eV. The polymer was found to have a slightly red-shifted absorption edge as compared to P1 APFO-3 but blue shifted as compared to P4 PSi-DBT. The molecular energy levels were determined by cyclic voltammetry to be −5.5 and −3.6 eV for the HOMO and LUMO energy levels, respectively. The photovoltaic properties of P5 were studied by blending the polymer with PCBM[60] in a 1:4 ratio in a bulk heterojunction device that gave a PCE of 3.6% with Jsc = 6.92 mA/cm2, a Voc of 0.89 V, and FF of 63%. The average film thickness, measured with atomic force microscopy (AFM), was around 70 nm. The high FF of the devices shows that the charge mobility is well-balanced and that no significant recombination loss occurred within the active material at this thickness.34 Furthermore, the high FF and Jsc

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SCHEME 8.3  Synthesis of bis borolane-functionalized silafluorene.

SCHEME 8.4  Synthesis of bis borolane-functionalized carbazole.

indicate that there are low serial resistances from the materials in this configuration. Interestingly, the devices were fabricated and characterized in air and demonstrated the high stability of P5 PCDTBT. Looking at the current output of scientific papers based on PCDTBT, it is clear there was a huge interest in this material that led to an in-depth study to understand the material property and improve the power conversion efficiency.35 Indeed, the power conversion efficiency of devices based on P5 PCDTBT in single junction structure rose above 7% with continuous optimization of the devices.35,36 Recently, Thompson’s group developed the synthesis of P5 PCDTBT with a direct arylation method,37 which is an emerging metal-catalyzed reaction. Different reaction conditions were explored such as the use of neodecanoic acid (NDA) and the use of a more affordable K 2CO3 base. The prepared polymer’s yield reached up to 70%. Photovoltaic devices that reach to a PCE of 2.08% were obtained together with PCBM[60] from the batch of polymers prepared by the direct arylation method, comparable to the performance achieved from a model polymer prepared by the Suzuki coupling reaction. The carbazole donor unit was combined with other acceptor units to produce materials with improved properties. Recently, a carbazole unit was combined with a quinoxaline unit (fluorinated and non-fluorinated) with a thiophene π-bridge to yield a D–A polymer,38 shown in Figure 8.7 (P6 and P7). While the fluorinated quinoxaline-containing polymer (P7) showed higher molecular weight (54.4

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kDa), the non-fluorinated quinoxaline-containing polymer (P6) showed significantly lower molecular weight (9.93 kDa) for reasons that are still unclear. The absorption onset of the polymers was found to be ~607 and 626 nm for P7 and P6, respectively. In the electrochemical study to determine the molecular energy levels, the HOMO of P7 fluorinated quinoxaline was lower by 0.09 eV as compared to the P6 non-fluorinated quinoxaline, showing the effect of fluorine in lowering the HOMO. The LUMO energy level of the two polymers differ only by 0.02 eV, showing the minor effect of the fluorination in adjusting the LUMO energy levels. Optimized photovoltaic devices produced a range of 4.7 to 5.19% PCE using chlorobenzene as a processing solvent. Cyclopentadithiophene (CPDT) and corresponding heteroatom-bridged aromatic units have been used as a donor unit in building D–A polymers together with acceptors such as BTD and thienopyrroledione (TPD). The synthesis of a CPDT molecular unit39 is shown in Scheme 8.5. The bridging atoms are varied from carbon to silicon to germanium to nitrogen, and even imine (C=N) functionalities were introduced to further tune CPDT. The silicon version of CPDT was synthesized as shown in Scheme 8.6. The germanium-replaced CPDT can be prepared by using dichloro-dialkyl germanium (R 2GeCl2) which can be prepared from the reaction between two equivalents of a Grignard reagent of the alkyl group and germanium tetrachloride. On the other hand, the synthesis of the dithienopyrole donor unit40 is accomplished as shown in Scheme 8.7. The bridge head imines are prepared from the reaction between the amines and CPDT-one and subsequently functionalized into distannyl compounds to make them ready for Stille coupling reactions. Some of the structures discussed in this group are shown in Figure 8.8 (P8–P19). CPDT (P8) (PCPDTBT)41 or silicon-bridged CPDT (P9) (PSBTBT) donor units combined with a BTD acceptor42 are the first two polymers to be studied thoroughly. Both polymers revealed excellent absorption coverage that extends up to 900 nm with a corresponding optical bandgap of 1.40 eV. The HOMO energy level of the polymers was estimated to be in the range of −5.3 to 5.05 eV. Obviously, these

SCHEME 8.5  Synthesis of CPDT.

SCHEME 8.6  Synthesis of silicon functionalized CPDT.

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SCHEME 8.7  Synthesis of distannyl functionalized dithienopyrrole.

FIGURE 8.8  CPDT and derivatives donor units-containing polymers (P8–P19).

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materials have a raised HOMO as compared to the previously discussed APFO polymers (for example P1 APFO-3), which is due to the stronger electron-donating effect of the CPDT unit. The LUMO of the polymers were measured to be −3.57 and −3.27 eV for P8 PCPDTBT and P9 PSBTBT, respectively. P8 PCPDTBT showed a PCE of 3.2% in combination with PCBM[70] in photovoltaic devices. On the other hand, the photovoltaic performance of the silicon-bridged polymer (P9) reached 5.1%. Interestingly, in this case, the silicon-bridged polymer (P9 PSBTBT (3 × 10−3 cm2/Vs)) also showed a hole mobility three times higher than that of P8 PCPDTBT. When studied with X-ray diffraction, the replacement of carbon by silicon improved the crystallinity of the polymer, which led to a better charge transport when it is blended with PCBM in devices.43,44 This result is similar to what is observed in APFO series polymers where the silicon-bridged systems showed higher charge carrier mobility. The silicon- (P10) and germanium- (P11) substituted CPDT-based donor units were also combined with a TPD acceptor to give two polymers (Psi and PGe).45 TPD was the choice of interest as an acceptor due to its inherently moderate electron acceptor nature and a unit that tends to give high open circuit voltage when incorporated into polymers. The two polymers were prepared by metal-catalyzed coupling reactions and gave high molecular weight polymers. The germanium-bridged polymer (P11) showed red-shifted absorption peak and absorption edge when compared to its silicon counterpart (P10), and these had optical bandgaps of 1.69 and 1.73 eV, respectively. The use of bigger atoms like germanium can even improve crystallinity and leads to better packing in solid state of the polymer chains by pushing the solubilizing side chains out from the backbone of the polymer. A theoretical study on optimized geometry with a methyl group attached as a substituent showed that the methyl groups are farther away from the first carbon of the thiophene ring in the germanium-bridged unit (3.27 Å vs. 3.11 Å), supporting the idea that the bigger atoms are preferred for a better π-π stacking of the polymer chains. Bulk heterojunction solar cells were fabricated using a blend of polymers with PCBM[70] as active layers in inverted device architectures ITO/ZnO/polymer:PCBM[70]/MoO3/Ag. P10 gave an average Jsc of 11.5 mA/cm2, a Voc of 0.89V, and FF of 65%, resulting in an average PCE of 6.6%. The DTG-containing polymer P11 PGe gave a higher Jsc of 12.6 mA/cm2 and FF (68%), a Voc of 0.85 V, and an average PCE of 7.3%. Further optimization of the devices increased the PCE to 8% for the germanium-containing polymer. The nitrogen-bridged CPDT (dithienopyrrole) as a donor unit with different alkyl side chain were combined with 2,1,3-benzothidiazole46 to prepare novel, near IR-absorbing, low bandgap D–A polymers and have been used for the fabrication of solar cells (P12–P14). The side chains attached to the dithienopyrrole were varied from 1-octylnonyl (P12) to 1-hexylheptyl (P13) and 1-pentylhexyl (P14) to study the effect of the length of side chains in the overall properties of the polymers. Generally, the polymers show strong absorption in the wavelength range of 600–900 nm with the absorption coefficient being enhanced as the length of alkyl chain decreases. The polymers showed similar HOMO energy levels, both falling in the range of −4.81 to −4.89 eV while the LUMO energy level of the polymers was −3.08 eV. The materials were used to fabricate photovoltaic devices and, in all cases, the Voc ranged from 0.4 to 0.55 V, which is low compared to the case of other heavy element bridging groups. The low Voc could arise from the raised HOMO energy level of the polymer. The best efficiency was recorded from the polymer with shorter side chain (P14), with PCE reaching 2.8% when blended with PCBM [60] in a 1:3 blend ratio. In 2013, the Bazan research group reported a series of polymers with a CPDT donor molecule functionalized with an imine group (C=N) at the bridging position and coupled with an alkoxy substituted benzodithiadizaole (P15–P19) (Figure 8.8).47 Figure 8.9 shows the resonance contributors of the imine and carbonyl-functionalized CPDT unit and polymers prepared from such units and the effect on bandgap and other properties of the polymers. The first structure is a CPDT donor unit functionalized with a keto group. The use of such donor units gives a bandgap of about 1.2 eV due to the contribution of the primary resonance contributor. However, the bandgap is lowered to 0.8 eV with structures when the oxygen is replaced with a carbon that is attached to more electron-withdrawing cyano groups. Another well-known system are copolymers of CPDT and BTD that give low bandgap polymers with localized HOMO and LUMO. The imine functionality is found to stabilize both the HOMO and LUMO and thus finely tuning the electronic structure of the polymers. The imine-functionalized polymers (P15–P19)

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FIGURE 8.9  Molecular structures of CPDT-based donor units.

showed optical bandgaps in the range of 1.40 to 1.46 eV. The molecular energy levels were studied with CV and the use of electron-withdrawing groups on the phenyl ring lowers both the HOMO and LUMO. For instance, para hexyl substituted (P15) gives a HOMO of −5.38 eV and LUMO of −3.73 eV. With the replacement of the hexyl side chain with fluorine, (P16) lowers both the HOMO and LUMO to −5.44 eV and −3.99 eV, respectively. P17 P3 showed a HOMO of −5.46 eV and LUMO of −4.00 eV. The HOMO and LUMO positions were found to be sensitive to the nature of the side chains and substituents as it is expected. In terms of molecular energy level, the HOMO and LUMO of the imine-functionalized polymers showed deeper energy level as compared to PCPDTBT. The photovoltaic properties of P15 P1 and P17 P3 were tested by blending the polymers with PCBM[70] processed from chlorobenzene in the presence of 3% DIO. P15 P1-based devices showed a Jsc of 11.2 mA/cm2, Voc of 0.56 V, FF of 49%, and overall performance of 2.88%. On the other hand, P17 P3 gave a better PCE, reaching 3.02%, with slightly lower Jsc (9.8 mA/cm2) and FF = 44% but higher Voc of 0.70 V. The higher Voc corresponds well with the deeper HOMO of P17 P3 as determined by CV. Dithienopyrrole, a nitrogen atom bridged of CPDT, is an electron-rich and planar system used for the synthesis of D–A conjugated polymers (P12–P14). Even though this unit gave an excellent charge carrier mobility in field effect transistors, its application in solar cell was limited due to the lower Voc produced owing to its electron-rich property. In order to circumvent this problem, Maes and coworkers developed a strategy to control the electron richness of the TPD unit by attaching a carbonyl group as a side chain (P20–P22).48 The synthesis of the carbonyl-substituted dithienopyrrole is shown in Scheme 8.8. The unit was copolymerized with a 0F, 1F, and 2F substituted quinoxaline acceptor to make three D–A polymers (P20–P22) (Figure 8.10). The polymers were prepared by Stille coupling reactions and obtained in good solubility in organic solvents. In the optical study, the non-fluorinated polymer (P20) showed a red-shifted absorption with an absorption edge of up to 703 nm at solid state. A clear blue shift was observed with mono- (P21) and difluorinated (P22) polymers. The HOMO and LUMO energy levels for the P20 were calculated to be −4.64 and −2.61 eV, respectively. The HOMO of the fluorinated polymers was found to be even lower, which is as expected due to the electron-withdrawing property of fluorine atoms attached to the quinoxaline group. A standard device structure, glass /ITO/PEDOT:PSS/active layer/Ca/Al, was used to investigate the photovoltaic properties of the polymers. The highest performance was obtained from P20, which had a 4.81% PCE with a Voc of 0.67 V, Jsc of 12.57 mA/cm2, and an FF of 54%. Even though the fluorinated polymers gave an even higher Voc, the performances were limited by the lower Jsc, giving a PCE of 2.78% and 2.5% for the 1F and 2F polymers, respectively.

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SCHEME 8.8  Synthesis of carbonyl side group substituted dithienopyrrole.

FIGURE 8.10  Dithenopyrrole-quinoxaline-based copolymers (P20–P22).

