Conjugated polymers. Properties, processing, and applications [Second edition] 9781138065703, 1138065706

1,181 189 142MB

English Pages [843] Year 2019

Report DMCA / Copyright


Polecaj historie

Conjugated polymers. Properties, processing, and applications [Second edition]
 9781138065703, 1138065706

Table of contents :
Content: Properties and Characterization of Conjugated Polymers. Photophysics. Conducting Polymers. Transport in Conjugated Polymers. Thermoelectrics. Electrochemistry of Conducting Polymers. Electrochromic Polymers. Mechanical Properties of Conjugated Polymers. Photorefractive Polymers. Optoelectronic Polymers. Processing and Morphology of Conjugated Polymers. Printing. Thermal Processing. Morphology Evolution. Conducting Polymer Composites. Oligomer Based Film Structure. Thin Film Structure. Soft X-ray Structural Characterization (RoXS). OFET Applications. Polymers for Organic Photovoltaics. Organic Photovoltaics Architecture. Organic Photovoltaics. Applications of Conjugated Polymers. Conjugated Polymer Sensors. Electrochromic Polymer Devices. Electronic Skin. Electrochemical Devices. Bio- Applications. Redox Active Polymers. Supercapacitors. Actuators and Artificial Muscles. Industrial Status of PEDVT. Aerospace Applications.

Citation preview

Conjugated Polymers Properties, Processing, and Applications

Conjugated Polymers Properties, Processing, and Applications

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: 978-1-138-06570-3 (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 ( 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  and the CRC Press Web site at

Contents Editors.. .....................................................................................................................vii Contributors.............................................................................................................. ix


Conjugated Polymer-Based OFET Devices........................................................1


Electrical Doping of Organic Semiconductors with Molecular Oxidants and Reductants..................................................................................................21

Mark Nikolka and Henning Sirringhaus

Stephen Barlow, Seth R. Marder, Xin Lin, Fengyu Zhang, and Antoine Kahn


Electric Transport Properties in PEDOT Thin Films. . ................................... 45


Thermoelectric Properties of Conjugated Polymers. . ....................................129


Electrochemistry of Conducting Polymers.................................................... 161


Electrochromism in Conjugated Polymers –  Strategies for Complete and Straightforward Color Control................................................................ 201

Nara Kim, Ioannis Petsagkourakis, Shangzhi Chen, Magnus Berggren, Xavier Crispin, Magnus P. Jonsson, and Igor Zozoulenko Kelly A. Peterson, Eunhee Lim, and Michael L. Chabinyc P. Audebert and F. Miomandre

Anna M. Ö sterholm, D. Eric Shen, and John R. Reynolds


Mechanical Properties of Semiconducting Polymers................................... 249


Magnetic Field Effects in Organic Semiconductors; Low and High Fields, Steady State and Time Resolved. . ....................................................... 277

Mohammad A. Alkhadra, Andrew T. Kleinschmidt, Samuel E. Root, Daniel Rodriquez, Adam D. Printz, Suchol Savagatrup, and Darren J. Lipomi

Eitan Ehrenfreund and Z. Valy Vardeny


Organic Electro-Optic Materials................................................................... 299


Establishing the Thermal Phase Behavior and its Inf luence on Optoelectronic Properties of Semiconducting Polymers. . ............................ 329

Larry Dalton

Natalie Stingelin





Poly(3-alkylthiophenes): Controlled Manipulation of Microstructure and its Impact on Charge Transport.............................................................. 351 Michael McBride, Guoyan Zhang, Martha Grover, and Elsa Reichmanis


Microstructural Characterization of Conjugated Organic Semiconductors by X-Ray Scattering . . ............................................................ 391 Maged Abdelsamie and Michael F. Toney


Soft X-Ray Scattering Characterization of Polymer Semiconductors.......... 427


Morphology Evolution and Interfacial Design of Conjugated PolymerBased Photovoltaics. . ...................................................................................... 459

Long Ye, Samuel J. Stuard, and Harald Ade

Yao Liu and Thomas P. Russell


The Relevance of Solubility and Miscibility for the Performance of Organic Solar Cells........................................................................................ 485 Stefan Langner, Jose Dario Perea Ospina, Chaohong Zhang, Ning Li, and Christoph J. Brabec


Processing-Structure-Function Relationships of Polymer-AcidTemplated Conducting Polymers for Solid-State Devices............................. 515 Melda Sezen-Edmonds and Yueh-Lin Loo


Conjugated Polymer Thin Films for Stretchable Electronics........................535


Conducting Polymers for Electrochemical Capacitors.................................. 561


Redox-Active Polymers as an Organic Energy Storage Material. . ................ 587


Electrochromics: Processing of Conjugated Polymers and Device Fabrication on Semi-Rigid, Flexible, and Stretchable Substrates................. 595

Aristide Gumyusenge, William McNutt, and Jianguo Mei

Luciano M. Santino, Yang Lu, Yifan Diao, Hongmin Wang, and Julio M. D’Arcy Kenichi Oyaizu and Hiroyuki Nishide

Matthew Baczkowski, Sneh Sinha, Mengfang Li, and Gregory Sotzing


Separation Techniques Using Conjugated Polymers.. ................................... 629


Organic Bioelectronics Based on Mixed Ion–Electron Conductors. . ........... 679


Conducting and Conjugated Polymers for Biosensing Applications.. .......... 697


Conjugated Poly/­O ligo-Electrolytes for Cancer Diagnosis and Therapy. . ....743


Biomedical Applications of Organic Conducting Polymers......................... 783

Cheng-Wei Lin, Wai H. Mak, Brian T. McVerry, and Richard B. Kaner

Magnus Berggren, Erik O. Gabrielsson, Daniel T. Simon, and Klas Tybrandt C. Pitsalidis, A.M. Pappa, C.M. Moysidou, D. Iandolo, and R.M. Owens Lingyun Zhou, Guillermo C. Bazan, and Shu Wang

Alexander R. Harris, Paul J. Molino, Caiyun Wang, Gordon G. Wallace, and Zhilian Yue

Index. . ...................................................................................................................... 813

Editors 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 where he is 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 to Gallipolis, Ohio at a young age, 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 in fields as diverse as advanced lithium-sulfur batteries, MEMS devices, photovoltaic cells, and biosensors, through several startups. 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.


Contributors Maged Abdelsamie Stanford Synchrotron Radiation Lightsource (SSRL), SLAC National Accelerator Laboratory Stanford University Menlo Park, California Harald Ade Organic and Carbon Electronics Lab (ORaCEL) Department of Physics North Carolina State University Raleigh, North Carolina Mohammad A. Alkhadra Department of NanoEngineering University of California, San Diego San Diego, California Pierre Audebert PPSM – CNRS – Ecole Normale Supérieure Paris-Saclay Paris, France Matthew Baczkowski Polymer Program University of Connecticut Storrs, Connecticut Stephen Barlow School of Chemistry and Biochemistry Center for Organic Photonics and Electronics Georgia Institute of Technology Atlanta, Georgia Guillermo C. Bazan Departments of Chemistry & Biochemistry and Materials, Center for Polymers and Organic Solids University of California Santa Barbara, California

Magnus Berggren Laboratory of Organic Electronics, Department of Science and Technology Linköping University Norrköping, Sweden Christoph J. Brabec Institute of Materials for Electronics and Energy Technology (iMEET) Friedrich-Alexander University Erlangen-Nürnberg Erlangen, Germany and Forschungszentrum Jülich GmbH Helmholtz-Institut Erlangen-Nürnberg (HI ERN) Egerlandstrasse Erlangen, Germany Shangzhi Chen Laboratory of Organic Electronics Department of Science and Technology Linköping University Norrköping, Sweden Xavier Crispin Laboratory of Organic Electronics Department of Science and Technology Linköping University Norrköping, Sweden Michael L. Chabinyc Materials Department University of California Santa Barbara Santa Barbara, California ix



Julio M. D’Arcy Department of Chemistry Institute of Materials Science & Engineering Washington University St. Louis, Missouri

Magnus P. Jonsson Laboratory of Organic Electronics Department of Science and Technology Linköping University Norrköping, Sweden

Larry Dalton Department of Chemistry University of Washington Seattle, Washington

Antoine Kahn Department of Electrical Engineering Princeton University Princeton, New Jersey

Yifan Diao Institute of Materials Science & Engineering Washington University St. Louis, Missouri

Richard B. Kaner Department of Chemistry and Biochemistry University of California, Los Angeles Los Angeles, California

Eitan Ehrenfreund Physics Department and Solid State Institute Technion Israel Institute of Technology Haifa, Israel Erik O. Gabrielsson Laboratory of Organic Electronics, Department of Science and Technology Linköping University Norrköping, Sweden Martha Grover School of Chemical and Biomolecular Engineering Georgia Institute of Technology Atlanta, Georgia Aristide Gumyusenge Department of Chemistry Purdue University West Lafayette, Indiana Alexander R. Harris Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia Donata Iandolo Department of Chemical Engineering and Biotechnology, University of Cambridge Cambridge, United Kingdom