8.2.2 Thiophene and Derivatives as a Donor Unit in Donor-Acceptor Polymers Thiophene is one of the most widely used donor moieties to construct D–A polymers. In these sections, selected thiophene-based D–A conjugated polymers will be discussed. 8.2.2.1 Thiophene​/Thie​nothi​ophen​e/Sel​enoph​ene-Q​uinox​aline​ Thiophene, thienothiophene, and selenophene are among the donor units that has been combined with quinoxaline by the Stille coupling reaction. Some of the most common polymers are shown in Figure 8.11. P23 TQ1 is a blue polymer in which the synthesis and optical properties were reported by Yamamato et al.49 for the first time. Similar structures were developed by other groups and used for solar cell applications.50 Attracted by the ease of synthesis of the polymer and by the promising optical properties,

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FIGURE 8.11  Structures of thiophene-based polymers (P23–P28).

Andersson’s group synthesized TQ1 in 2010 and used the material for photovoltaics.51 P23 showed a broad absorption band characterized by strong π-π stacking as revealed from the stronger absorption band at longer wavelength in the solid of the UV-visible test. The band weakened when it is in solution and when it is run at higher temperature, which shows the higher aggregation tendency of the polymer in solid state. The polymer showed ideal HOMO (−5.7 eV) and LUMO (−3.3 eV) energy levels for photovoltaics. The deeper HOMO is needed to get high Voc while the well-situated LUMO energy level will provide enough driving force for charge separation which is essential for achieving higher Jsc. The photovoltaic performance of the hero cell reached 6% PCE when blended with PCBM[70] in 1:3 ratio and processed with ODCB with corresponding Jsc of 10.5 mA/cm2, Voc = 0.89 V, and FF = 64%. Yang’s group studied the same polymer (P23 TQ1) and PC71BM blend to fabricate photovoltaic devices with the addition of solvents such as 1,8-octanedithiol (ODT), 1,8-diiodooctane (DIO), diphenylether (DPE), and 1-chloronaphthalene (CN)) to ODCB processing solvent with the intention of optimizing the nanomorphology. The best device achieved a 7.08% PCE with a Voc of 0.91 V, Jsc of 12.2 mA/cm2, and FF of 64% with the use of 5% CN solvent additive to ODCB to process a 1:2 blend ratio of P23 TQ1 and PCBM[70]. The Jsc and Voc were improved in this case, which was related to the well-developed nanomorphology. The highperforming devices gave a balanced hole and electron mobility, which improved Jsc and FF.52 The use of thieno[3,2-b]thiophene as a donor unit together with a quinoxaline acceptor resulted in a polymer with a higher tendency of aggregation as characterized by UV-vis spectroscopy.53 As compared to P23 TQ1, the use of the thienothiophene donor unit (P24) pushed the absorption edge to the red side by 30 nm due to the enhanced conjugation length. The application of this polymer as donor, coupled with PCBM[70] in photovoltaic devices with standard geometry, allowed it to obtain solar cells with a photovoltaic performance of about 5%. The use of selenophene as donor unit to quinoxaline was also studied (P25).54 The replacement of thiophene by selenophene pushed the absorption onset to longer wavelength and also increased the strength of the absorption band in the longer wavelength of the absorption profile indicating the aggregation/packing tendency of the polymer, which could be promoted by the intermolecular Se-Se interactions.55 Photovoltaic devices were prepared from this polymer using the configuration of glass/ITO/PEDOT:PSS/active layer/LiF/Al) by spin coating from ODCB solutions. In a preliminary study, the highest efficiency obtained from this polymer was 1.7% PCE. Both the Jsc and FF were found to be lower as compared to P23 TQ1-based devices, probably due to the unoptimized nanomorphology. Another structural modification to the thiophene/selenophene-quinoxaline based polymers is the

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attachment of one and two fluorine atoms to quinoxaline (P26–P28) in order to tune the properties of the polymers. Some of the fluorinated structures are shown in Figure 8.11.56 With the addition of fluorine on the quinoxaline, the optical absorption edge generally shifted to shorter wavelength. However, the HOMO and LUMO energy levels were lowered as determined by CV as compared to the non-fluorinated counterparts. The difluorinated quinoxaline-containing polymer with thiophene as a donor (P26), for instance, was found to have HOMO and LUMO energy levels of −6.04 and −3.76 eV, respectively. The polymer was used as a photoactive material in OPV to give a PCE of 1.6%.57 The PCE increased to 4.41%56 from the same material with careful optimization of the devices. The use of mono-fluorinated quinoxaline unit was coupled with thiophene to give thiophene-quinoxaline-based polymers (P27).58 The photovoltaic performance was tested in standard BHJ devices by blending the polymer with PCBM[70]. A first optimization step resulted in a device with a PCE of 4.01%, FF of 43.3%, Jsc of 9.75 mA/cm2, and a Voc of 0.95 V. An additional thermal treatment at 110 °C for one minute led to an improved PCE of 5.3%, with Jsc of 10.1 mA/cm2, FF of 60%, and Voc of 0.87 V. 8.2.2.2 Thiophene-Isoindigo Donor-Acceptor Polymers The isoindigo unit is a naturally occurring pigment that is an isomer of the famous indigo dye. Structurally, the isoindigo is a symmetrical molecule consisting of two electron-deficient lactam rings fused together. Reynolds et al. reported the first synthesis of p-type isoindigo-based small molecules in 2010, with interesting optical and electrochemical properties suitable for solar cell application.59 The Andersson group reported an early D–A conjugated polymer possessing thiophene and isoindigo units called P29 PTI-160 synthesized by Stille coupling reaction. Later, an oligothiophene donor unit was combined with an isoindigo unit to get P30 P3TI61 (Figure 8.12). Both polymers were prepared in high molecular weight and found to have good solubility in organic solvents. In the UV-visible spectra, a strong ICT band is observed, which implies the strong electron-withdrawing characteristics of the isoindigo unit. In addition, P30 P3TI showed stronger absorption in the shorter wavelength region as compared to PTI-1 due to the absorption contribution of the three thiophene units in the former in the high energy region. The HOMO and LUMO energy levels were calculated from a square wave voltammogram to be −5.85 and −3.88 eV for P29 PTI-1 while the corresponding values for P30 P3TI were measured to be −5.82 and −3.83 eV, respectively. Interestingly, the polymers showed similar energy levels with a slightly raised value for P30 P3TI. Photovoltaic devices with a structure of ITO/PEDOT:PSS/ polymer:PC71BM/LiF/Al were fabricated using P29 PTI-1 as a donor polymer to give a PCE of 3.0% and high Voc of 0.89 V with a Jsc of 5.4 mA/cm2 and FF of 63%. Alternatively, using the same device configuration processed from ODCB with the addition of DIO, P30 P3TI gave a 6.3% PCE with a Voc of 0.70 V, Jsc of 13.1 mA/cm2, and FF of 69%. The higher performance of P30 P3TI comes from the improved Jsc and FF, which might have arisen due to the morphology change. Further increasing the oligothiophene length to six thiophene rings yielded a material with improved absorption and crystallinity and hence charge carrier mobility in photovoltaic devices.62 A blend of P31 P6TI and PCBM[70] gave a PCE of 7.25% with Jsc of 16.24 mA/cm2, which is among the highest performing isoindigo-based low band gap polymers. Five isoindigo-terthiophene based polymers (P32–P35), including P30 with different length and topography of side chains attached on the isoindigo, were prepared by the Stille coupling reaction63 (Figure 8.12). The linear side chains improve nanostructured order compared to the branched 2-ethylhexyl side chains, as observed by grazing-incidence wide-angle X-ray scattering. The C6 alkyl side chain-containing polymer (P32) and C8 (P30) gave a higher PCE of 5.2%.

8.2.3 Benzodithiophene as a Donor Unit for the Synthesis of Donor-Acceptor Polymers Benzodithiophene is an aromatic molecule built from two thiophene rings fused to a central benzene ring with an interesting and peculiar characteristic. The BDT unit is known for a rigid and planar structure, allowing an extended π-conjugation and favorable inter-chain π-π stacking. In addition, the

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FIGURE 8.12  Thiophene-isoindigo-based polymers (P32–P35).

centrosymmetric BDT unit (most commonly used) can be easily synthesized and functionalized with different side chains,64,65 such as an alkoxy group, alkyl thiophene, benzene, and so on. The synthesis of the centrosymmetric BDT core unit is approached according to the synthetic steps shown in Scheme 8.9.66,67 Encouraged by the unique nature of BDT, a large number of D–A polymers were developed since its first application as a donor moiety.66,68 Some of the acceptors that have been coupled with BDT include quinoxaline, TPD,69 DPP,70,71 and thienothiophene.72,73 The following few sections will focus on the discussion of D–A polymers developed more recently. 8.2.3.1 Benzodithiophene-Thienothiophene-Based Donor-Acceptor Polymers The introduction of the thieno[3,4-b]thiophene unit into the backbone of the D–A polymer is known to enhance the quinoidal character of the backbone of the polymer and hence leads to a narrower bandgap. As shown in Figure 8.13, the formation of the quinoid form happens readily as the aromatic character of the upper thiophene is increased during the transformation from the aromatic to the quinoid form.

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SCHEME 8.9  Synthesis of the centrosymmetric BDT and derivatives.

FIGURE 8.13  Aromatic and quinoid isomers of thieno[3,4-b]thiophene.

From the structure of D–A polymers developed in the past, the PTB family that is built from various derivatives of BDT and thieno[3,4-b]thiophene are worth mentioning. The first type of such polymers is known in the literature as P36 PTB1, which was reported by Yu’s group in 2009 (Figure 8.14).74 In this polymer, an alkoxy side chain was attached to the BDT core while an alkyl group is appended to the thienothiophene ester group. The polymer was found to have a longer onset of absorption that extends up to 774 nm, unlike the benchmark high band gap P3HT polymer, which has an absorption onset of ~650 nm. The material was used to construct a bulk heterojunction photovoltaic device together with PCBM[70]. A 1:1.2 blend of polymer and acceptor with a thickness of 100–110 nm gave a Jsc of 15.0 mA/ cm2 with a Voc of 0.56 V and FF of 63.3% and yielding an overall PCE of 5.30%. The same group developed a series of polymers to tune the structure of the polymer by introducing different side chains on the BDT and on the ester group of the thienothiophene unit yielding a series of polymers.75 Among the new polymers, a fluorine atom was introduced to the thienothiophene unit to give a finely tuned structure P37 (PTB-4), as shown in Figure 8.14. With the introduction of a fluorine atom attached to the thienothiophene ring, the HOMO of the polymer was lowered (−5.12 eV) as compared to the polymers without the fluorine atom (for instance P36 PTB-1, HOMO = −4.9 eV). The lowered HOMO in P37 PTB-4 is good for achieving a higher Voc. The hole mobility of the polymers was measured according to method based on the space charge limited current (SCLC) model. The hole mobility of P36 PTB1 and P37 PTB4 were measured to be 4.7 × 10−4 and 7.7 × 10−4 cm2/Vs, respectively. The introduction of the fluorine atom into the polymer is one reason for the observed enhanced hole mobility in P37 PTB-4, which might arise from the favorable inter-polymer π-stacking due to the interaction between the electron-rich and electron-deficient aromatic units in the polymer structure. Photovoltaic properties of the polymers were investigated in solar cell structures of ITO/PEDOT:PSS/polymer:PC61BM(1:1 wt. ratio)/ Ca/Al which gave 6% PCE (Voc = 0.74 V, Jsc = 13.0 mA/cm2, FF= 61.4%) processed from an o-DCB and DIO solvent mixture. A slight modification of the side chain of P37 PTB-4, i.e. replacing the side chains

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FIGURE 8.14  Structure of benzodithiophene-thienothiophene polymers (P36–P41).