Nara Kim Laboratory of Organic Electronics Department of Science and Technology Linköping University Norrköping, Sweden Andrew T. Kleinschmidt Department of NanoEngineering University of California, San Diego San Diego, California Stefan Langner Institute of Materials for Electronics and Energy Technology (iMEET) Friedrich-Alexander University Erlangen-Nürnberg Erlangen, Germany Mengfang Li Polymer Program University of Connecticut Storrs, Connecticut Ning Li Institute of Materials for Electronics and Energy Technology (iMEET) Friedrich-Alexander University Erlangen-Nürnberg Erlangen, Germany Eunhee Lim Materials Department University of California Santa Barbara Santa Barbara, California



Cheng-Wei Lin Department of Chemistry and Biochemistry University of California, Los Angeles Los Angeles, California Xin Lin Department of Electrical Engineering Princeton University Princeton, New Jersey Darren J. Lipomi Department of NanoEngineering University of California, San Diego San Diego, California Yao Liu Beijing Advanced Innovation Center for Soft Matter Science and Engineering Beijing University of Chemical Technology Beijing, China and Department of Polymer Science and Engineering University of Massachusetts Amherst, Massachusetts Yueh-Lin Loo Department of Chemical and Biological Engineering Princeton University and Andlinger Center for Energy and the Environment Princeton University Princeton, New Jersey Yang Lu Institute of Materials Science & Engineering Washington University St. Louis, Missouri Wai H. Mak Department of Chemistry and Biochemistry University of California, Los Angeles Los Angeles, California

Seth R. Marder School of Chemistry and Biochemistry School of Materials Science and Engineering Center for Organic Photonics and Electronics Georgia Institute of Technology Atlanta, Georgia Michael McBride School of Chemical and Biomolecular Engineering Georgia Institute of Technology Atlanta, Georgia William McNutt Department of Chemistry Purdue University West Lafayette, Indiana Brian T. McVerry Department of Chemistry and Biochemistry University of California, Los Angeles Los Angeles, California Jianguo Mei Department of Chemistry Purdue University West Lafayette, Indiana Fabien Miomandre PPSM – CNRS – Ecole Normale Supérieure Paris-Saclay Paris, France Paul J. Molino Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia Chrysanthi M. Moysidou Department of Chemical Engineering and Biotechnology, University of Cambridge Cambridge, United Kingdom Hiroyuki Nishide Department of Applied Chemistry Waseda University Tokyo, Japan



Mark Nikolka Cavendish Laboratory University of Cambridge Cambridge, United Kingdom

Adam D. Printz Department of NanoEngineering University of California, San Diego San Diego, California

Anna M. Österholm School of Chemistry and Biochemistry Georgia Institute of Technology Atlanta, Georgia

Elsa Reichmanis School of Chemical and Biomolecular Engineering and School of Chemistry and Biochemistry

Kenichi Oyaizu Department of Applied Chemistry Waseda University Tokyo, Japan

and School of Materials Science and Engineering Georgia Institute of Technology Atlanta, Georgia

Roisin M. Owens Department of Chemical Engineering and Biotechnology, University of Cambridge Cambridge, United Kingdom

John R. Reynolds School of Chemistry and Biochemistry, School of Materials Science and Engineering Georgia Institute of Technology Atlanta, Georgia

Anna-Maria Pappa Department of Chemical Engineering and Biotechnology, University of Cambridge Cambridge, United Kingdom

Daniel Rodriquez Department of NanoEngineering University of California, San Diego San Diego, California

Jose Dario Perea Ospina Institute of Materials for Electronics and Energy Technology (iMEET) Friedrich-Alexander University Erlangen-Nürnberg Erlangen, Germany Kelly A. Peterson Materials Department University of California Santa Barbara Santa Barbara, California Ioannis Petsagkourakis Laboratory of Organic Electronics Department of Science and Technology Linköping University Norrköping, Sweden Charalampos Pitsalidis Department of Chemical Engineering and Biotechnology, University of Cambridge Cambridge, United Kingdom

Samuel E. Root Department of NanoEngineering University of California, San Diego San Diego, California Thomas P. Russell Department of Polymer Science and Engineering University of Massachusetts Amherst, Massachusetts 01003, United States Beijing Advanced Innovation Center for Soft Matter Science and Engineering Beijing University of Chemical Technology Beijing, China and Materials Sciences Division Lawrence Berkeley National Laboratory Berkeley, California Luciano M. Santino Department of Chemistry Washington University St. Louis, Missouri



Suchol Savagatrup Department of NanoEngineering University of California, San Diego San Diego, California

Michael F. Toney Stanford Synchrotron Radiation Lightsource (SSRL) SLAC National Accelerator Laboratory Menlo Park, California

Melda Sezen-Edmonds Department of Chemical and Biological Engineering Princeton University Princeton, New Jersey

Klas Tybrandt Laboratory of Organic Electronics, Department of Science and Technology Linköping University Norrköping, Sweden

D. Eric Shen School of Chemistry and Biochemistry Georgia Institute of Technology Atlanta, Georgia Daniel T. Simon Laboratory of Organic Electronics, Department of Science and Technology Linköping University Norrköping, Sweden Sneh Sinha Polymer Program University of Connecticut Storrs, Connecticut Henning Sirringhaus Cavendish Laboratory University of Cambridge Cambridge, United Kingdom Gregory Sotzing Polymer Program University of Connecticut and Department of Chemistry University of Connecticut Storrs, Connecticut Natalie Stingelin School of Materials Science and Engineering and School of Chemistry and Biochemistry Georgia Institute of Technology Atlanta, Georgia Samuel J. Stuard Organic and Carbon Electronics Lab (ORaCEL) Department of Physics North Carolina State University Raleigh, North Carolina

Valy Vardeny Department of Physics & Astronomy University of Utah Salt Lake City, Utah Gordon G. Wallace Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia Caiyun Wang Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia Hongmin Wang Institute of Materials Science & Engineering Washington University St. Louis, Missouri Shu Wang Beijing National Laboratory for Molecular Sciences, Key Laboratory of Organic Solids Institute of Chemistry Chinese Academy of Sciences Beijing, China and University of Chinese Academy of Sciences Beijing, China Long Ye Organic and Carbon Electronics Lab (ORaCEL) Department of Physics North Carolina State University Raleigh, North Carolina Zhilian Yue Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia


Chaohong Zhang Institute of Materials for Electronics and Energy Technology (iMEET) Friedrich-Alexander University Erlangen-Nürnberg Erlangen, Germany Fengyu Zhang Department of Electrical Engineering Princeton University Princeton, New Jersey Guoyan Zhang School of Chemical and Biomolecular Engineering Georgia Institute of Technology Atlanta, Georgia


Lingyun Zhou Beijing National Laboratory for Molecular Sciences, Key Laboratory of Organic Solids Institute of Chemistry Chinese Academy of Sciences Beijing, China and University of Chinese Academy of Sciences Beijing, China Igor Zozoulenko Laboratory of Organic Electronics Department of Science and Technology Linköping University Norrköping, Sweden

1 Conjugated PolymerBased OFET Devices

Mark Nikolka and Henning Sirringhaus

1.1 Introduction........................................................................................... 1 1.2 State of OFET Techn​ology​/Appl​icati​ons/ C​ommer​ciali​zatio​n Efforts................................................................... 3 1.3 Recent Developments in Polymer OFET Materials – From Crystalline Polythiophenes to Donor–Acceptor Polymers............. 3 1.4 Charge Transport in Polymer OFETs.................................................5 1.5 Role of Disorder.....................................................................................7 1.6 Charge Carrier Mobility and Artefacts............................................ 10 1.7 Stability of OFETs................................................................................ 12 1.8 Outlook.................................................................................................. 15 References......................................................................................................... 16

1.1 Introduction Easy processability, mechanical flexibility and the endless possibilities of molecular modifications to achieve, for example, performance improvements or analyte selectivity in sensing applications have staged organic field-effect transistors (OFETs) as one of the future trendsetting technologies. To date, OFET integration in applications ranging from rudimentary sensors and circuits to flexible low-­ resolution displays for electronic paper has been demonstrated. Concomitantly, more and more promising applications for OFETs such as fully flexible organic light-emitting diode (OLED) and organic TFT addressed liquid crystal displays (OLCDs), image sensing applications (e.g. X-Ray sensors), organic logic circuits and sensors for wearable applications come within reach.1–4 These novel applications require the use of high-performance, high-stability organic semiconductors with good uniformities imposing tight constraints on the choice of organic semiconductors. For a long time, small molecular semiconductors were deemed the best materials class to meet these rigorous requirements owing to high-mobility band transport observed in covalently bonded small molecule single crystals5,6; charge transport in such systems is widely thought to approach fundamental limits imposed by the weak van der Waals bonding and the polaronic nature of charge carriers in these molecular solids. Conjugated polymers, on the other hand, were considered too disordered to achieve similarly high charge carrier mobility and for a long time were considered inferior for industrial applications. Nevertheless, whereas small molecular organic semiconductors such as rubrene or pentacene potentially are able to deliver high performances in research laboratories, their high degree of crystallinity, polycrystalline nature and the inevitably resulting device-to-device variations in performance have caused concerns for large area applications, especially in the key displays market where device uniformity is as important as device 1


Conjugated Polymers

performance. Due to their good film forming properties and more uniform thin film microstructures, conjugated polymers could offer potential advantages in order to meet these tight uniformity requirements. Therefore, conjugated polymers that can form percolating networks of chains allowing for both intrachain as well as interchain charge transport have attracted renewed interest in the community. These research efforts have led to a number of breakthroughs that have allowed conjugated polymers to approach performance values that were formerly thought impossible. A lot of research has been conducted to improve the performance of polymers, with field-effect charge carrier mobilities having improved by six orders of magnitude over the past 30 years. In this respect, much of the latest advancements have been made by the class of amorphous/semi-crystalline donor–acceptor polymers. Some of these new donor–acceptor polymers already demonstrated extraordinary charge transport properties, such as hole and electron mobilities exceeding 1 cm2/Vs,7,8 allowing them to match and partially even exceed the performance of amorphous silicon (a-Si) thin film transistor (TFT) technology – a material class that is, for instance, applied in many LCD backplanes nowadays. Nevertheless, despite extensive research, charge transport in these materials is still insufficiently understood, which constrains the design of new high-performance materials and makes it difficult to have clear molecular design criteria for materials that can deliver improved device performance. Furthermore, low operational and air stability of conjugated polymer devices remains a major concern to date, leading to product lifetimes that are too short for demanding large-scale applications. In this chapter, we review the current state of the art as well as address a range of challenges that still have to be overcome. We will particularly focus on the role and origin of environmental and operational stress instabilities, the control of the degree of energetic disorder, as well as the issue of device non-uniformity and non-ideality of device characteristics that can lead to difficulties in estimating accurately reported charge carrier mobility values. Many of the above issues pose a major challenge for the upscaling of polymer OFETs and we will review which steps have to be taken to addressing them (Figure 1.1) 9–11.

FIGURE 1.1  Left: Schematic structure of an active matrix LCD and OLED display representing key applications for polymer OFET devices. Images adapted from Sears, The evolution of display technology; In the case of an LCD display a matrix of TFTs modulates the light emitted from a strong backlight by polarizing liquid crystals; for an OLED display, the TFT array directly addresses the light emitting OLED pixels requiring high currents, performances and stability. Right: Evolution of charge carrier mobilities of semiconducting polymers. The asterisk denotes materials which were aligned using a special technique10 – The quoted mobility value can nevertheless be considered inflated and significantly below 10 cm2/Vs. (With permission from Himmelberger, et al., MRS Commun. 5, 1–13 (2015)).

Conjugated Polymer-Based OFET Devices


1.2 State of OFET Techn​ology​/Appl​icati​ons/ C​ommer​ciali​zatio​n Efforts Over the last decade, organic semiconductors have steadily left the realms of an emerging technology and in many areas have successfully been commercialized. One of the most prominent examples in this respect is the use of OLED displays for smart phones and high-resolution television screens. Here, Samsung has pioneered the development of active-matrix organic light-emitting diode (AMOLED) displays which incorporate OLEDs driven by a backplane of inorganic thin-film transistors. Also, for lighting applications OLEDs have enjoyed popular demand and companies such as Philips, Osram, Novaled, LG Chem and Konica Minolta have successfully brought large area OLED lighting panels to the market. In contrast, OFET technology to date has only been commercialised in certain niche applications. Although applications such as simple logic circuits or chemical and biological sensors have repeatedly been demonstrated in research laboratories, no significant commercialization attempt has, thus far, fully succeeded. The most advanced commercialization efforts to date have been in flexible electronic paper (E-paper) displays by companies including Polymer Vision and Plastic Logic. E-paper does not have demanding drive requirements and at the time these requirements were well matched to the performance that could be delivered by OFETs. However, E-paper is also not capable of delivering vibrant colour and video displays which has limited its market success. Nevertheless, Plastic Logic in Germany has developed a stable, high yielding manufacturing process for OFET-based flexible E-paper displays and is selling these displays for niche applications that benefit from the bistability, low power consumption and excellent readability in bright sunlight conditions, such as electronic readers or signage displays. The recent advancements in the performance of conjugated polymer OFET technology are now allowing it to target more mainstream markets such as high-resolution video rate LCD and OLED displays. For instance, by now the ability to reliably achieve performances above those of amorphous silicon (a-Si) renders polymer OFETs as an attractive alternative to a-Si in state-of-the-art LCD backplane displays. The opportunity to process light weight displays on flexible foil gives OFET driven LCDs an edge over its traditional competitor. The UK based company FlexEnable, which was spun-out of Plastic Logic, has developed and demonstrated a flexible OFET-LCD platform that in price, reliability and yield is comparable to standard amorphous-LCD technology. This offers OFET technology an attractive entry into the large-scale display market12 and enables flexible, light-weight, unbreakable colour and video displays for applications in automotives or consumer electronics. Looking ahead, the industrial “gold-standard” in the display industry is gradually shifting further towards even more demanding applications such as current-driven OLED displays. In this display class, the transistor directly addresses the light emitting pixel and hence imposes an unprecedented requirement on both mobility and stability on the OFETs. Despite these harsh performance requirements, some display manufacturers have already shown prototypes of fully flexible OLED backplane displays based on organic TFT technology. Companies such as Sony, Samsung and LG have all demonstrated fully flexible, full-color OLED displays and independently announced significant investment into this technology. Although not yet driven by polymer OFETs, there is scope and interest for their integration in the near-term future, taking polymer OFETs a step closer to industrialization and full-scale commercialization. However, this will require further improvement in both materials’ performance – ­field-effect mobilities of 10–15 cm2/Vs are required for OLED driving and operational stability.

1.3 Recent Developments in Polymer OFET Materials – From Crystalline Polythiophenes to Donor–Acceptor Polymers The first ever polymer field-effect transistor reported by Koezuka et al. in 1987 was based on electrochemically deposited polythiophene and exhibited a rather moderate field-effect mobility of 2 10−5 cm2/Vs.13 Only one year later, the addition of solubilizing side chains resulted in the first solution processed


Conjugated Polymers

polythiophene transistors exhibiting similarly low performances.14 The reason for choosing polythiophene in both cases was based on a design strategy that had evolved and prevailed in the field for a long time: The design of rigid-rod-like polymers that, very similarly to small molecular semiconductors, are able to form crystalline structures with enhanced π–π stacking and long-range order. Polythiophene with its rigid backbone of conducting conjugated rings was hence deemed the best candidate to achieve similar performances as had been observed for small molecules such as pentacene, rubrene or sexithiophene before.15–20 For almost two decades polythiophenes remained the material of choice with charge carrier mobility values steadily improving to reach values above 0.1 cm2/Vs by the late 1990s. To a significant degree, this achievement was enabled through a better control of thin-film morphology and the control the growth of ordered microcrystalline domains.21,22 This approach of tailoring the material’s crystallinity has led to a range of excellent rigid-rod polymer candidates and brought forth materials such as poly(​2,5-b​is(3-​tetra​decyl​t hiop​hene-​2-yl)​t hien​o[3,2​-b]th​iophe​ne) (pBTTT), which for the first time could exhibit field effect mobilities above 1 cm2/Vs23,24 and hence compete with current amorphous silicon (a-Si) TFT technology that is crucial for commercialization in products such as LCD-backplane displays. The sketch presented in Figure 1.2 shows the prevailing morphology observed for rigid-rod polymer exhibiting clearly defined π–π-stacking of the polymer backbones oriented mostly edge-on with respect to the substrate plane. In the mid-2000s, the new class of donor–acceptor (D–A) polymers emerged which fundamentally deviated from the hitherto prevailing rigid-rod design motive. In contrast to the relatively short electron donating units of polythiophenes, donor–acceptor polymers comprise alternating electron rich donor units and electron poor acceptor units along their backbone. Owing to the presence of both electron rich and electron poor moieties, as well as a relatively small bandgap usually ranging between 1.0 and 1.9 eV, donor–acceptor polymers generally exhibit pronounced ambipolar charge transport which means that they are capable of exhibiting good electron and hole conductance when applying positive and negative gate voltages to the OFET. Although initially substantially poorer performing than polythiophenes, donor–acceptor polymers recently have led to a significant increase in reported charge carrier mobilities, with many reports claiming field-effect mobilities approaching and even exceeding values of 10 cm2/Vs.25–28 Although in many of these reports mobilities are extracted from non-ideal device characteristics which could potentially lead to an overestimation of mobility values (an issue we will address later in this chapter), these performances nevertheless display the potential of D–A type polymers approaching the performance of small molecular semiconductors. This is particularly surprising when considering, that D–A-type polymers are significantly less crystalline than polythiophenes or the formerly mentioned small molecules. In Figure 1.3 some examples of recent high-performance donor–acceptor polymers are shown, along with some more traditional rigid-rod polymers, sometimes termed homopolymers (i.e. having a purely electron donating character). Understanding the origin of

FIGURE 1.2  Film morphology of a crystalline polythiophene polymer (Left) compared to an amorphous/semicrystalline donor–acceptor polymer (Right). Each square represents one monomer unit having chemical structures as represented in Figure 1.3. Side chains are omitted in this sketch.