with 2-ethylhexyl side chain on both BDT and thienothiophene, yielded a polymer known as P38 PTB-7, with good solubility in organic solvents, a high number average m ­ olecular weight (97.5 kDa), and PDI of 2.1. The hole mobility of P38 PTB-7 was found to be 5.8 × 10−4 cm2/Vs, which is quite high considering the negative effect the branched side chains might have on charge mobility. The device prepared from a mixture of P38 PTB-7 and PCBM[70], processed with a mixture of o-DCB and 3% DIO in the solvent, gave a PCE of 7.4%76 with corresponding Jsc of 14.5 mA/cm2, FF of 69%, and Voc of 0.74V. In 2012, Cao’s group used P38 PTB-7 and blended it with PCBM[70] in an inverted structure using a water soluble PFN polymer and MoO3 as an electron and hole transporting layer, respectively, to give over 9% PCE.77 The increase in performance is mainly due to the higher current harvested from the device (>17 mA/cm2) which might have arisen from the ohmic contact created for an effective photogenerated charge carrier collection and optimum photon harvest in the device. Another polymer in the same class is PTB7-Th (P39) (Figure 8.14), in which a thiophene aromatic rings are attached to the BDT core unit instead of the alkoxy group. P39 PTB7-Th showed an even narrower bandgap (1.59 eV, with an absorption onset of ∼780 nm), which is closer to the optimum bandgap (∼1.1–1.5 eV) for a single junction cell.77 Furthermore, P39 PTB7-Th has a HOMO of –5.22 eV (∼0.07 eV deeper than that of P38 PTB7), which is also very close to the ideal energy levels. It is therefore expected that the P39 PTB7-Th devices can simultaneously deliver higher Voc and Jsc than that of P38 PTB7-based devices. Indeed, with the continuous modification of the interfacial layer and use of PFN as a cathode interface layer, a blend of P39 PTB7-Th and PCBM [70] in a 1:1.5 blend ratio gave over 10%78 PCE with a Voc of 0.825 V, Jsc of 17.43 mA/cm2, and FF of 73.78%. In addition, Chen et al. used P39 PTB7-Th and PCBM[70] with ITO/ZnO-C60 as a cathode and MoO3/Ag as an anode.79 The power conversion efficiency increased with the Zno-C60 interfacial layer relative to that of ZnO, with the PCE being 9.35% and 7.64%, respectively. All three of the corresponding parameters (Voc, Jsc, and FF) were improved with the modified interface layer. It was reported that the electron mobility of ZnO-C60 is significantly higher than that of ZnO by a factor of 50. The use of this cathode material improved the electron transport and reduced the chance of electron/hole recombination at the interface, which is a plausible reason for the PCE improvement. Li and coworkers optimized the devices based on P39 PTB7-Th and PCBM[70], using mixed solvent additives of DIO and NMP to ODCB. An addition of 1.5% DIO and 1.5% NMP gave a high PCE of 10.8%80

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with corresponding Voc = 0.82 V, Jsc = 19.1 mA/cm2, and FF = 69.1% in simple, conventional devices. The films processed from binary solvent additives gave the highest and most balanced hole and electron mobility of 7.61x10−3 cm2/Vs and 6.52 × 10−3 cm2/Vs, respectively, which is consistent with the high FF and better photovoltaic performance of the devices. The role of alkyl chains has been explored in different D–A polymers. Alkyl chains play an important role in controlling solubility, packing, crystallinity, and device performances. The selection of the side chain (whether it is straight or branched or short or long or alkyl chains with a heteroatom) has to be done carefully as it affects the quality of the materials. One of the side chains that has gained focus recently is the alkylthio substitution as a side group. Li and coworkers reported a P39 PTB7-Th version with an alkylthio attached to the BDT unit81 instead of the 2-ethylhexyl side chain to prepare P40 PBDTT-S-TT (Figure 8.14). As compared to P39 PTB7-Th, the sulphur-containing D–A polymer (P40) showed a red-shifted onset of absorption and the HOMO energy level was pushed down by 0.11 eV. Polymer solar cell with a blend of P40 PBDTT-S-TT and PCBM[70] gave a Voc of 0.84 V, leading to a higher PCE of 8.42% without any processing additive or post treatment. Meanwhile the batch of P39 PTB7-Th the group worked on gave a Voc = 0.77 V and PCE of 7.42% with the addition of DIO. In a similar approach, Hou and coworkers replaced the branched side chain of the BDT unit in P40 PBDTT-S-TT with a linear octylthio side chain to take advantage of side chain engineering in tuning properties of the polymer and prepared P41 PBDTT-TS1.82 Optical study showed that the absorption onset extends to 820 nm corresponding to an optical bandgap of 1.51 eV, which is lower than P39 PTB7-Th (1.58 eV). The HOMO and LUMO levels of P41 PBDT-TS1 are respectively determined to be −5.33 eV and −3.52 eV by electrochemical cyclic voltammetry. The polymer was used to fabricate solar cells with a configuration of ITO/PEDOT:PSS/Polymer:PCBM[70]/Mg/Al, processed from ODCB with the addition of 3% DIO to give a PCE of 9.19%. An XRD study showed that the film prepared from this polymer has a pronounced feature of π-π stacking, which is a behavior not commonly seen in other types of polymers. Due to the apparent high degree of ordering of P41 PBDT-TS1, the hole mobility (1 × 10−2 cm2/Vs) was found one order of magnitude higher than P39 PTB7-Th (2.83x10−3 cm2/Vs) which positively contributed to the enhancement of Jsc and FF. Two BDT and thienothiophene based polymers were prepared and characterized (P42 and P43)83 (Figure 8.15). P42 was built from an alkoxy substituted BDT donor unit and a sulfonyl substituted thienothiophene group. On the other hand, a thiophene-substituted donor unit was used for the synthesis of P43 with the same acceptor used as the earlier polymer. Moreover, thiophene π-bridge was introduced between the donor and the acceptor. Stille coupling reaction between the respective monomers gave the desired polymers. The optical study of the polymer films shows an absorption edge of 750 and 780 nm for P42 and P43, respectively. In the CV study, the thiophene-bridged polymer (P43) showed a slightly raised HOMO energy level. The hole mobility shows that P42 (4.56 × 10-4 cm2/Vs) evinced a lower mobility as compared to P43 (2.76 × 10-3 cm2/Vs), indicating that elongating the backbone by introducing thiophene is a good method for improved charge carrier mobility. Photovoltaic devices were prepared from these polymers by blending them with PCBM[70] processed from ODCB with 3% DIO. While P42 gave a maximum PCE of 6.36%, the thiophene-bridged polymer (P43) gave an improved PCE of 7.81%. The simple modification in the monomer synthesis resulted a 25% increase in the PCE. The synthesis and characterization of a series of polymers prepared from thienothiophene and BDT was similar with P39 PTB7-Th, but this time thiophene was inserted between the BDT and thienothiophene unit as a π-bridge (P44–P47) (Figure 8.15). Also, a varying number of fluorine atoms (0F, 1F, 2F, 3F) were introduced to investigate the effect of molecular energy modulation by introducing various number of fluorine atoms in the BDT-TT based polymers in a systematic way72 (P44–P47). The study revealed that when increasing the number of fluorine atom substitution, the HOMO energy level was lowered. Thus, the HOMO was measured to be −4.90, −4.95, −5.15, and −5.20 eV for PBT-0F (P44), 1F (P45), 2F (P46), and 3F (P47) polymers, respectively. Photovoltaic devices were fabricated according to the configuration ITO)/ PEDOT:PSS/polymer:PCBM[70]/Ca/Al and the polymer and acceptor blend was spin coated from o-DCB and 3% DIO. The fluorinated polymers gave a diminished domain size

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FIGURE 8.15  Structure of polymers (P42–P47).

and interpenetrating network with evident polymer fibrils, which is a morphology needed for efficient charge separation and transport,84 so the high Jsc and FF were obtained from the fluorinated polymers. Indeed, a higher photovoltaic performance was achieved from P47 PBT-3F, reaching 8.6% with corresponding Voc of 0.78 V, Jsc of 15.2 mA/cm2, and FF of 72.4%. The effect of fluorine was systematically studied85 by Cao’s group in 2014 (Figure 8.16). The authors synthesized random terpolymers consisting BDT, theinothiophene and a different loading of fluorinated thienothiophene to yield polymer with 0% (P48), 25% (P49), 50% (P50), 75% (P51) and 100% (P39) of the fluorinated thienothiophene. The structures of the polymers are shown in Figure 8.16. The polymers showed comparable absorption onset in both solution and solid state and showed comparable molecular energy levels with HOMO that range between −5.22 and −5.28 eV and LUMO in the range of −3.65 and −3.68 eV. The photovoltaic properties of the polymers were studied, and the performance increased with a higher loading of the fluorinated thienothiophene unit. The best performance was 8.36% PCE with corresponding Jsc of 16.74 mA/cm2, Voc of 0.79 V, and FF of 63% for the 100%F (P39) polymer. The performance with 75%F (P51), 50%F (P50), 25%F (P49), and 0%F (P48) was measured to be 7.28%, 7.01%, 6.38%, and 6.69%, respectively. One reason for the high performance of the 100% fluorinated polymer (P39) is due to the higher hole mobility (2.75 × 10−5 cm2/Vs) as compared to the other polymers, though the crystallinity of the 50%F (P50) and 75%F (P51) polymers were found to be higher than that of 100%F. This example shows that even though crystallinity of materials is a needed property, it does not guarantee high performance when used in solar cells. A crystalline material with poor connectivity between the crystallites, for instance, can behave badly in solar cell devices.

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FIGURE 8.16 Benzodithiophene-thienothiophene terpolymers with different loading of fluorinated thienothiophene.

FIGURE 8.17  BDT-TPD-based polymers (P52–P54).

8.2.3.2 Benzodithiophene-TPD-Based Donor-Acceptor Polymers Thieno[3,4-c]pyrrole-4,6(5H)-dione (TPD) has proved to be a versatile aromatic unit to construct a range of high bandgap polymers in the past. The imide structure of TPD provides the resulting polymers with simplicity of synthesis, excellent stability, and structural flexibility in tuning the solubility and packing property of the polymers.86 One of the donor units that was used to construct successful D–A polymers in the past was BDT and herein some of the polymers are discussed. Leclerc’s group reported P52 BDT-TPD, with structure shown in Figure 8.17,87 for the solar cell application (PBDTTPD) P52 for the first time. The polymer was prepared by the Stille cross-coupling reaction between a stannylated BDT and a dibromo TPD monomers catalyzed with a palladium catalyst. The optical properties of the polymer were studied, giving an optical bandgap of 1.8 eV (absorption onset of 685 nm). The HOMO and LUMO energy levels of the polymer were found to be −5.56 and −3.75 eV, as determined from the onset of oxidation and reduction of the cyclic voltammetry, respectively. Photovoltaic properties of the polymer was investigated in a device structure of ITO/PEDOT:PSS/ polymer:PCBM[70]/LiF/Al with an active device area of 1.0 cm2, which gave 5.5% PCE with corresponding Jsc of 9.81 mA/cm2, a Voc of 0.85 V, and a FF of 66% processed in air. The Fréchet group reported three polymers (P52–P54) (Figure 8.17) with different side chains attached on the TPD unit, keeping the donor unit to be the same, in order to study the correlation of side chain modulation to the property of the polymers.88 Note that P52 is the same structure as previously reported by Leclerc’s group. As shown in Figure 8.17, 2-ethylhexyl, 3,7-dimethyloctyl, and octyl side chains were attached onto the TPD unit. By replacing the shorter and bulkier 2-ethylhexyl side chain in P54 P1 with

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FIGURE 8.18  Structure of BDT-TPD polymers (P55 and P56) and non-fullerene acceptors (IDIC and ITIC).

the longer and less bulky 3,7-dimethyloctyl (P53 P2) and octyl side chain (in P52 P3), broader and redshifted absorption spectra with more defined vibronic structure was obtained, which is indicative of the planarization of the conjugated backbone and more efficient packing of the polymer. CV was carried out to determine the electrochemical HOMO levels, giving values ranging from −5.57 to −5.4 eV. The photovoltaic properties were studied in a device structure of ITO/PEDOT:PSS/PCBM[60]/Ca/Al processed from CB and DIO in order to optimize the morphology and resulted a PCE of 4.0%, 5.7%, and 6.8% for P54 P1, P53 P2, and P52 P3, respectively. The higher performance of P52 P3 based device is due to the much-ordered microstructure of the polymer in solid state due to the less dense octyl side chain as seen from the XRD. Later, further optimization efforts on P52 P3-based devices enabled the achievement of an 8.5% PCE.69 Two BDT-TPD-based polymers with fused thienothiophene introduced as a π-bridge were prepared and used as a donor unit in combination with non-fullerene acceptors shown in Figure 8.18.86 The two polymers differ by the presence of the methoxy group as a side group to the P56 polymer while the alkyl groups and the backbone is the same as in P55 (Figure 8.18). Stille cross-coupling reaction was utilized to make P56 PMOT16 and P55 PBDTT-6ttTPD. The polymers exhibit good solubility in chloroform, CB, and o-DCB at room temperature. The absorption onset of the polymers at solid state were found to be similar falling in the range between 667–669 nm with corresponding optical bandgap of around 1.85 eV. P56 PMOT16 (the methoxy group containing polymer) has a stronger vibronic peak at 608 nm as compared to the other polymer, suggesting the role of the methoxy group to encourage the polymer chains ordering at solid state.89,90 Also, a blend of the polymers with the non-fullerene acceptors pushes the absorption edges to 800 nm, which is an important contribution from the acceptors for harvesting more photons. The two polymers were used in solar cell fabrication in combination with the non-fullerene acceptors. P55 PBDTT-6ttTPD-based devices blended with ITIC in a 1:1 ratio and processed from chloroform and 0.5% DIO solvent system in a device ITO/ ZnO/active layer/MoO3/Ag gave an efficiency of around 9.8%, with Voc of 0.9 V, Jsc = 15.6 mA/cm2 and

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FIGURE 8.19  Structure of BDT-TPD polymer prepared by the continuous-flow method (P57).