Conjugated Polymer-Based OFET Devices


FIGURE 1.3  Chemical structure of the common rigid-rod polymers PBTTT, P3HT as well as popular donor– acceptor polymers such as N2200, CDT–BTz, DPP–TT–T and IDT–BT. Acceptor units (e.g. DPP, BT) are circled in light gray and donor units (e.g. IDT, CDT, (oligo)thiophene) are circled in dark gray.

the superior performance of these structurally more disordered polymers over polythiophenes has been a major challenge in the field and more recently the key to opening the door for polymer transistors to be used in highly demanding applications.

1.4 Charge Transport in Polymer OFETs To understand the origin of high-performance in conjugated polymer OFETs, we have to take a look at the mechanism for charge transport in these materials. We should, however, firstly address the elephant in the room and point out that studies of charge transport are naturally entangled with the choice of device architecture and factors such as the device geometry, contact resistance and charge transport interfaces. Some of these factors have successfully been addressed through extensive research effort over the past decades and optimized device architectures have been conceived (Figure 1.4). For instance, the use of fluorinated polymer dielectrics with low dielectric constants such as CYTOP29 or TEFLON30 or conventional inorganic dielectrics such as aluminium oxide (Al2O3), hafnium oxide (HfO),31 benzocyclobutene (BCB) or silicon dioxide (SiO2)1 treated with dipole-shielding self-assembled monolayers (SAMs) based on octadecyltrichlorosilane (OTS),32 hexamethyldisilazane (HMDS) or phosphonic acid (PA)33 have given charge transport interfaces with low interfacial trap densities. Figure 1.5a illustrates the importance of shielding the charge carriers in the accumulation layer of the OFET from the randomly oriented dipole moments in the gate dielectric to avoid a widening of the density of states (DOS) at the crucial charge transport interface; this in turn has been shown to result in a detrimental localization of charge carriers and as a result, a reduction in charge carrier mobility. Evidence that careful optimization of both charge transport interface and the control of processing conditions have now become possible is served by the routine realization of notoriously sensitive n-type charge transport and lightemitting polymer OFETs that are very sensitive to electron trap states at the interface.34–37 A bottleneck in any OFET architecture still remains the injection/extraction of charge carriers into the semiconducting polymer. In contrast to silicon–MOSFETs where carriers are injected from highly doped regions of either n++ or p++ doped silicon, OFETs require a matching of the electrode’s work function to the polymer’s HOMO or LUMO level (Figure 1.5b). For many materials, this remains a challenge as the choice of (stable) electrode metals is limited; the resulting mismatch of energy levels can hence lead to considerable contact resistances. For many current generation donor–acceptor polymers thus far, the HOMO level is positioned between 5.0 and 5.3 eV which allows for relatively efficient injection from gold electrodes facilitating p-type OFETs with sufficiently low contact resistances.


Conjugated Polymers

FIGURE 1.4  Device architecture related factors that can lead to poor OFET performance: a) Presence of dipoles in the dielectric broadens the DOS at the interface resulting in a localization of charges and a reduction in charge carrier mobility. (With permission from Sirringhaus, B. H. Device physics of solution-processed organic fieldeffect transistors. Adv. Mater. 17, 2411–2425 (2005).); b) Work function misalignment between the position of the semiconductor’s LUMO level (or electron affinity, EA)/HOMO level (or ionization potential, IP) and injecting metal electrode’s work function leads contact resistance and non-ideal device performance.

FIGURE 1.5  Schematic representation of an amorphous polymer with the backbone broken up in conjugated segments and charges hopping along these segments (b) Broadening of the HOMO and LUMO levels of a polymer due to energetic and structural disorder. A range of a Gaussian density of states (DOS) (black line) is approximated by an exponential DOS (red line). (With permission from Brondijk, Device Physics of Organic Field-Effect Transistors, PhD thesis (University of Groningen, 2012)).

Nevertheless, despite the above mentioned device specific factors, the biggest challenge still remains the design of polymer semiconductors that are able to transport charges efficiently. In conjugated polymers, variations in the backbone structure (e.g. kinks, defects, variations of torsion angles, etc.) as well as small variations in the intermolecular environment give rise to a range of delocalization lengths experienced by charge carriers along different polymer chains or parts of polymer chains. This results in energy level fluctuations and consequently the polaronic energy levels within a polymer film are inhomogeneously distributed and widened. Additionally, despite there being a few counterexamples, the presence of solubility enhancing side chains generally make it more difficult for the conjugated units to aggregate and to exhibit close π–π stacking. This allows only for weak interactions to take place between isolated chains contrasting the much stronger intermolecular coupling seen in most small molecular systems. Many of these effects are commonly summarized under the term disorder. One of the earliest and simplest models describing charge transport in such disordered materials is the mobility edge (or multiple trapping and release) model developed by Nevill Mott in 1967.38 The model makes the assumption that electronic states are localised below a certain energy level (called mobility edge) and charge transport occurs by thermal excitation of carriers from these trap states to spatially extended, mobile states above the mobility edge. Although being developed originally for the description of amorphous silicon,39 the mobility edge (ME) model had some success in describing charge transport in organic


Conjugated Polymer-Based OFET Devices

semiconductors such as sexithiophene.40 Also for solution processed conjugated polymers the ME model has had some phenomenological success, as for instance demonstrated by charge transport modeling work on the fluorine copolymer PQT-12.41,42 Nevertheless, in the case of polymers, the structural aspects outlined above (mostly) prevent energy band formation, such that charge transport is widely agreed to happen by hopping of charge carriers instead.1 To describe this type of transport, Allen Miller and Elihu Abrahams developed a relatively simple model which predicts the frequency νij of charge carrier hopping between the shallow impurity sites i and j in a lattice with weak coupling and at cryogenic temperature.43 The resulting Miller Abrahams hopping rate is given by

ε j − εi



(− α Rij ) e − νij = ν0e 


for ε j > εi , i.e. upward hop for ε j < εi , i.e. downward hope

where νij is the maximum hopping rate (sometimes this is referred to as hopping attempt frequency) and α is the inverse localization length indicating how well a charge carrier is able to tunnel across the distance R ij. It should be noted here, that the model is elementarily made up of only two independent components: A first quantum mechanical tunneling contribution as well as a second temperature activated component. Consequently, it can be energetically more favourable to hop over a large distance with a small activation energy instead of hopping to a near site with a large activation energy. Since only energetic site differences derived from a density of states (often Gaussian or exponential) are considered, the approach of Miller–Abrahams hopping allows to describe charge transport in disordered systems without the knowledge of complex transfer integrals or reorganization energies; this would be strictly required for more complex models such as the popular Marcus model that also accounts for the polaronic nature of charges inside of organic semiconductors. In fact, in a simplified picture knowing the number of molecular sites as an estimate for the number of hops N needed to travel across a device with distance L, it is possible to calculate the drift velocity v of charge carriers. From this and the applied electric field E, the charge carrier mobility μ = v/E can easily be extracted as a parameter characterizing the efficiency of charge transport in a semiconductor. Due to the relative simplicity of the Miller–Abrahams hopping rate, it has been adapted in various charge transport models termed variable range hopping (VRH) models. Here, for instance, Bässler has pioneered one of the first charge transport models based on hopping in a system with spatial and energetic disorder.44 The Bässler model itself assumes a hopping rate following Miller–Abrahams expression along with a broadening of the HOMO and LUMO levels described by a Gaussian DOS – a shape which is supported by the observation of Gaussian shaped optical spectra.44 Recent adaptations of the Bässler model are the charge transport models developed by Vissenberg–Matters45 as well as Brondijk et al.46, which are similar to the Bässler model47 but make varying assumptions about the shape of the DOS and allow modeling of the current that would be expected in an OFET architecture. In this case, however, there are a number of free fit parameters in the model which make the model parameters difficult to deduce unambiguously and with a high level of confidence.