FF = 67.8% for the best device. Interestingly, a similar device with PCBM[70] as an acceptor only gave a PCE of 6.59% with Voc of 0.78 V, Jsc = 12.2 mA/cm2, and FF= 69.2%. It is important to note here the contribution of the ITIC acceptors toward the improved Jsc. In the case of P56 PMOT16, when blended with ITIC in devices, it gave lower performance due to the reduction of Jsc and FF, though a high Voc of over 1 V was achieved. The overall performance for the best device for the blend of P56 PMOT16 and ITIC was 7.63%. A blend of P56 PMOT16 and PCBM[70] gave a comparable performance of 7.74%. The other non-fullerene acceptor used was IDIC (Figure 8.18), with lower energy levels than ITIC. It is interesting that a blend of P56 PMOT16 and IDIC PSCs showed PCEs of around 10% and low energy loss of 0.68 eV, which surpasses that of P55 PBDTT-6ttTPD:IDIC due to the increase in the Voc of the device. In 2015, the group of Maes reported a polymer P57 BDT-TPD (Figure 8.19) synthesized by continuous-flow method,91 which is a technique known to be effective for real life OPV production. The polymer prepared by the continuous-flow method gave one of the highest performances in polymer solar cells. A conventional device structure of glass/ITO/PEDOT:PSS/P57 PBDTTPD:PC71BM/Ca/Al gave a PCE of 7.9%. Replacing the calcium with a PT-based polyelectrolyte improved the photovoltaic performance to 9.1% with a clear improvement of Jsc to 12.5 mA/cm2. 8.2.3.3 BDT-Quinoxaline-Based Donor-Acceptor Polymers Quinoxaline has been a very popular acceptor for constructing D–A polymers. Some of the highperforming materials for solar cell application developed in the past are quinoxaline-based polymers. Herein, some of the BDT-quinoxaline-based polymers are discussed. For a better understanding of the progresses of quinoxaline-based polymers, interested readers are invited to look at recent review papers.64 2D-conjugation is a technique to tune properties of polymers through improving light harvesting capability by giving a broad and strong absorption band and improving solid state ordering of the polymer chains. In 2012, Hou’s group investigated the effect of 2D-conjugation in BDT-quinoxaline-based D–A polymers. The first polymer studied has a 2-ethyloxy alkyl side chain attached on the BDT core unit while the second polymer is a 2D system with a 2-ethylhexylthiophene used as side group on the BDT core unit92 (P58 and P59). The polymers are shown in Figure 8.20. The polymers were prepared by a Stille coupling reaction. The effect of the 2D conjugation was first studied by measuring optical absorption behavior of the two polymers. It is revealed that P59 P2 (2D system) gave a red-shifted absorption edge by 26 nm as compared to the polymer that possess an alkoxy side chain (P58 P1). The thermal stability as studied by TGA was found to be quite different with the onset of decomposition for the 2D polymer being 430oC while the alkoxy substituted polymer had onset of decomposition at about 320oC, showing the role of 2D structures in improving thermal stability. The polymers were used in fabricating solar cell devices by blending them with PCBM[70] using the device structure of glass​/ITO/​PEDOT​:PSS/​polym​er:PC​BM[70​]/Ca/​A l. The alkoxy substituted BDT containing polymers gave a PCE of 3.06% with Jsc of 7.0 mA/cm2, Voc of 0.71 V, and FF of 61.5%, while the nearly

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FIGURE 8.20  BDT-quinoxaline-based polymers (P58–P64).

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identical device based on P59 P2 gave a higher PCE of 5.0 %, with Jsc of 10.13 mA/cm2, Voc of 0.76 V, and FF of 64.3%. Clearly, the Jsc was enhanced by the use of 2D conjugated polymer for solar cells, which could originate from the better light harvesting capability due to the redshifted absorption. Moreover, the hole measurements showed a hole mobility of 4 × 10−5 cm2 /Vs and 1.04 × 10−4 cm2 /Vs for P58 P1 and P59 P2, respectively which is another reason for the enhanced Jsc of the 2D-polymer-based solar cell as compared to the alkoxy substituted polymer. The effect of fluorination in BDT-quinoxaline-based polymers has also been studied. The first polymer based on fluorinated quinoxaline and BDT was reported in 2012 by Chen et al.93 (P60). The introduction of the two fluorine atoms to the quinoxaline unit (positions 6 and 7) resulted in a lowered HOMO energy level (−5.5 eV) as compared to the non-fluorinated quinoxaline-based polymers (HOMO = −5.12 eV).92 The polymer was found to aggregate at solid state as shown from the red-shifted absorption of the film as compared to the solution in the UV-visible absorption spectrum. BHJ solar cells with structure glass/ITO/PEDOT:PSS/active layer/Ca/Al were fabricated and characterized. An optimized 1:1 (wt./wt.) P60 (higher molecular weight):PCBM[70] active blend gave a PCE as high as 8% when processed from o-DCB in presence of 3% of 1,8-diiodooctane (DIO). The superior Jsc (18.2 mA/cm2) correlates well with the high hole mobility (2.3 × 10−3 cm2/Vs) of the polymer. In a similar class of materials, the effect of a thiophene or thienothiophene π-bridge on the properties of the polymers were studied (P61 and P62).94 The polymers that possessed thieno[3,2- b]thiophene as a π-bridge (P62) were found to have a more planar and linear backbone, allowing a better π-delocalization with an extended conjugation length and resulting in a red-shifted absorption, lower optical bandgap, and a higher tendency to form ordered and crystalline domains. The photovoltaic properties of the polymers were investigated using a conventional device structure, glass/ITO/PEDOT:PSS/active layer/ LiF/Al. PCEs of 2.18% and 5.6% were achieved for 1:1 (wt./wt.) P61:PCBM[60] and P62:PCBM[60] blends, respectively, processed from o-DCB. The thienothiophene-bridged polymer (P62) showed a doubled Jsc and improved FF, which can be ascribed to the enhanced light absorption, charge mobility, and improved BHJ morphology. The performance of the P61-based device was improved with the addition of 3% DIO (v/v) to the active solution that led to a PCE of 5.7%. The effect of side chain BDT-quinoxaline-based polymers were studied. P63 had 2-octylthiophene as a side group while 2,3-hexyl thiophene was attached to the BDT unit95 in P64 (Figure 8.20). As expected, the polymers showed quite different optical behavior in both solution and solid state. For instance, the polymer with the octyl side chain (P63) showed red-shifted absorption with an absorption edge of 730 nm as compared to 700 nm for the other polymer. The photovoltaic properties of the polymers were studied and gave a PCE of 5.7% and 3.4% for as cast and annealed P63 PFQBDT-TR1 and P64 PFQBDT-T2R2 based devices, respectively. The P63 PFQBDT-TR1-based solar cell yields a current density of 11 mA/cm2, almost double that of the corresponding P64 PFQBDT-T2R2-based device (6.5 mA/ cm2), which is in agreement with the different optical properties, both in pristine or blended films and molecular packing of the polymer chains. 8.2.3.4 BDT with Benzodithiophene-dione Benzodithiophene-dione as an acceptor in constructing D–A polymers together with BDT resulted many successful polymers. One of the first polymers in this class was P6596 (Figure 8.21), which gave a low PCE of 0.73%, probably due to the low charge carrier mobility. Then, a thiophene π-bridge was introduced between the BDT donor unit and the acceptor. In addition, the attachment of the 2-ethylhexylside chain solubilizing alkyl group was employed to give a polymer (P66) (Figure 8.21). Moreover, a thiophene group was added as a side group to the BDT to give a modified polymer.97 A blend of the polymer with PCBM[60] in 1:1 ratio gave a higher PCE of 6.67% with Voc = 0.86V, Jsc = 10.68 mA/cm2, and FF = 72.3%. The non-fullerene acceptors (ITIC-Figure 8.18) have also been used together with (PBDTBDD) (P66).98 The polymer showed a higher tendency of aggregation in both solution and solid state as seen from the continuous reduction of the high wavelength peak when measured at higher temperatures. Such

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FIGURE 8.21  BDT-benzodithiophene-dione-based polymer (P65 and P66).

behavior is good as it will allow pure polymer domains and reduce geminate recombination. The use of ITIC acceptor complements with the optical absorption of the donor polymer and hence more photons are harvested. Moreover, the matching energy levels of the ITIC with the donor material avoids energy losses. The material was tested in organic photovoltaics using ZnO and MoO3 as cathode and anode interfacial layer, respectively. The P66 PBDTBDD:PC71BM control device a PCE of 7.45%, while the device prepared by blending the same polymer with ITIC showed a Jsc of 16.80 mA/cm2, a Voc of 0.899 V, and an FF of 74.2%, yielding a PCE of 11.21%. The electron and hole mobilities were measured to be 3.13 × 10−4 cm2/Vs and 2.10 × 10−4 cm2/Vs, respectively, with low ratio of electron and hole mobility (μe/μh = 1.49), which accounts the reason for the high Jsc and FF. The benzodithiophene-dione acceptor unit was combined with a BDT monomer with an alkoxyphenyl group and with either thiophene or selenophene used as a π-bridge (P67 and P68), and first reported by Li’s group (Figure 8.22).99 On the BDT, an m-ethylhexyloxy phenyl ring was attached to form a 2D-conjugated system. The phenyl substituent is known to give a deeper HOMO energy level and hence higher Voc than similar polymers with thienyl substituents are expected to generate when the polymers are used in solar cells. The use of selenophene as a π-bridge has many advantages. First, due to the lower aromaticity of selenium as compared to thiophene, it can encourage the quinoid formation leading to improved planarity, increased conjugation length and lower bandgap.100–103 Second, due to the bigger size of selenium, it has more polarization of charge, which in turn promotes polymer chain interaction, and hence materials with better charge mobility are prepared. The polymers were prepared by the Stille coupling reaction using palladium as a catalyst. Consistent with previous studies, the selenium-bridged material (P68) showed higher hole mobility and red-shifted optical absorption as compared to the thiophene counterpart (P67). The HOMO and LUMO energy levels were determined from a CV experiment and found to be −5.42/−3.36 eV and −5.35/−3.31 eV for PBPD-Th (P67) and PBPD-Se (P68), respectively. The polymers were used to fabricate photovoltaic devices in a conventional device structure and by using zirconium acetylacetonate (ZrAcac) as a cathode interfacial material. The photoactive films were processed from a blend of ODCB and DIO. P68 PBPD-Se:PC71BM-based PSCs exhibited a significantly improved PCE of 9.8% with an enhanced Jsc of 14.9 mA/cm2 and a slightly lower Voc of 0.90 V and FF of 73% in comparison with a PCE of 8.4% with a Voc of 0.95 V, a Jsc of 12.4 mA/cm2, and FF of 71% for P67 PBPD-Th:PC71BM-based devices. The ratio of hole to electron mobility was found to be 0.97 for P68 PBDP-Se:PCBM[70] while the value

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FIGURE 8.22  BDT-benzodithiophene-dione-based polymer (P67–P69) and IT-4F.

for P67 PBDP-Th:PC71BM was estimated to be 1.22. Obviously, the values indicate that the P68 PBDPSe:PCBM[70] blend film gave a balanced hole and electron mobility, which is beneficial for getting high FF. The higher Jsc obtained from the P68 PBDP-Se-based device is due to the higher crystallinity as seen in XRD with smaller lamella stacking and π-π spacing. By introducing fluorine atoms to the ITIC molecule, a modified non-fullerene acceptor called IT-4F (Figure 8.22) was prepared which has been used as an acceptor together with a donor polymer (P69 PBDB-T-SF) that led to a PCE of 13%.22 The optical absorption of the polymer was recorded, giving an optical edge of 688 nm, with corresponding optical band gap of 1.80 eV. This polymer showed an enhanced absorption main peak as compared to the non-fluorinated polymer, showing the positive role of the fluorine toward increasing intermolecular π-π packing. Moreover, the low-lying HOMO and LUMO levels of P69 PBDB-T-SF and IT-4F may have certain advantages, such as good chemical stability and large polarization, which are beneficial for improving their photovoltaic performance. Organic solar cell devices with an inverted structure of ITO)/ZnO/active layer/MoO3/Al were prepared to investigate the photovoltaic performance of the P69 PBDB-T SF:IT-4F blend film that yielded the highest performance of 13% PCE in single junction devices. An extended, fused BDT (two more thiophenes are fused on the backbone) and benzodithiophenedione acceptor with thiophene used as a π-bridge was developed to give a new polymer (P70) (Figure 8.23).104 When combined with ITIC-Th1 (Figure 8.23), it gave a PCE of 9.6%.