1.5 Role of Disorder The lack of long-range delocalized transport, as seen for inorganic semiconductors such as silicon, and the presence of hopping transport make conjugated polymers extremely sensitive to the distribution of available states. For a long time, it has hence been the prevailing polymer design motif to mimic the crystalline structure of inorganic materials in an attempt to narrow the distribution of states and even achieve some degree of long-range charge carrier delocalization. The control of morphology and crystallinity has therefore been a fundamentally integral part of any OFET research activity. It is for this reason that the emergence of high-performing D–A copolymers with their structurally disordered


Conjugated Polymers

amorphous to semi-crystalline morphology has been puzzling the community. To understand the superior performance of some of these D–A copolymers, it is essential to consider the two key factors that enable good charge transport properties. A first key factor that has been shown to be important in achieving high carrier mobilities is a sufficiently high long-range order of the polymer chains. One way of achieving such long-range order is a sufficiently high molecular weight that ensures that there are efficient chain interconnections between aggregates/crystallites. Recently several studies were able to shed some light on this aspect by combining a molecular- and microstructural picture of charge transport in polymers. For instance, Salleo and coworkers were able to demonstrate that polymers are fundamentally limited by lattice disorder which can be overcome by charge transport along the extended polymer chains. Therefore, for efficient charge transport it is not necessary that the film is highly crystalline, short-range intermolecular aggregation or nanometre sized crystalline domains are sufficient to allow efficient long-range charge transport as long as these domains are connected by tie-chains.48 The transport through these so-called tie-chains thus enables seemingly structurally disordered conjugated polymers to exhibit high charge carrier mobilities. Figure 1.6 illustrates this concept for the case of a semi-crystalline, weakly ordered and fully amorphous polymer. The tie-chains in the weakly ordered polymer are enabling efficient charge transport through the structurally disordered microstructure; this allows for similar percolation of charges as in the case of a more crystalline polymer. However, although a high molecular weight might constitute a means to achieving long-range order in various polymer semiconductors, it does not necessarily result in a good transistor material. In some cases, long-range ordering can only be achieved through additional means that can induce backbone elongation and inter-grain connectivity. For instance the combination of pre-aggregation in solution, as well as shear-alignment, has been shown to result in significantly improved charge carrier mobilities of 2 cm2/Vs for the polymer N2200.49 Also, in the case of the more crystalline polymers P3HT and PBTTT, the use of pre-aggregation and/or external alignment has shown to be highly successful route towards enhanced charge carrier mobilities.50,51 However, because of this long-range alignment of polymer backbones or, equivalently, aggregation, pronounced device-scale mobility anisotropy will often be inevitable. As a result, generally the charge carrier mobility along the alignment direction of the polymer backbone will be enhanced over the performance observed in the isotropic case. The charge carrier mobility perpendicular to the alignment direction, on the other hand, can be substantially reduced.52 A second key factor has been identified to be a low degree of energetic disorder in the polymer film. Although achieving long-range order constitutes an important ingredient for high performances, it does not fully explain why recent generation D–A copolymers generally tend to outperform crystalline polythiophenes or polymers with high long-range order by up to an order of magnitude. For this, it is important to understand that individual polymer chains themselves possess a strong local mobility anisotropy with varying barriers encountered for intra- and interchain charge transport. The energetic disorder exhibited by a polymer is a representation of this disorder and the resulting distribution of accessible states. In simple terms, it therefore reflects the energetic landscape encountered by a charge when moving

FIGURE 1.6  Microstructure of conjugated polymer films. Representative of microstructure of (a) semi-­ crystalline, (b) weakly ordered, and (c) amorphous polymer films. (With permission from Noriega et al., Nat. Mater. 12, 1038–1044 (2013)).

Conjugated Polymer-Based OFET Devices


through the polymer. Assessing this disorder, unfortunately, is not straightforward, making it difficult to correlate polymer structure to charge transport properties. However, recent studies using a combination of techniques such as photothermal deflection spectroscopy (PDS), thermoelectric Seebeck measurements53 as well as charge transport measurements, in combination with charge transport models54–56, have demonstrated that the degree of energetic disorder is exceptionally low in these high-mobility D–A copolymers. In this way, it has been possible to explore the origin of high performance in some donor–acceptor polymers to an extraordinarily planar, torsion-free polymer backbone. The resulting low energetic disorder essentially ensures that the charge carriers can move efficiently along the tie chains in between aggregates and are not impeded in their motion by chain conformations that constitute transport bottlenecks along the polymer chain. Figure 1.7 shows the planar backbone of the polymer IDT–BT compared to the somewhat twisted backbone of the rigid-rod polymer pBTTT and how it materializes itself in an unusually low disorder that is approaching a disorder-free domain; this is retained even once the polymer is in a completely amorphous phase (Figure 1.7c). Such a high degree of backbone planarity can be detected by pressure dependent Raman measurement (Figure 1.7b)57 or by calculating the gas-phase torsion potentials of IDT–BT (Figure 1.7d). As a consequence of this low disorder

FIGURE 1.7  Resilience of torsion-free polymer backbone conformation to side-chain disorder. a) Simulations of the backbone conformation of IDT–BT and PBTTT in side-chain-disordered and non-interdigitated structures. The side chains and hydrogen atoms are omitted for clarity. Yellow, sulphur atoms; blue, nitrogen atoms; b) Pressure dependence of the intensity ratio of the Raman transitions at 1,542 cm−1 and 1,613 cm−1 (top) and the Raman spectrum of IDT–BT measured using a diamond-anvil cell (bottom). a.u., arbitrary units; c) Simulation of the backbone conformation of IDT–BT in the amorphous phase. A single chain from the simulated unit cell has been highlighted in bright yellow (other colours as in a). d) Calculated gas-phase torsion potentials of IDT–BT and PBTTT. For PBTTT, the potential for torsion between the thiophene and thienothiophene units is shown. (With permission from Venkateshvaran et al., Nature 515, 384–388 (2014)).


Conjugated Polymers

backbone configuration, charge carriers, though still based on hopping, can access almost all available states yielding the observed high charge carrier mobilities reliably above 1 cm2/Vs. Indeed, the feature of low energetic disorder can be found in all families of high-performing donor–acceptor polymers explaining why these systems can excel other materials with superior crystallinity. Therefore, polymer structures with only occasional close contact between individual chains but low energetic disorder and planar backbone have evolved as a novel design motif for the next generation of high-performance polymer semiconductors.

1.6 Charge Carrier Mobility and Artefacts Although it might appear trivial, it remains a fundamental challenge to accurately determine and compare the performance of conjugated polymer OFETs. In this respect, the charge carrier mobility (μ) is naturally the main benchmark to compare the performance of OFETs made from various materials and in different architectures, though it is worth noting that for industrial applications device-dependent metrics, such as absolute ON current or transconductance at a given applied voltage may be more relevant metrics. To extract an OFET’s charge carrier mobility, a classical model developed for ideal silicon MOSFETs is widely applied. This model assumes that the potential drops linearly across the transistor’s channel (also referred to as the gradual channel approximation) and that additionally, no states are present in the semiconductors band gap and hence, mobility is independent of charge carrier concentration. The MOSFET model then gives a simple relation between drain current and gate and drain voltage that allows extraction of charge carrier mobility µ in both the linear and saturation regime according to,

where L W VG VDS VTh

I Dlin = µ

W Ci (VG − VTh )VSD L

I DSat = µ

2 W Ci (VG − VTh ) L


VSD > VG − VTh

is the transistor’s channel length (i.e. the distance between the source and the drain electrode) is the channel width (i.e. the overall width of the source or drain contact) is the applied gate voltage the applied drain-source voltage is the transistor’s threshold voltage

The saturation charge carrier mobility (often quoted as a single reference value due to its importance for high-current applications such as OLED backplanes) then can be extracted by taking the square root’s slope of measured transfer characteristics. Figure 1.8 displays how the saturation mobility can be extracted in an ideal or close to ideal device as for instance observed for many of the low-disorder polymers we discussed previously.53 In such cases, we would expect a low threshold voltage, a sharp subthreshold swing as well as output characteristics that is linear for low voltages and shows flat saturation behaviour as displayed in Figure 1.8. However, strictly speaking, many of the assumptions made in the MOSFET model are invalid for most OFETs as organic semiconductors often do not exhibit the same “ideal” behaviour as for instance seen for silicon transistors. For instance, in contrast to MOSFETs that operate in inversion mode, OFETs are operated in accumulation mode, and hence artefacts such as the formation of backchannels are often observed, leading to output characteristics that do not fully saturate. The broad distribution of states encountered in conjugated polymers, trapping at impurities or the semiconductor–dielectric interface, misalignment between the injecting electrode and the semiconductor’s energy levels are some of the additional factors that can lead to non-ideal FET characteristics and an inflation of extracted charge


Conjugated Polymer-Based OFET Devices


Gate Voltage (VG /V)

Drain current (ID /A)

Sqrt (ID)


Sub-Treshold Slope (V/dec)

Drain Voltage (VD /V)

FIGURE 1.8  Drain current as a function of gate voltage (transfer characteristics, left) and drain current as a function of Source-drain Voltage (Output characteristics right) for an ideal polymer OFET with negligible hysteresis and a small threshold voltage.

carrier mobilities. In the following, we list four more common phenomena that have been observed in polymer devices and are known to lead to an inflation of reported charge carrier mobility values: i) A semiconductor with a broad distribution of states (i.e. large disorder) will tend to exhibit a genuinely field-dependent charge carrier mobility. This originates in the growing number of states that become accessible at higher gate-voltages. In these systems, the square root of the drain current will usually be super-linear at higher fields and consequently, higher fields will result in higher charge carrier mobilities. In such cases, a large discrepancy between the charge carrier mobility extracted in the linear and the saturation regime can be expected. Consequently, the linear mobility should be quoted as a lower bound to the device’s performance. ii) The presence of polarizable dielectrics and the presence of ions in the gate dielectric can cause the value of the gate dielectric capacitance, Ci, to be underestimated, leading to an overestimation of carrier mobility. This may manifest itself as well in a large hysteresis in the measured transfer characteristics. In such circumstances, it is absolutely crucial to determine the steady state capacitance of the gate dielectric. iii) The presence of contact resistance will lead to a shift in threshold voltage. At higher gate voltages this initial resistance is gradually overcome leading to a steeper than usual rise in current and hence inflated charge carrier mobility. This effect usually gives rise to an S-shaped or (if the threshold voltage is large enough) exponential shape of the square root resulting in a severe overestimation of reported charge carrier mobility values (illustrated schematically in the longdashed line curve in Figure 1.9)58; iv) The square root of the drain current exhibits a kink with higher slopes being observed at lower gate voltages (illustrated schematically in the short-dashed line curve in Figure 1.9). Interestingly, this phenomenon occurs mostly for low-bandgap donor–acceptor polymer devices with a bottom gate architecture comprising SAMs such as OTS.26 Various explanations have been given for the origin of this phenomenon, relating it to contact resistance or to an unconfined accumulation layer at lower gate voltages.59,60 Although the exact origin for inflated charge carrier mobility will have to be established in each individual case (and in some cases a combination of mechanisms will simply make it impossible to pinpoint individual causes), it is important to use a robust mobility extraction method that allows a conservative comparison of the true current-carrying capability of different transistor materials and structures. One example of addressing the inflation of charge carrier mobilities has recently been proposed by Choi et al. who define a reliability factor for mobility extraction in non-ideal devices.61 The method is based on comparing the measured on-current at the highest gate voltage to that of an ideal transistor and the determination of the effective mobility that such an ideal transistor would need to have to match the


Conjugated Polymers

FIGURE 1.9  Illustration of mobility inflation on the basis of the square root of ID.

measured on-current. The reliability factor, which relates the effective mobility to the claimed maximum mobility extracted from the measured transfer characteristics, can be used to avoid the reporting of inflated charge carrier mobilities and lead to a more robust reporting of device performances. Likewise, it is generally the case that the above artefacts affect the extracted saturation mobility more than the linear charge carrier mobility. This is also the reason why many conjugated polymer OFETs will exhibit a saturation mobility that can be up to an order of magnitude higher than the reported linear mobility. Hence, the ratio between these two mobilities should also give a good estimate of how accurate and representative the extracted charge carrier mobilities are. This is why in an industrial setting, the on-current of OFETs with a given architecture is a much more applicable figure of merit and is used to assess performance. Nevertheless, although mobility overestimation is a current problem especially for values exceeding 10 cm2/Vs there are many examples of reported devices that exhibit close-toideal linear and saturation mobilities well above 1 cm2/Vs; this approaches industrial requirements for OLCD and OLED backplane displays as long as a sufficiently high stability can be warranted.

1.7 Stability of OFETs As discussed above, the rapid improvement seen in OFET performance, in particular enabled by novel D–A copolymers, in many areas has brought OFETs significantly closer towards real-world applications. It is for this reason that recently their reliability under representative operational and environmental conditions has been called into closer focus. We would, therefore, like to continue this discussion from an application point of view and take a look not only at the performance of polymer OFETs but also their stability. In fact, for industrial applications, the long-term stability under shelf-storage as well as bias–stress will crucially determine the suitability of OFETs for commercialization. In this respect, the disordered structure of most conjugated polymers is not only reflected in complex charge transport mechanisms but often in unsatisfactory levels of stability. Since organic semiconductors are only weakly held together by van der Waals interactions, they are prone to degradation and potential structural changes during operation and ingress by environmental species. In this context, it has been understood for some time that electrons in high-lying energy levels, i.e. materials with low electron affinity, exhibit a high susceptibility to electron trapping in the presence of water and/or oxygen species in ambient conditions62, which for a long time has been a major hurdle for the realization of n-type OFETs. In the current generation of high-performance p-type and ambipolar OFETs this issue is addressed by

Conjugated Polymer-Based OFET Devices


designing molecules with deep-lying HOMO and LUMO levels. Generally, ionization potentials deeper than approx. 5.1 eV are considered to improve stability against oxidative p-type doping by atmospheric species, such as oxygen, significantly. While this improves the stability of conjugated polymers in oxygen rich environments, these materials have been found to show a tendency to degrade in oxygen poor environments such as nitrogen or vacuum.63 Here, it is important to point out that degradation due to environmental species is generally creating both shallow and deep traps for charge carriers. Shallow trapping happens on a fast time scale and thus it can be easily observed during single transfer or output measurements. Deep traps manifest themselves in long-term threshold voltage shifts during environmental exposure or operational stress tests. Recently, studies have shown that in current generation D–A copolymers, many of these traps are related to water molecules and can be passivated with molecular additives that are incorporated into the film to occupy voids within the films and passivate/displace water molecules away from the charge transporting polymer backbone.63 In this way significant improvements in the environmental and operational stability can be achieved, but also improvements of other device parameters such as contact resistance, that is affected by charge trapping when charges travel from the contacts through the bulk of the polymer to the active interface. One mechanism by which water is able to create shallow trap states could be related to induced torsion along the polymer backbone and thus, an increased disorder; this mechanism renders high-performing donor–acceptor polymers with a narrow distribution of states particularly susceptible. For applications that require prolonged operation, such as in active-matrix addressed light emitting diode displays, a high level of threshold voltage stability is furthermore of paramount importance. Naturally, the effects of environmental degradation described above apply to prolonged FET operation as well and have to be addressed. However, during bias or current stress, additional trapping mechanisms take place which might not occur during a single transfer or output measurement. When an FET is operated for a prolonged time, charge carriers can continuously get trapped in the dielectric, in localized states at the dielectric/semiconductor interface or in deep-localized states in the organic semiconductor. These trapped charges form an immobile charge density distribution that shields the gate induced field and thus, gives rise to an ongoing threshold voltage shift scaling with the stress duration.64 Since most of the mobile charges are not affected by this degradation mechanism, the trapped carriers often do not lead to a change in the shape of the transfer characteristics or the extracted fieldeffect mobility, as would be the case during shallow trapping. The major culprit of stress instability has nevertheless also been related to water molecules65,66 involving the electrochemical oxidation of residual water at the semiconductor-dielectric interface according to the basic electrochemical reaction

4 H + + O 2 + 4 e − → 2H 2 O

which converts mobile holes on the organic semiconductor to more immobile protons. The protons generated in this reaction are assumed to diffuse into the dielectric until an equilibrium between the surface concentration [os+] of holes in the accumulation layer and the volume concentration [H+] of protons in the dielectric (in the vicinity of the surface) is obtained. Due to its diffusive nature, the resulting stress function has been shown to exhibit the shape of a stretched exponential, and thus scientifically justifies a fit of the time-dependent threshold voltage shift according to

β   t  ∆V = V0  1 − exp  −    τ  

where β represents an empirical exponent τ is the trap relaxation time The above interpretation can describe accurately the stressing behaviour observed for many polymer OFETs as well as explain the slow recovery exhibited by OFETs after stressing.65,66 There have been


Conjugated Polymers

several approaches to address this issue to achieve satisfactory levels of device performance. The use of molecular additives to displace water molecules in the porous polymer films has recently been shown to result in high-performance polymer OFETs that exhibit levels of environmental and threshold stability comparable to the widely used amorphous silicon transistor technology.63 Figure 1.10 summarizes how the incorporation of a small amount of a molecular additive such as the small molecules TCNQ, F4-TCNQ or Aminobenzonitrile (ABN) results in improved environmental OFET stability (Figure 1.10a) as well as improved charge carrier injection. The transistor data serves as an example of how ideal, textbook-like devices with low contact resistance can be obtained that manifest themselves in linear output characteristics at low source-drain voltages (Figure 1.10b); this becomes even more explicit by extrapolation of the contact resistance using transfer line measurements (Figure 1.10d). The beneficial effect of using molecular additives likewise manifests itself in a significantly improved stability under bias-stress conditions reflective of the high-demanding operational conditions in an active-matrix OLED backplane display.