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FIGURE 8.23  Molecular structure of P70 and ITIC-Th and ITIC-Th1.

FIGURE 8.24  The structures of PBnDT-DTBT(0F), DTfBT(1F), and DTffBT (2F) (P71–P73).

8.2.3.5 BDT-Benzothiadiazole Benzothiadiazole is one of the well-known acceptors and has been combined with different donor units. You’s group reported three BDT-BTD-based105 polymers, differing only in the number of fluorine atom substitutions, to see the effect of fluorine substitution in this class of polymers (P71–P73) (Figure 8.24). The structures of the polymers included in the study are shown in Figure 8.24. PBnDT-DTBT (P71) has no fluorine while PBnDT-DTfBT (P72) and PBnDT-DTffBT (P73) contain 1F and 2F attached on the BTD unit of the polymer. The three polymers were synthesized by the Stille coupling reaction. The optical study showed only slightly differing optical bandgaps of 1.65, 1.67, and 1.73 eV for the polymers with 0F (P71), 1F (P72), and 2F (P73) polymers, respectively. The most fluorinated polymer showed higher optical bandgap, which is consistent with previous observations. The polymer with 2F polymer (P73) yields an absorption coefficient of 4.4 × 10−4 cm−1, slightly higher than that of both 0F (P71) and 1F (P72) polymers (∼4.0 × 10−4 cm−1). This slight increase in absorption for 2F (P73) can be attributed to it having larger crystallites with a more face-on orientation. Also, the HOMO of the polymer with more fluorine was found to be deeper (−5.53 eV) as compared to −5.42 eV for the polymer without fluorine atoms, while the polymer with 1F (P72) has a HOMO of −5.48 eV. Photovoltaic devices were prepared by blending the polymers with PCBM[60] in a 1:1 ratio. The polymer with 2F (P73) gave a PCE of 7.16% with Voc of 0.90 V, Jsc of 12.2

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FIGURE 8.25  BDT-BT polymers (P74–P77).

mA/cm2, and FF of 62.1%. The corresponding values for the 1F (P72) polymer were 0.85 V, 11.4 mA/cm2, and FF of 50.6%, resulting in a PCE of 5.28%. The non-fluorinated polymer showed a performance of 4.53% with Voc of 0.78 V, Jsc of 11.7 mA/cm2, and 45.6% FF. The higher Voc from the devices with 2F (P73) as compared to both 1F (P72) and 0F (P71) aligns with the deeper HOMO of the polymer in the electrochemistry study. The Jsc increase with the 2F (P73) could also be accounted for by the higher absorption coefficient of the polymer. The authors also claim that the addition of fluorine particularly improved morphology; the doubly substituted polymer showed a greater face-on polymer and improved π-π stacking along with pure polymer and fullerene domains due to the lower miscibility. Janssen’s group prepared four BDT-BTD-based polymers with thiophene and furan bridging groups and also with and without the fluorination of the BTD (P74–P77).106 The difference between BDT-BT-2T (P74) and BDT-FBT-2T (P75) is the fluorination of the BT unit in the latter and thiophene is used as a π-bridge. P76 and P77 differ in the fluorination of the BT unit in the later and furan is used as π-bridge (Figure 8.25). The polymers were prepared by the Stille coupling reaction from the respective monomers. The optical properties were studied with UV-visible spectroscopy and shows that the thiophene-bridged polymers have relatively red-shifted absorption than the furan counterparts. While the four polymers show similar LUMO energy levels as determined by CV, the HOMO of the polymers are found to be different with the fluorinated polymers (P75 and P77), giving deeper HOMO as compared to the non-fluorinated polymers (P74 and P76). The performance of the polymers in solar cells was tested by combining with PCBM[70] processed from ODCB for the three polymers (CB/CN was used for P75 BDT-FBT-2T) with a device structure of ITO/PEDOT:PSS /polymer:fullerene/LiF/Al. The thiophene-bridged polymers behaved better in solar cell than the furan-bridged polymers. The highest device efficiency recorded was 7.7% and came from a blend P75 BDT-FBT-2T and PCBM[70]. The other polymers showed efficiency in the range from 3.7 to 5.4% PCE. One reason for the higher performance of the P75 BDT-FBT-2T-based device is partly due to the high hole mobility (9.3x10−2 cm2/Vs). The BDT-BTD-based polymer (P78)107 (Figure 8.26) was developed and used as donor polymer in solar cell together with PCBM[70]. An optimized device gives a PCE of 8.30% together with PCBM[70] in 1:1.5 blend ratio.

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FIGURE 8.26  Structure BDT-BTD polymer (P78).

FIGURE 8.27  Structures of BDT-BTz-based polymers (P79–P82).

8.2.3.6 BDT-triazole Polymers BTz is a weak electron acceptor as compared to the most commonly used BTD unit due to the basic, central nitrogen atom that donates an electron to the ring through resonance. Thus, the LUMO of the resulting polymer, which most of the time is located on the acceptor unit, will be raised resulting in a bandgap increase (high bandgap polymer). The basic nitrogen atom of the BTz unit is also useful since solubilizing side chains can be easily attached. Figure 8.27 shows two polymers with fluorinated BTz or non-fluorinated BTz combined with BDT108 developed by You’s group and used to fabricate photovoltaic devices (P79 and P80). The absorption edge of the fluorinated polymer is 2.0 eV while the nonfluorinated polymer has an optical bandgap of 1.98 eV. Fluorination lowers the HOMO by 0.07 eV, with the HOMO of the fluorinated polymer being −5.36 eV according to the CV of the polymer. Similarly, the LUMO of the fluorinated polymer is lowered by 0.18 eV with the LUMO of the fluorinated polymer being −3.05 eV. The photovoltaic properties of the two polymers were studied by blending with PCBM[60]. A device prepared from a blend of P80 PBnDT-FTAZ:PC61BM-based BHJ cells, showed a Voc of 0.79 V, a Jsc of 12.45 mA/cm2, and a very notable FF of 72.2%, leading to the highest overall power conversion efficiency of 7.1% with an active layer thickness of 250 nm. Interestingly, the polymer gave a PCE of 6% despite a thicker film (1 micrometer). On the other hand, the non-fluorinated polymer gave a maximum power conversion efficiency of 4.36%. The higher power conversion efficiency of the fluorinated polymer (P80) could be due to the higher hole mobility of the devices as compared to the non-fluorinated polymer (P79) (1.03 × 10−3 cm2/Vs vs. 2.94 × 10−4 cm2/Vs). The performance of FTAZ (P80) combined with a fluorinated and non-fullerene acceptor (ITIC-Th1 and ITIC) (Figure 8.23) was tested in a solar cell in order to exploit the contribution of the non-fullerenated acceptors toward charge carrier formation. One of the acceptors was fluorinated to improve molecular interactions, which is essential to improve Jsc and FF. The use of ITIC-TH as an acceptor gave a PCE of 8.88% while the fluorinated acceptor (ITIC-Th1) gave a PCE of more than 12.1%109 due to the improvement of Jsc and FF at the same time. The polymer and ITIC-Th1-blended film gave a

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higher hole and electron mobility while the ratio between these two values are smaller as compared to the relatively lower hole and electron mobility of the polymer and ITIC-Th blended films with the ratio being higher. The thioalkyl-substituted BDT has been combined with a mono-fluorinated BTz with either thiophene (P81) or furan (P82) as a π-bridge to give two high bandgap D–A polymers by the Peng group110 (Figure 8.27). The polymers were characterized by optical and electrochemical techniques to reveal the suitability of the materials for solar cell applications. The optical study showed a well-defined absorption pattern with absorption peaks that arise from the ordered nature of the polymer chains in both solid and solutions. Such behavior is good for devices to perform well as it promotes better charge carrier mobility. The absorption edges were found to be 1.83 and 1.81 eV for the furan (P82) and thiophene (P81) bridged polymers, respectively. The slightly lower bandgap of the thiophene-bridged polymer indicates the potential of the thiophene unit in promoting conjugation length. The HOMO and LUMO of the thiophene-bridged polymer (P81) were −5.32 and −3.26 eV while the corresponding values for the furanbridged (P82) polymers were −5.38 and −3.28 eV. The HOMO of both polymers were found to be deeper than the HOMO of the polymers with alkyl substituted BDT containing polymer previously reported by the You group.108 The thiophene-bridged polymer (P81) gave a PCE of 7.74% (Jsc = 12.36 mA/cm2, Voc = 0.88 V, FF= 71.2%) when blended with PCBM[70]) in a single junction regular device when used as donor material in organic solar cells. The furan-bridged polymer (P82) gave a PCE of 6.25% with slightly lower Jsc = 11.81 mA/cm2 and FF=58.2% and higher Voc = 0.91 V as compared its counterpart. From TEM study on the morphology of the devices, the thiophene-bridged system gave a well-ordered morphology that improved the Jsc and FF. The thiophene-bridged polymer (P81) has been used as a high band gap donor in a tandem device structure together with a low a bandgap DPP-based polymer, which gave an impressive PCE of 9.4%.

8.2.4 Indacenodithiophene and its Derivatives as a Donor Unit in the Construction of Donor-Acceptor Polymers The indacenodithiophene (IDT) unit has an extended fused-ring and rigid coplanar structures, which are beneficial for p-electron delocalization and preventing rotational disorder, which reduces reorganization energy and enhances charge carrier mobility.111 Besides a well-developed synthetic methodology for the synthesis of the core unit (Scheme 8.10), the IDT unit has reaction sites for easy modifications. For instance, alkyl side chains can be easily attached to the bridging atom for enhancing solubility and tuning the properties of the final materials. Another key advantage of the IDT unit is the possibility of introducing atoms like silicon, nitrogen, and germanium and the possibility of functionalization to make it ready for coupling reactions. In order to extend the conjugating further, groups like thienothiophene can be used instead of thiophene. In this section, few D–A polymers based on an IDT donor and derivatives will be discussed. For a complete overview on this class of polymers, the following review papers are suggested for the readers.13,64 8.2.4.1 Functionalization of the Bridging Atom A series of indacenodithiophene-based donor polymers consisting of carbon, silicon, and germanium as a bridging atom were synthesized with different molecular weight and polymer purification methodology in order to study the influence of the bridging atom and the purification technique (P83–P85)112 (Figure 8.28). The acceptor used was BTD. After Stille coupling, the polymers were purified by Soxhlet extraction as well as recSEC. As seen with other class of materials, the bridging atom affects molecular packing and other optoelectronic properties of the polymers. DSC study, for instance, showed that both the silicon- (P84) and germanium-bridged (P85) polymers showed a melting and crystallization peak while the P83 IDT-BT polymer happened to be amorphous with no obvious thermal transitions. The molecular energy levels of the polymers were found to be closely similar, with the HOMO being in the range of −5.2 to −5.3 eV while the LUMO was in the range of −3.5 to −3.6 eV. To evaluate the photovoltaic

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SCHEME 8.10  Synthesis of the IDT core unit.

performances of the polymers, BHJ solar cells with a conventional device structure consisting of ITO/ PEDOT:PSS/Polymer:PCBM[70]/Ca/Al were prepared and tested under simulated solar light. The high molecular weight germanium-bridged polymer (P85) gave a PCE of 6.5% with corresponding Jsc = 13.95 mA/cm2, Voc = 0.85 V, and FF = 55%, which is comparable to the carbon-bridged polymer (P83). The highest performance for the high molecular weight silicon-bridged polymer (P84) was 5.6% with corresponding Jsc = 12.68 mA/cm2, Voc = 0.88 V, and FF = 50%. Grazing-incidence X-ray diffraction (GIXD) studies showed that both the silicon- (P84) and germanium-bridged (P85) polymers have weak scattering peaks associated with the lamellar repeat between molecules. However, the germanium-bridged materials are slightly stronger than the silicon-bridged polymer. Jen’s group developed a selenium counterpart replacing thiophene with selenium113 and combining it with difluorobenzothiadiazole (DFBT) (P87). The structure is shown in Figure 8.28. In the optical study, the selenium-bridged polymer (P87) showed red-shifted absorption. Selenium is known for higher polarizability due to the reduced aromaticity compared to thiophene. This quinoid character is encouraged with the selenium version and increased conjugation and decreased bandgap. In fact, the selenium version showed a 27 nm red-shift as compared to the thiophene. The HOMO and LUMO energy levels were estimated from the onsets of the oxidation and reduction waves. The selenium ring in P87 PIDSe-DFBT had little effect on the HOMO while it lowered the LUMO by 0.08 eV as compared to P86 PIDT-DFBT. The hole mobility measurement showed that the high molecular weight P87 PIDSeDFBT showed improved mobility as compared to both the low molecular weight component of the same polymer and P86 PIDT-DF, which indicates the improved packing of the P87 PIDSe-DFBT polymer chains in the solid state. The effect of selenium substitution on the photovoltaic properties of the materials was investigated using a device configuration of ITO/PEDOT:PSS/polymer:PC71BM/Bis-C60/Ag, giving a 6.79% PCE from the high molecular weight of P87 PIDSe-DFBT, while the highest efficiency from the P86 PIDT-DFBT-based device delivered 6.02% PCE.

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FIGURE 8.28  IDT-BT-based polymer structures (P83–P88).

8.2.4.2 Further Extension of the Fused System In order to enhance charge mobility and device performance of the fused-ring polymers, the IDT unit can be further extended by incorporating two thieno[3,2-b]thiophene (TT) units. The extended IDT system with two outward TT units replacing the thiophene moieties on IDT was designed to form a novel seven-ring indacenodithieno[3,2-b]thiophene (IDTT) donor unit. The copolymerization of IDTT with DFBT produced a polymer (P88).114 The PCE improved to 7.03% from 5.97%. The charge mobility improved in PIDTCPDT-DFBT (P89) due to the increase in the effective conjugation length and planarity 111111 Indeed, the device based on P89 PIDTCPDT-DFBT showed a respectable PCE of 6.46% with a Voc of 0.75 V, Jsc of 14.59 mA/cm2, and FF of 59%. Finally, a ladder-type polymer with a germanium atom as a bridging group in which two thienothiophene units are held together by the bridging group has been prepared by the Heeney group115 (P90) and used to prepare photovoltaic devices that gave a PCE of 7.2% (Figure 8.29).

8.2.5 Summary and Outlook D–A polymers were intensively investigated for optoelectronic applications due to advantages such as easy synthetic approaches for the development of materials with desired optical, electrochemical, and

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FIGURE 8.29  Structure of P89–P90.

device properties via chemical modification of the monomers. Moreover, solution processability of polymers and the prospect of using them in larger scale photovoltaics production primarily on flexible substrates by roll-to-roll printing was a reason for greater focus and deeper investigation of D–A polymers in the past few decades. Hence, an abundant library of D–A polymers was developed and tested in organic photovoltaics. Through the synthesis of novel conjugated polymers with desired properties and careful optimization of devices prepared from these materials blended with fullerene and nonfullerene-based acceptors, impressive results were achieved. D–A-based polymers that gave as high as 13% PCE together with a non-fullerene acceptors in single junction devices were reported. Even though the improvement in organic photovoltaics in the past few decades was immense, there are still challenges to be overcome for the field to have a substantial contribution toward generation of cleaner and cheaper energy:

1. The design and synthesis of novel materials needs to continue with a focus on materials that are high-performing, but also on simple polymers with cheaper production cost. 2. New, improved, and more environmentally friendly synthetic methods need to be developed. 3. D–A polymers with higher stability in ambient conditions should be developed and tested. 4. The focus of organic photovoltaic research in the past was mainly on improving the PCE in a laboratory environment: a small device area processed from chlorinated solvents by spin coating mostly processed inside glove box under inert atmosphere. Environmentally friendly processing techniques and large area device fabrication needs to be emphasized in the future.

Acknowledgments The authors would like to thank Flinders University and the Australian Research Council (DP160102356) for financial support.

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9 Conjugated Polymers for n- and p-Type Charge Transport 9.1 Introduction.......................................................................................325 9.2 p-Type Charge Transport.................................................................326 9.3

Zachary S. Parr, Zhijie Guo, and Christian B. Nielsen

Polythiophene-Based Systems • Donor-Acceptor Systems • Molecule:Polymer Blends

n-Type Charge Transport.................................................................359 Indigo- and Isoindigo-Based Systems • Diketopyrrolopyrrole-Based Systems • Rylene Diimide-Based Systems • Other Structural Systems

9.4 Ambipolar Charge Transport..........................................................395 9.5 Conclusions and Outlook................................................................ 400 References........................................................................................................411

9.1 Introduction Electronic charge transport is a fundamental process of great importance in most areas of organic electronics whether one is studying, for instance, organic light-emitting diodes (OLEDs), organic photovoltaic (OPV) devices, or organic field-effect transistors (OFETs). Several techniques and devices have been used to investigate charge transport properties in π-conjugated organic materials with the most widely recognized technique for benchmarking charge transport being the field-effect transistor. Ando and coworkers from Mitsubishi Electric Corporation reported the first organic polymer-based field-effect transistor in 1986, using electropolymerized polythiophene prepared from 2,2′-bithiophene (Tsumura, Koezuka, and Ando, 1986). After electrochemical de-doping of the as-deposited polymer, an electrical field-effect was observed giving rise to an extracted hole mobility around 10−5 cm2/Vs and a transconductance of 3 nS. More than 30 years have passed since this report, and the field of organic electronics has matured significantly with market-ready applications such as e-paper displays and electronic sensors, not least driven by the vast improvements in understanding and control of charge transport properties in organic polymers and the shift to solution-processable materials facilitating the use of high-resolution printing technologies (Fukuda and Someya, 2017). In this chapter, we will discuss these significant advancements over recent years that have seen OFETs surpass amorphous silicon-based thin film transistors in terms of performance and steadily approach high-end polycrystalline silicon and oxide-based devices. We will focus on materials development and highlight fundamental considerations and significant advances in the design of new π-conjugated materials with a specific focus on π-conjugated polymers and various blends thereof. We will discuss underlying design strategies toward efficient charge transport for both n- and p-type materials covering, for 325

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example, backbone modifications including extended polycyclic aromatic motifs and donor-acceptor structures, side chain modifications, substitution of sulfur with heavier chalcogens, and incorporation of non-covalent conformational locks. We will furthermore discuss other aspects that play an important role for charge transport in polymeric systems including the role of polymer molecular weight, impurities, additives, and device processing conditions. The chapter is divided into three sections covering p-type (hole transport), n-type (electron transport) and ambipolar charge transport materials and further subdivided by various materials classes that have been of particular interest to the community over the last decade. We have chosen to focus on charge mobility as the defining metric for high-performance polymer systems and have chosen a cross section of systems focusing mainly on p-type materials with maximum mobilities in excess of 1 cm2/Vs and n-type materials with mobilities above 0.1 cm2/Vs. Although we have chosen to define high mobility as our benchmark for this review, there has been some debate over exaggeration of charge mobility values where some effects have led to significant non-ideality in transfer curves and some hysteresis in cycling of the devices (Choi, Cho, and Frisbie, 2018). Charge mobility is one of the defining metrics for semiconducting materials that limits the efficiency of any devices synthesized from these materials; however, mobility extraction has been a challenging and contentious topic. Ordinarily, mobility is extracted from the transfer characteristics of a thin film transistor using equations derived for a perfect semiconducting MOSFET-type device. Organic devices suffer from significant shortcomings where extraction of mobility from these equations can lead to both over- and under-estimation of charge mobility. The most significant problems with this approximation are the presence of contact resistance from Schottky barriers, which cannot be accounted for fully using these equations but nevertheless cause significant barriers to charge transport and have been suggested to have a 10-fold increase in estimation of carrier mobility and, conversely, resistive contacts where there is insensitivity to gate voltage resulting in large underestimation of mobility, which can occur regardless of characteristic kinks in the transfer curves (Liu et al., 2017). Specific details of the extraction of mobility for conducting polymers and the attempts made to mitigate the effects of these problems will be discussed more thoroughly in Chapter 1 (Volume 2). We will, however, attempt to highlight examples where significant non-idealities have been reported for certain structures throughout this chapter.

9.2 p-Type Charge Transport 9.2.1 Polythiophene-Based Systems Arguably one of the most thoroughly investigated π-conjugated polymers, poly(3-hexylthiophene) (P3HT) has been studied in detail since the early 1990s when synthetic methodologies were developed to control the regioregularity (McCullough and Lowe, 1992; McCullough et al., 1993; Chen and Rieke, 1992). Maintaining a high degree of regioregularity in P3HT by avoiding head-to-head and tail-to-tail couplings, both regarded as structural defects, facilitates a structurally ordered solid-state packing with a predominantly edge-on texture and hole mobilities of up to 0.1 cm2/Vs in the direction parallel to the substrate as tested in an OFET configuration (Sirringhaus et al., 1999). Kline and coworkers found that thin film crystallinity decreased with increasing molecular weight, thereby affording microstructures with less defined grain boundaries and higher charger carrier mobilities (Kline et al., 2003). Recent work by Reichmanis and coworkers further highlights the importance of microstructural ordering for efficient charge transport in P3HT films (Kleinhenz et al., 2016). Through a pre-aggregation approach in solution, induced by sonication and subsequent ageing, they were able to form micrometer-long P3HT fibers that could be solution cast into thin films with fewer grain boundaries and mobilities up to 0.15 cm2/Vs. Using a solution-deposited polymer-electrolyte gate dielectric consisting of polyethylene oxide (PEO) and lithium perchlorate, Frisbie and coworkers initially afforded hole mobilities up to 0.70 cm2/Vs for P3HT and subsequently up to 3.4 cm2/Vs (Panzer and Frisbie, 2006 and 2007). Switching to an iongel dielectric comprising a PEO-polystyrene triblock copolymer and an ionic liquid, they subsequently

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reported average saturated mobilities up to 1.8 cm2/Vs with high on/off ratios of 105 (Cho et al., 2008). Meanwhile, maximum saturated hole mobilities of 0.4–0.5 cm2/Vs are generally among the highest values reported for P3HT with conventional FET devices (Baeg et al., 2010; Fei et al., 2015). Nawaz and coworkers have more recently worked in detail with so-called defect-free P3HT samples that do not contain any head-to-head or tail-to-tail couplings according to nuclear magnetic resonance (NMR) spectroscopic analysis (Nawaz et al., 2016). Using a cross-linked poly(vinyl alcohol)-based dielectric material, they afforded a maximum hole mobility of 1.20 cm2/Vs which was further increased to 2.8 cm2/Vs upon PEDOT:SS treatment of the dielectric; although, the increased mobility was associated with a significant drop in the on/off ratio. In a follow-up study, the authors used a floating film transfer method to create aligned films of defect-free P3HT with a dichroic ratio around 2.5 (Nawaz, Kumar, and Hümmelgen 2017). When tested in OFETs, mobilities as high as 8.0 cm2/Vs were reported with average values around 6.3 cm2/Vs. It should be noted that the reported mobilities are extracted from non-linear transfer curves and, as such, there are some doubts about the effective carrier mobilities, though the data clearly suggests that increasing the regioregularity and affording aligned films are beneficial for the charge transport properties. Using a nanoimprinting approach, Barbero and coworkers fabricated a P3HT film with micrometer sized features and subsequently measured an average vertical mobility around 3 cm2/Vs in those P3HT pillars, which is an improved of more than four orders of magnitude compared to the vertical transport in a non-patterned film of P3HT (Skrypnychuk et al., 2016). The efficient charge transport in this system is ascribed to predominant intra-chain transport due to effective alignment of polymer chains in the vertical direction during nanoimprinting. Although these latter results are not immediately transferable to a cheap and scalable manufacturing process, they show that control of microstructure and the elimination of significant grain boundaries are the crucial aspects for achieving efficient charge transport in P3HT-based systems. In 2017, Park and coworkers sought to address the issue with P3HT grain boundaries in the solid state from a molecular design perspective (Son et al., 2016). They simply reduced the side chain density in P3HT by replacing in a random fashion of up to 33% of the 3-hexylthiophene units with unsubstituted thiophene. Although the overall crystallinity in the solid state is reduced, the lower side chain density affords a more planar polymer backbone that gives rise to enhanced aggregation and charge transport across grain boundaries. Consequently, the average saturated mobility increased from 0.17 cm2/Vs for P3HT to 1.31 cm2/Vs for the random copolymer with 33% unsubstituted thiophene moieties. Lower feed ratios of thiophene:3-hexylthiophene (17% and 25%) also gave rise to significantly enhanced mobilities around 0.9 cm2/Vs. Rather than reducing steric interactions between side chains by partly removing them, Heeney and coworkers explored synthetic approaches to rigidify the polythiophene backbone through fluorination (Fei et al., 2015). Investigating poly(3-octylthiophene) (P3OT) and its fluorinated analogue F-P3OT rather than P3HT due to issues with solubility upon fluorination, they found that fluorination did indeed result in a higher degree of backbone planarization and thus solid-state aggregation. On the other hand, fluorination did not increase the overall degree of crystallinity in the solid state, which could potentially enhance detrimental grain boundary effects. The enhanced backbone planarization and interchain interactions are reflected in an increase in the average linear mobility from 0.14 cm2/Vs for P3OT to 0.70 cm2/Vs for F-P3OT. Although the examples above exploring synthetic approaches to planarize the polythiophene backbone have emerged over the last couple of years, this line of thought was explored much earlier with polythiophene systems such as PQT12, PBTTT, and PC12TV12T, as well as other closely related fused thiophene systems (see Figure 9.1) (Ong et al., 2004; McCulloch et al., 2006; Kim et al., 2011; Fong et al., 2008; Biniek et al., 2013). Here, the side chain density was decreased in a more systematic and regioregular fashion by alternating a dialkylated bithiophene unit with either an unsubstituted bithiophene (PQT12) or thienothiophene (PBTTT) unit, while PC12TV12T introduced an unsubstituted vinylene unit for every two dialkylated bithiophenes. PBTTT, in particular, has been thoroughly studied as a model system due to its liquid crystalline properties, which allow for thin film processing from the mesophase thereby affording much larger crystalline domains than observed with thin films of P3HT for example.