FIGURE 1.10  Improving polymer FET performance and the environmental and operational stability through the use of molecular additives (a) Linear (V DS = –5 V, dashed lines) and saturation (V DS = –50 V, solid lines) transfer characteristics of IDTBT OFETs with (right panel) and without (left panel) 2 wt.% of TCNQ additive. Measurements were taken successively for the as-prepared device, after 24 hours exposure to first air and then nitrogen environments and after a 12 h anneal in nitrogen. The device structure is shown as an inset (channel length L = 20 µm, channel width W = 1 mm); (b) Output characteristics of an OFET with 2 wt.% of TCNQ additive; (c) Electron affinity of the F4TCNQ (top), TCNQ (middle) and ABN (bottom) additives used; (d) Transmission line measurements of the normalized channel resistance as a function of channel length for FETs comprising IDTBT (blue squares), IDTBT after air exposure (black diamonds) and IDTBT with 2 wt.% of TCNQ (green triangles), ABN (magenta triangles) or F4TCNQ (red circles). The contact resistance can be extracted from an extrapolation to zero channel length; (e) Constant current-stress measurements at 2 µA and room temperature comparing the threshold voltage shift of neat IDTBT OFETs with and without additives, in nitrogen. The recovery kinetics after removing the current stress are also shown. (With permission from Nikolka et al., Nat. Mater. 16, 356–362 (2017)).

Conjugated Polymer-Based OFET Devices


A different approach has been the use of layered bilayer dielectrics where threshold shifts have been shown to be compensated to well under 1 V over the course of several hours of operation.67,68 These examples demonstrate OFET with levels of stability that are compatible with current industrial requirements for the key OFET markets of LCD and are approaching those of current-driven active-matrix OLED displays. A further upcoming area of interest is the stability of OFETs for sensing and bioelectronics applications.69–74 Conjugated polymers are highly topical due to their suitability for bio-interfacing as well as an infinite toolkit of potential surface and bulk functionalization routes. However, whereas in an active-matrix driven display the OFETs will be part of an encapsulated stack, for sensing it is desirable to have the polymer semiconductor exposed to the environment that should be sampled. Simultaneously, desired high levels of sensitivity and selectivity require OFETs with unprecedented inherent and underlying stability in challenging environments such as water or biofluids. This requirement stands in stark contrast to the susceptibility of particularly high mobility conjugated polymers to water traps which we have addressed earlier on. Nevertheless, for low-performance systems such as the polymer PTAA sufficient stability for OFET operation in water and even ionic environments have been demonstrated. An impressive demonstration of the sensing capabilities of OFETs has been the detection of heavymetal ions in sea water environments in which said ions will diffuse into the semiconductor and selectively dope and de-dope it.75 An alternative approach is organic bio-electronic transistors demonstrated by Torsi et al., where the detection of odorant binding proteins to a functionalized gate-electrode is enhanced through a water-gated OFET.76 However, due to the direct exposure of the semiconductor to water environments, all these devices operate in a domain between OFETs and organic electrochemical transistors (OECTs) and stability issues from non-analyte related doping, ion migration and charging of the semiconductor need to be addressed. It remains yet to be seen if the current advancements in environmental and operational stability observed for high-performance D–A polymer OFETs can be translated into this challenging field of non OECT-based biological sensors as well.

1.8 Outlook We have shown in this chapter that issues related to the stability and performance of conjugated polymer OFETs have progressively been addressed to a degree that allows for the emergence of first polymer OFET driven products and prototypes. Polymer OFETs have by now shown sufficient performance, stress stability and lifetimes under representative conditions for typical configurations in OLCD displays and are approaching those required for OLED applications. Polymer-OFETs are starting to be a real economic alternative as evident by their emerging use in products such as flexible OLCD display backplanes. However, for conjugated polymer based OFET devices to truly penetrate into the key OLED display market, the challenge remains the development of novel polymer semiconductors that reliably are capable of exceeding “real” linear and saturation charge carrier mobilities above 10 cm2/Vs while maintaining the currently demonstrated high levels of stability. The design of novel D–A type polymers with extended networks of interconnected tie-chains, low torsional backbones and transport that approach a disorder-free domain will hence be of fundamental importance to reach this goal. These achievements combined with novel device architectures, such as vertical OFETs designs77 allowing for ultra-short channel length, are staging polymer OFET as an attractive technology for the next generation of backplane displays. The large investments into OLED technology by companies such as LG and Samsung will generate further traction in this direction with the potential of offering truly flexible all-organic displays in the near- to mid-term future. While in display applications organic semiconductors are used to replace conventional inorganic semiconductors and offer functional benefits such as the compatibility with low-temperature flexible substrates and potentially lower manufacturing costs, for sensing applications OFETs enable truly stretchable, skin-compatible electronics and the wide synthetic tuning of the molecular structures to,


Conjugated Polymers

for example, enhance the sensitivity to particular analytes. OFETs are currently being widely explored for sensors in artificial skin, wearable and implantable electronics. Although currently a lot of focus is put on the development of novel OECT-type sensing architectures, these require the permanent presence and stability of a surrounding ionic medium, potentially posing constraints on the selectivity of these architectures. OFET architectures offer potential advantages to OECT architectures in this respect and OFETs might also find applications in selective gas sensing78 as well as selective sensing in fluids with unstable ionic concentrations such as bio-fluids. However, for commercial applications, the response, reproducibility, as well as reliability of OFET sensing architectures, have to be substantially improved over the current state of the art.


1. Sirringhaus, H. 25th anniversary article: Organic field-effect transistors: The path beyond amorphous silicon. Adv. Mater. 26, 1319–1335 (2014). 2. Xu, J., Wang, S., Wang, G.J.N., Zhu, C., Luo, S., Jin, L., Gu, X., et al. Highly stretchable polymer semiconductor films through the nanoconfinement effect. Science 355, 59–64 (2017). 3. Takeda, Y., Hayasaka, K., Shiwaku, R., Yokosawa, K., and Shiba, T. Fabrication of ultra-thin printed organic TFT CMOS logic circuits optimized for low-voltage wearable sensor applications. Sci. Rep. 6, 25714 (2016). 4. Perinot, A., Kshirsagar, P., Malvindi, M.A., Pompa, P.P., Fiammengo, R., and Caironi, M.. Directwritten polymer field-effect transistors operating at 20MHz. Sci. Rep. 6, 38941 (2016). 5. Troisi, A. Prediction of the absolute charge mobility of molecular semiconductors: The case of rubrene. Adv. Mater. 19, 2000–2004 (2007). 6. Hasegawa, T., and Takeya, J. Organic field-effect transistors using single crystals. Sci. Technol. Adv. Mater. 10, 024314 (2009). 7. Zhang, W., Smith J., Watkins S.E., Gysel R., McGehee M., Salleo A., Kirkpatrick J., et al. Indacenodithiophene semiconducting polymers for high performance air stable transistors. J. Am. Chem. Soc. 132, 11437–11439 (2010). 8. Bronstein, H., Chen Z., Ashraf R.S., Zhang W., Du J., Durrant J.R, Tuladhar P.S., et al. Thien​o[3,2​ -b]th​iophe​ne-di​ketop​y rrol​opyrr​ole-c​ontai​ning polymers for high-performance organic fieldeffect transistors and organic photovoltaic devices. J. Am. Chem. Soc. 133, 3272–3275 (2011). 9. www.s​ears.​com/a​r ticl​es/tv​s-ele​c tron​ics/t​elevi​sions​/the-​evolu​t ion-​of-di​splay​-tech​nolog​y.htm​l. The evolution of display technology. Sears (2018). 10. Tseng, H., Phan H., Luo C., Wang M., Perez L.A., Patel S.N., Ying L., et al. High-mobility fieldeffect transistors fabricated with macroscopic aligned semiconducting polymers. Adv. Mater. 26, 2993–2998 (2014). 11. Himmelberger, S., and Salleo, A. Engineering semiconducting polymers for efficient charge transport. MRS Commun. 5, 1–13 (2015). 12. http:​//www​.flex​enabl​​/tech​nolog​y/olc​d-man​ufact​uring​. FlexEnable Limited (2017). 13. H. Koezuka, Tsumura, A., and Ando, T. Field-effect transistor with polythiophene thin film. Synth. Met. 18, 699–704 (1987). 14. Assadi, A., Svensson, C., Willander, M., and Inganäs, O Field-effect mobility of poly (3-hexylthiophene). Appl. Phys. Lett. 53, 1–4 (1988). 15. Park, S. K., Member, S., Anthony, J. E., and Jackson, T. N. Solution-processed TIPS-pentacene organic thin-film-transistor circuits. IEEE Electron Device Lett. 28, 877–879 (2007). 16. Lin, Y., Gundlach, D. J., Nelson, S. F., and Jackson, T. N. Stacked pentacene layer organic thin-film transistors with improved characteristics. IEEE Electron Device Lett. 18, 606–608 (1997). 17. Street, R. A. The benefit of order. Nat. Mater. 5, 171–172 (2006). 18. Garnier, F., Horowitz, G., Peng, X., and Fichou, D. An all-organic “soft” thin film transistor with very high carrier mobility. Adv. Mater. 2, 592–594 (1990).