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FIGURE 9.1  Representative thiophene-based polymers for high-performing p-type OFETs.

McCulloch and coworkers initially reported a maximum saturated hole mobility of 0.72 cm2/Vs while Müllen and coworkers afforded maximum values around 1.3 cm2/Vs for highly aligned microstructures fabricated by dip-coating (Wang et al. 2012b). Similarly, PC12TV12T was found to exhibit a maximum hole mobility around 1.05 cm2/Vs and a relatively high solubility in common organic solvents due to the four dodecyl chains per repeat unit. Frisbie and coworkers moreover applied their previously mentioned polymer electrolyte dielectrics to PQT12-based devices and achieved a linear mobility of 0.9 cm2/Vs with the PEO/perchlorate system and an average saturation hole mobility of 1.6 cm2/Vs with the ion-gel dielectric (Panzer and Frisbie, 2007; Cho et al., 2008). While the polythiophenes discussed above generally have ionization potentials (IPs) around 4.8–5.1 eV, recent work has focused on structural modifications that lower the IP without adversely affecting the charge transport properties for applications such as organic electrochemical transistors and thermoelectric generators. In this context, McCulloch, Rivnay, and coworkers have explored polythiophenes decorated with oligoethylene glycol sidechains (Figure 9.2) rather than alkyl chains to facilitate not only electric but also ionic charge transport in the solid state (Nielsen et al., 2016) The direct attachment of oxygen atoms onto the polythiophene backbone has a planarizing effect through intramolecular S-O interactions with neighboring thiophenes, while the IP is simultaneously lowered significantly to roughly 4.4 eV due to the electron donation by resonance into the conjugated π-system. The low IPs of these polymers prevented fabrication of reliable OFETs; however, charge carrier mobilities of 0.28 and 0.95 cm2/Vs were extracted from organic electrochemical devices for g2T-T and g2T-TT, respectively. Work by Katz and coworkers synthesized a series of electron rich PQT12 derivatives with thioalkyl chains that replaced the alkyl chains and/or 3,4-ethylenedioxythiophene (EDOT) units that replaced the unsubstituted thiophene units (Li et al., 2017). The thioalkyl substitution reduced the IP from 5.1 to 5.0 eV, while EDOT substitution lowered the IP to 4.8 eV; the effects were additive as both substitutions combined lowered the IP to 4.7 eV. These materials were designed for investigations of thermoelectric properties, and the authors found high conductivities of 140–350 S/cm upon doping with NOBF4 and F4TCNQ despite the fact that relatively low hole mobilities of 1 × 10−2 and 2 × 10−3 cm2/Vs were found for un-doped samples of PQTS12 and PDTDE12, respectively.

9.2.2 Donor-Acceptor Systems The concept of a donor-acceptor (D–A) system is a simple one and is borne from the need to modulate the highest occupied molecular orbital level (HOMO) and lowest unoccupied molecular orbital (LUMO)

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FIGURE 9.2  Representative thiophene-based polymers for emerging applications such as thermoelectric generators and electrochemical transistors.

level to achieve a tunable band gap, which facilitates absorption and photo-excitation across more of the visible and near infra-red spectrum in order to access solar energy allowing for more efficient solar cells. This has led to improvements in materials design for other applications like OLEDs, OPVs, OFETs, and OECTs. A variety of structural backbone or side chain modifications have led to improvements in morphology, solubility, and charge carrier mobility. Donor–acceptor polymers consist of an electron donating monomer with higher HOMO level—often an electron rich thiophene-based moiety— and an electron-withdrawing monomer with some of the most successful being the diketopyrrolopyrrole (DPP), benzothiadiazole (BT), and isoindigo (IIG) derivatives. An alternating copolymer of the two monomers results in π-orbital hybridization where the resulting HOMO is raised in energy relative to the HOMO of the donor moiety and the resulting LUMO is lowered; in energy relative to the LUMO of the acceptor moeity, which gives the polymer a modulated band gap narrower than that of either of the two monomers. This orbital structure is observed as an intramolecular charge transfer band and is generally observed in the UV-Vis absorption spectrum at a higher wavelength and more intense absorption than the absorption band from the π-π* transition (Fei et al., 2017). Further introduction of heteroatoms into the backbone is also beneficial as it has a planarizing effect through non-bonding electrostatic interaction of heteroatoms on adjacent monomers, resulting in a reduction in torsion of the backbone, increasing the orbital hybridization, and promoting improved interchain packing (Jackson et al., 2013). By regioregular and random copolymerization of the two monomers, many high-performance materials have been achieved with a huge variety of donor acceptor moieties. 9.2.2.1 CPDT-Based Systems Cyclopentadithiophene (CPDT)-based systems have been well studied as a potential high charge transport motif over the last few years. The CPDT unit consists of a 2,2’-bithiophene unit fused in the 3,3’-positions with an sp3 hybridized carbon. This has a planarizing effect and hence reduces reorganization energy of the polymer when processed from solution (Zhang et al., 2007). It also enables improved π-stacking , which facilitates improved charge transport between π-conjugated polymer chains with improved π-π interactions leading to improved charge hopping between polymer chains (Qiu et al., 2016). The sp3 carbon is then substituted with various solubilizing side chains to enable processability. When CPDT is copolymerized with BT, it has produced some of the highest field effect mobilities for

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a semiconducting polymer, 21.5 cm2/Vs, with considerable engineering of material morphology (Luo et al., 2014). The polymer CPDT-BT was one of the earlier and relatively simple systems to be synthesized as a donor-acceptor-type conjugated polymer and could be directly solution cast, reporting saturated hole mobility of 0.17 cm2/Vs after annealing, at a time when few systems exceeded 0.1 cm2/Vs (Zhang et al., 2007). Despite this improved charge mobility, the polymers were observed by wide angle X-ray scattering to be extremely amorphous with no higher order diffractions for lamellar (100) crystals organization. Only extremely diffuse diffractions in the π-π (010) plane were observed with a packing distance of 3.7 Å, contrary to popular wisdom for promoting charge transport in P3HT and PBTTT polymers at the time, which suggested increased crystallinity was critical to performance. Mobility for this polymer has since been improved upon by tighter control of molecular weight (Tsao et al., 2009). In this work, the authors were able to produce CPDT-BT polymer samples in two molecular weights of Mn = 12 kDa and 50 kDa, which yielded a saturated hole mobility of 0.67 cm2/Vs for a spin-coated sample of higher molecular weight sample, and a dip-coated OFET device yielded maximum saturation mobility values of 1.4 cm2/Vs. When examined by GIWAXS, the polymer exhibited a π-stacking of 3.9 Å and a considerable anisotropy for the dip-coated sample, suggesting that this method yielded good organization of the polymer. AFM experiments showed small, circular self-assembled areas with grain boundaries for the spin-coated sample and aligned chains into fibrous structure, which facilitated high charge transport along the direction of the polymer chains. A high dependence on molecular weight exhibited by CPDT-BT systems was shown by Tsao et al (Tsao et al., 2011), where they we able to obtain a mobility of 1.2 cm2/Vs for a polymer of Mn = 51 kDa, measured by GPC in THF, analogous to the previous reports. The authors also measured the polymer by GPC in 1,2,4-trichlorobenzene (TCB), where they extracted a molecular weight of 25 kDa for this polymer and were able to measure a polymer of higher Mn of 35 kDa, exhibiting a mobility of 3.3 cm2/Vs. CPDT-BT polymers exhibit close to model linear behavior in transfer curves with very little hysteresis or charge-injection limited inhomogeneities in transfer curves, which have been exhibited by many high-performance polymers to date (Choi, Cho, and Frisbie, 2018). CPDT-BT-based polymers have shown extremely high charge mobility through careful engineering of polymer morphology. Wang et al (Wang et al., 2012a) were able to achieve an average mobility of 4.4 cm2/Vs with a high mobility of 5.5 cm2/Vs for single fiber measurements grown via drop casting of a polymer from low concentration in a solvent-rich atmosphere allowing for high molecular organization. While this method is easily able to afford a high mobility and very effective polymer fibers that display near ideal transfer characteristics, it is very unlikely to be industrially applicable since the fibers must be individually manipulated and the throughput is low. A number of attempts have been made to modify the conjugated backbone of CPDT-BT in order to increase polymer planarity and improve packing and charge transport. One of the most successful methods has been to introduce regioregular polymers of CPDT and pyridal[2,1,3]thiadiazole (PT, Figure 9.3), which has exhibited some of the highest charge mobilities for this type of polymer for certain orientations of PT unit (Ying et al., 2011). Polymers of two regio-isomers of CPDT-PT with the nitrogen atoms oriented head-to-head or head-to-tail with respect to each other (Figure 9.3) as well as a regiorandom system were synthesized in comparable molecular weight. The regioregular polymers showed a much higher charge mobility of 0.6 cm2/Vs and 0.4 cm2/Vs for head-to-head and head-to-tail polymers, respectively, whereas regiorandom polymers exhibited a charge mobility of two orders of magnitude lower. The comparatively low molecular weight of the polymers would likely explain the lower field effect mobilities of CPDT-PT; additionally, when synthesized in a higher molecular weight of 300 kDa, field effect mobilities of around 2 cm2/Vs were extracted (Tseng et al., 2012) for a device with no surface modifications. Much higher mobilities of a 6.7 cm2/Vs were found for macroscopic organization of polymer chains on a nano-groove surface, with the increased mobility for the spin-coated device of higher molecular weight polymers likely due to superior interconnection of crystal grains resulting in reduced grain boundaries and therefore higher mobility. High mobility OFETs from solution processed CPDT-BT have been achieved by nano-grooving of the polymer surface to enhance molecular order and, therefore, charge transport. Using the nano-grooving

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FIGURE 9.3  CPDT-PT polymer structures synthesized. (From Ying et al. 2011. “Regioregular Pyridal[2,1,3]thiadiazole π-Conjugated Copolymers.” Journal of the American Chemical Society 133 (46): 18538–41.)

method of modifying substrate surface morphology to facilitate alignment of polymer chains on the surface and utilizing a capillary action method of depositing the polymer on the surface from the solvent to further enhance polymer chain alignment, a maximum hole mobility of 21.3 cm2/Vs was extracted for CPDT-PT with Mn = 140 kDa and 18.5 cm2/Vs for CPDT-BT (Luo et al., 2014). Comparatively, very organized crystal packing could be observed in GIWAXS with a π-stacking of 3.5 Å and several higher order diffractions in the plane of the polymer indicative of high crystallinity, as well as no low-angle peaks in the out of plane direction, suggesting a high degree of chain alignment. Tseng et al. (Tseng et al., 2014) were able to achieve a maximum charge mobility of 23.7 cm2/Vs for CPDT-PT in the saturation regime using a method of scratching nano-grooves on an OFET substrate surface before deposition of the polymer. Interestingly, when drop cast, the authors found a typical molecular weight dependence of charge mobility exhibiting a positive correlation between molecular weight and charge mobility, whereas for the scratched substrate the highest mobility was obtained for a molecular weight of 50 kDa, and the second highest for a molecular weight of 300 kDa with sharp diffraction peaks for π-π and lamellar crystal orientations. Very low activation energy for charge transport suggested efficient interchain charge hopping in fibers and grain boundaries due to molecular organization. Band-like charge transport has been reported for CPDT-BT polymers with very a high charge mobility of 6.5 cm2/Vs (Yamashita et al., 2014). The polymers were processed via drop casting polymer solution on the surface of an ionic liquid and subsequent compression of the polymer into a film that was deposited on a silica OFET substrate after which electrodes were evaporated onto the polymer. Molecular organization was facilitated by floating the polymer on the surface of ionic liquid at a temperature of 120˚C, which allowed high polymer flexibility and uniaxial orientation of polymer chains. It was then compressed and deposited on a substrate. The films showed high anisotropy in GIWAXS and cross-polarized optical microscopy with polymer chains oriented parallel to the channel; however, the lack of strong π-stacking diffraction suggested disorganized packing in the polymer fibers. Inverse temperature dependence measurements of mobility at elevated temperatures of around 300 K suggested that, at least in part, charge carriers obeyed a band-like model. Magnetic field-dependent hall voltage measurements showed charge carriers with dependence on magnetic field and a high coherence from calculation of the Hall coefficient. The data indicated that the polymer is on the transition between hopping dominated charge transport and band dominated charge transport at least at higher gate voltages and room temperature whereas at lower voltages transport is dominated by a hopping model (Yamashita et al., 2016). Modification of side chains has been a successful method to achieve higher performance OEFT materials via control of solid-state packing and morphology. Interestingly, the unbranched 1-hexadecyl (C16) chain has exhibited the highest molecular weights and mobilities whereas branched side chains synthesized by Tsao et al. (Tsao et al., 2011) showed a highest mobility of 0.4 cm2/Vs for a 3,7-dimethyloctyl (DMO, Figure 9.4) side chain in a polymer of 16 kDa with similar dependence on molecular weight and π-stacking distance of 3.7 Å for both C16 and DMO side chains. In an effort to improve the solubility of higher molecular weight CPDT-BT, polymers containing 2-ethylhexyl (2EH) and 5-ethylnonyl (5EN) side chains (Figure 9.4) were synthesized (Lee et al., 2015a) in Mn of 21.0 kDa and 38.0 kDa and PDI of 1.60 and 2.03, respectively, with only 5EN. Both polymers exhibited lower charge mobilities of 0.004

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FIGURE 9.4  Side chain modifications of CPDT-BT.

cm2/Vs and 0.1 cm2/Vs for 2EH and 5EN, respectively, with only 5EN displaying a diffuse π-π diffraction at 3.65 Å, with qualitative observation by AFM of the fiber network similar to C16-CPDT-BT polymer; although, it is likely that this might be improved upon by careful control of the film morphology and molecular weight. Modification of polymer side chains with alkene induced rigidity and distorted polymer side chain packing (Hinkel et al., 2014). Two polymers were synthesized with a cis- and transalkene moiety in the side chain (Figure 9.4), which displayed mobilities of 0.39 cm2/Vs and 0.61 cm2/Vs. Despite the increase in steric disorder of the side chains, there was very little effect on mobility despite a more disorganized side chain packing with less interdigitation of side chains with both polymers having comparable π-stacking distances of 3.7 Å. The slight difference in mobility was attributed to differences in molecular weight between the two polymers by the authors as a result of the cis-isomer being more soluble than the trans-isomer. The synthetic protocol chosen for preparing CPDT-based polymers is also important with higher mobilities observed for polymers prepared via palladium-catalyzed crosscoupling reactions whereas oxidative coupling via FeCl3 and electrochemical polymerization yielded lower molecular weight oligomer-like materials with three orders of magnitude lower charge carrier mobilities (Horie et al., 2012). Efforts have been made to modify the CPDT unit to modulate performance in OPVs and OFETs via modification of a bridging atom with a heteroatom of silicon, germanium, or nitrogen, with limited improvements on hole mobility. Polymers of Si-CPDT-BT showed mobility of 0.01 cm2/Vs for a polymer with Mn = 12 kDa, with this polymer exhibiting a very high molecular weight dependence since a fraction of Mn = 10 kDa for n-octyl side chains and 0.01 cm2/Vs had a charge carrier mobility of four orders of magnitude lower in optimal processing conditions; for a polymer with 2EH side chains, the mobility improved to 0.1 with a Mn of 18 kDa, which was consistent with molecular weight dependence of mobility (Beaujuge et al., 2012). A larger molecular weight polymer was synthesized of 26 kDa, but an OFET for this fraction was not reported likely due to difficulties in obtaining a homogenous film due to poor solubility of the polymer; however, this polymer did report the highest PCE in OPV morphology for any of the systems, suggesting it has the best charge transport and morphology as a bulk heterojunction system. Aggregation in solution for Si-CPDT-BT was observed (Schulz et al., 2017) for a polymer with Mn of 11 kDa for which the authors simulated a maximum charge carrier mobility of 0.18 cm2/Vs. The authors were unable to conventionally spin coat a thin film at room temperature due to polymer aggregation and very low solubility. Measurements of the films showed slightly shorter π-π diffractions for the Si-CPDT-BT derivative of 3.48 Å vs. 3.75 Å for conventional CPDT-BT with 2EH side chains and 3.70 Å for F-CPDT-BT. Further theoretical calculations showed a higher backbone torsion of Si-CPDT-BT of 19˚ for Si-CPDT-BT vs. 11˚ for CPDT-BT, hindering 1D charge transport, and thus lowering overall mobility. AFM analysis of these films suggests a more amorphous morphology with domains of crystallites as opposed to the fibrillar structure of CPDT-BT and thus increased charge carrier trapping and hopping and hence lower mobility. The authors were only able to achieve a modest improvement in charge mobility values by the alignment of polymer chains by high temperature-rubbing experiments or by annealing at the melting point of the polymer. In the study, there was a modestly higher charge transport for Si-CPDT-BT vs. the CPDT-BT control polymer, which was probably the result of

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poorer film morphology and packing with 2EH side chains compared to the well-known high mobility polymer C16-CPDT-BT. A copolymer of silicon-bridged CPDT (SiCPDT) with 2EH side chains with a fully fluorinated BT monomer was synthesized as a trimer and polymerized with bis-stannyl-thiophene (CPDSi-BT-T), which yielded a polymer with a moderate Mn of 18 kDa (Nketia-Yawson et al., 2015). DFT calculations showed a coplanar backbone with a wavy backbone structure, which correlated with GIWAXs experiment. The authors therefore suggested that the polymer had a packing structure of alternating up down orientations in the π-stacking direction which was characterized by a relatively short distance of 3.42 Å. The polymer exhibited a reliable hole mobility of around 2.82 cm2/Vs with a P(VDFTrFE-CTFE) dielectric that yielded the most ideal OFET behavior. The measured mobility was 9.05 cm2/ Vs for a P(VDF–TrFE) dielectric. In an effort to improve planarity and packing, a number of variations on the CPDT backbone have been explored including germanium-bridged CPDT units, which showed a hole mobility of 0.11 cm2/Vs with an Mn of 15 kDa with charge transport likely being molecular weight limited (Fei et al., 2011). A selenophene analogue of CPDT-BT (CPDSe-BT) has also been synthesized (Fei et al., 2017), which showed a much lower mobility of 0.14 cm2/Vs despite a good molecular weight of Mn = 75 kDa. Extension of the conjugation of the CPDT unit with thienothiophene motifs to form a cyclopentadithienothiophene system (CPDTT-BT) in order to extend the planar ladder-type structure resulted in a maximum charge mobility of 0.67 cm2/Vs with Mn = 41 kDa; although, there was significant deviation from the ideal transistor behavior observed in most CPDT systems (Zhong et al., 2015). AFM measurements showed a significantly less continuous morphology of distinct crystallites with many domain boundaries and low molecular organization, which likely hindered higher charge transport by formation of a significant number of traps and lower interchain hopping. Wide angle with π-stacking distance was also larger for this polymer and more complex with a number of peaks observed at 3.5–3.9 Å. As with many attempts to fluorinate the BT CPDT-type polymer backbones, the authors found in this polymer that there was a significant decrease in molecular weight and a slight decrease in maximum hole mobility to 0.37 cm2/Vs (Albrecht et al., 2012). The polymer showed pronounced ambipolarity with an electron mobility of 0.17 cm2/Vs due to the lowered LUMO combined with the possibility of formation of continuous films. 9.2.2.2 IDT-Based Systems One of the most successful strategies to improve charge carrier mobility in conducting polymers has been to increase planarity of the polymer backbone by utilizing larger fused ring systems, which enhance π-stacking and decrease molecular organization energy, which facilities more effective interchain charge hopping and promotes fewer grain boundaries and, hence, fewer charge traps. Of the fused ring systems, one of the more successful has been the IDT monomer. When initially reported in 2011 that introduction of a stoichiometric dioctyl-fluorene monomer to a poly-triarylamine (PTAA) polymer had a fivefold increase in average charge carrier mobility and that a dioctyl-indenofluorene monomer had an order of magnitude increase in charge carrier mobility to 0.04 cm2/Vs over a control unmodified PTAA, which exhibited a mobility of 0.004 cm2/Vs with no loss in charge transport over a period of months with near ideal transistor transfer curves, it was suggested that this increase was as a result of improved interchain π-orbital overlap (Zhang et al., 2009). Having demonstrated the improvement in charge carrier mobility by the indenofluorene monomer, the same group synthesized the IDT monomer and copolymerized it with both 2,1,3-BT and thieno[3-2,b]bithiophene (TT) (Zhang et al., 2010). The IDT-BT polymer showed a considerably improved mobility of 1.0 cm2/Vs, and the IDT-TT showed a charge mobility of 0.15 cm2/ Vs. Although both polymers showed improved charge mobility over PTAA polymers, they both displayed some non-ideal behavior in the saturation regime attributed to electrode fermi-level and polymer HOMO-level mismatch by the authors. The IDT-BT showed semi-crystalline behavior in XRD and a diffuse π-stacking peak at 4.1 Å with an edge-on packing orientation, whereas IDT-TT proved to be amorphous with a slightly shorter π-stacking at 4.0 Å. The BT comonomer seems to be structurally critical to successful polymers containing the IDT-type monomer due to its ability to adopt a conformation where

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Conjugated Polymers: Perspective, Theory, and New Materials

aromatic protons are syn to the aromatic protons on IDT. DFT confirmed lower torsional disorder in the backbone of IDT-BT vs. IDT-TT. It is interesting that large charge carrier mobilities were observed for the polymers despite the longer π-stacking distance. DPP polymers often exhibit π-stacking of around 3.5 Å, a lack of very high crystallinity, and edge-on packing mode. Bronstein et al. (2011b) explored the effect on field effect mobility of side chain length in IDT-BT polymers. They found that the longest unbranched C16 alkyl chain yielded the highest field effect mobility of 1.2 cm2/Vs for C16-IDT-BT and a reasonable value of 0.57 cm2/Vs for the 2EH side chain. They also found a much lower mobility for the 1-octyl (C8) side chain, probably as a result of high Mw and large PDI affecting film formation and, similarly, a low mobility for the shortest side chains with the lowest molecular weight. It is clear in many of the studies that it is often difficult to decouple effects of PDI and molecular weight when on charge transport while modifying side chains. Care should be taken to compare molecular weights of side chains and molecular weight (Tsao et al., 2011). The IDT-BT microstructure was studied in detail (Zhang et al., 2013) on a polymer with high molecular weight and narrow polydispersity, which yielded record-high charge mobility in both saturation and linear OFET regimes of 2.0 cm2/Vs and 3.6 cm2/Vs, respectively. Thin film microstructure was studied by polarizing spectroscopies, VASE, NEXAFS, and GIXD. Measurements of average backbone orientation by VASE showed low crystallinity and much less backbone alignment extracted from backbone dichroic ratios in IDT-BT compared with crystalline polymers such as PBTTT (DeLongchamp et al., 2007); although it still shows a large proportion of alignment with only a small percentage of polymer chains unaligned with respect to substrate plane. Line shape analysis also suggested a large portion of highly rigid chains in both alignments, even in amorphous domains. NEXAFS similarly showed a broadly well-aligned polymer surface with nearly a 60:40 ratio of surface face-on aligned chains to edge-on at the surface, which was comparable to chemically derived graphene. This value is important as the first few nanometers of surface probed by NEXAFS are those that constitute the active region in a BGTC transistor used in this study. GIXD data confirmed previous results for C16-IDT-BT, showing a broadly amorphous polymer microstructure and correlation with polarized spectroscopies with broad agreement that much of the polymer was orientated similarly, irrespective of whether the domain was crystalline or amorphous. They therefore suggested that less crystalline polymers show pseudo-1D charge transport mechanisms along polymer chains with only small areas of crystallinity necessary for high charge transport. It should be noted that IDT-BT has high solubility in a range of chlorinated and nonchlorinated solvents, suggesting some degree of conformational freedom in the solution phase, which is unusual for such a rigid polymer. Energetic disorder in IDT-BT has been measured via correlation of temperature with gate voltage and found to have almost disorder-free metal-oxide-like charge transport (Venkateshvaran et al., 2014) with mobility independent of gate voltage across the entire temperature range, as well as much reduced charge trapping in IDT-BT vs. PBTTT and PSeDPPBT (also measured by the authors). Urbach energy was also calculated for a number of polymer systems and was found to be lower for IDT-BT than for most other systems measured. Improvements on the IDT-BT system have been elusive despite a number of modifications of backbone with thiophene- or thiazole-flanked BT monomers (Planells et al., 2014) or cyanation and fluorination on thiophene-flanked BT monomers (Casey et al., 2015), which reduced the backbone or planarity of the polymer and, hence, lowered charge mobility to