Conjugated Polymer-Based OFET Devices


19. Gamier, F., Yassar, A., Hajlaoui, R., Horowitz, G., Deloffre, F., Serve, B., Ries, S., et al. Molecular engineering of organic semiconductors: Design of self-assembly properties in conjugated thiophene oligomers. J. Am. Chem. Soc. 115, 8716–8721 (1993). 20. Williams, W. G. Hole mobility in rubrene. Discuss. Faraday Soc. 51, 61–66 (1971). 21. Sirringhaus, H., Brown, P.J. Friend, R.H., Nielsen, M.M., Bechgaard, K., Langeveld-Voss, B.M.W., Spiering, A.J.H., et al. Two-dimensional charge transport in self-organized, high-mobility conjugated polymers. Nature 401, 685–688 (1999). 22. Wang, G. Swensen, J., Moses, D., and Heeger, A.J. Increased mobility from regioregular poly (3-hexylthiophene) field-effect transistors. J. Appl. Phys. 93, 6137 (2003). 23. McCulloch, I., Heeney M., Bailey C., Genevicius K., MacDonald I., Shkunov M., Sparrowe D., et al. Liquid-crystalline semiconducting polymers with high charge-carrier mobility. Nat. Mater. 5, 328–333 (2006). 24. DeLongchamp, D.M., Kline, R.J., Jung, Y., Germack, DS., Lin, E.K., Moad, A.J., Richter, L.J., et al. Controlling the orientation of terraced nanoscale “ribbons” of a poly(thiophene) semiconductor. ACS Nano 3, 780–787 (2009). 25. Bucella, S.G. Luzio, A., Gann, E., Thomsen, L., Mcneill, C.R., Pace, G., Perinot, A., et al. Macroscopic and high-throughput printing of aligned nanostructured polymer semiconductors for MHz largearea electronics. Nat. Commun. 6, 8394 (2015). 26. Li, J., Zhao, Y., Tan, H.S., Guo, Y., Di, C.A., Yu, G., Liu, Y., et al. A stable solution-processed ­polymer semiconductor with record high-mobility for printed transistors. Sci. Rep. 2, 754 (2012). 27. Kim, G., Kang S.J., Dutta G.K., Han Y.K., Shin T.J., Noh Y.Y., and Yang C A thienoisoindigo-naphthalene polymer with ultrahigh mobility of 14.4cm2/V·s that substantially exceeds benchmark values for amorphous silicon semiconductors. J. Am. Chem. Soc. 136, 9477–9483 (2014). 28. Luo, C., Ko, A., Kyaw, K., Perez, L.A., Patel, S., Wang, M., Grimm, B., and et al. General strategy for self-assembly of highly oriented nanocrystalline semiconducting polymers with high mobility. Nano Lett. 14, 2764-2771 (2014). 29. Pecunia, V., Nikolka, M., Sou, A., Nasrallah, I., Amin, A., McCulloch, I., and Sirringhaus, H.. Trap healing for high-performance low-voltage polymer transistors and solution-based analogue amplifiers on foil. Adv. Mater. 29, 1606938 (2017). 30. Baeg, K., Noh, Y., and Kim, D. Charge transfer and trapping properties in polymer gate dielectrics for non-volatile organic field-effect transistor memory applications. Solid State Electron. 53, 1165–1168 (2009). 31. Acton, O. Ting G., Ma H., Ka J.W., Yip H.L., Tucker N.M., Jen A.K.Y. p-s-phosphonic acid organic monolayer/sol–gel hafnium oxide hybrid dielectrics for low-voltage organic transistors. Adv. Mater. 20, 3697–3701 (2008). 32. Lei, Y., Deng, P., Li, J., Lin, M., Zhu, F., Ng, T.W., Lee, C.S., et al. Solution-processed donor- ­acceptor polymer nanowire network semiconductors for transistors. Sci. Rep. 6, 24476 (2016). 33. Klauk, H., Zschieschang, U., Pflaum, J., and Halik, M. Ultralow-power organic complementary circuits. Nature 445, 745–748 (2007). 34. Chua, L., Zaumseil, J., Chang, J.F., Ou, E.C.W., Ho, P.K.H., Sirringhaus, H., and Friend, R.H. General observation of n-type field-effect behaviour in organic semiconductors. Nature 434, 194–9 (2005). 35. Zaumseil, J., Friend, R. H., and Sirringhaus, H. Spatial control of the recombination zone in an ambipolar light-emitting organic transistor. Nat. Mater. 5, 69–74 (2006). 36. Smits, E.C.P. Setayesh, S., Anthopoulos, T.D., Buechel, M., Nijssen, W., Coehoorn, R., Blom, P.W.M., et al. Near-infrared light-emitting ambipolar organic field-effect transistors. Adv. Mater. 19, 734–738 (2007). 37. Sirringhaus, B. H. Device physics of solution-processed organic field-effect transistors. Adv. Mater. 17, 2411–2425 (2005). 38. Mott, N. F. Electrons in disordered structures. Adv. Phys. 16:61, 49–144 (1967).


Conjugated Polymers

39. Mott, N. The mobility edge since 1967. J. Phys. C Solid State Phys. 20, 3075 (1987). 40. Horowitz, G., Hajlaoui, R., and Delannoy, P. Temperature dependence of the field-effect mobility of sexithiophene. Determination of the density of traps. J. Phys. III Fr. 5, 355–371 (1995). 41. Salleo, A., Chen, T. W., and Völkel, A. R. Intrinsic hole mobility and trapping in a regioregular poly (thiophene). Phys. Rev. B 70, 115311 (2004). 42. Salleo, A. Charge transport in polymeric transistors. Mater. Today 10, 38–45 (2007). 43. Miller, A., and Abrahams E. Impurity conduction at low concentrations. Phys. Rev. 120, 745–755 (1960). 44. Carlo, A. M., Study, S., and Bässler, H. Charge transport in disordered organic photoconductors. Phys. Status Solidi 175, 15 (1993). 45. Vissenberg, M. C. J. M., and Matters, M. Theory of the field-effect mobility in amorphous organic transistors. Phys. Rev. B 57, 13 (1998). 46. Brondijk, J.J., Roelofs, W.S.C., Mathijssen, S.G.J., Shehu, A., Cramer, T., Biscarini, F., Blom, P.W.M, et al. Two-dimensional charge transport in disordered organic semiconductors. Phys. Rev. Lett. 109, 056601 (2012). 47. Brondijk, J. J. Device Physics of Organic Field-Effect Transistors, PhD thesis (University of Groningen, 2012). 48. Noriega, R., Rivnay, J., Vandewal, K., Koch, F.P.V., Stingelin, N., Smith, P., Toney, M.F., et al. A general relationship between disorder, aggregation and charge transport in conjugated polymers. Nat. Mater. 12, 1038–1044 (2013). 49. Wang, G., Huang, W., Eastham, N.D., Fabiano, S., Manley, E.F., Zeng, L., Wang, B., et al. Aggregation control in natural brush-printed conjugated polymer films and implications for enhancing charge transport. Proc. Natl. Acad. Sci. 114, E10066–E10073 (2017). 50. Persson, N. E., Chu, P. H., McBride, M., Grover, M., and Reichmanis, E. Nucleation, growth, and alignment of poly(3-hexylthiophene) nanofibers for high-performance OFETs. Acc. Chem. Res. 50, 932–942 (2017). 51. Biniek, L., Leclerc, N., Heiser, T., Bechara, R., and Brinkmann, M. Large scale alignment and charge transport anisotropy of pBTTT films oriented by high temperature rubbing. Macromolecules 46, 4014–4023 (2013). 52. Wu, D., Kaplan, M., Ro, H.W., Engmann, S., Fischer, D.A., DeLongchamp, D.M., Richter, L.J., et al. Blade coating aligned, high-performance, semiconducting-polymer transistors. Chem. Mater. 30, 1924–1936 (2018). 53. Venkateshvaran, D., Nikolka, M., Sadhanala, A., Lemaur, V., Zelazny, M., Kepa, M., Hurhangee, M., et al. Approaching disorder-free transport in high-mobility conjugated polymers. Nature 515, 384–388 (2014). 54. Kronemeijer, A. J., Pecunia, V., Venkateshvaran, D., Nikolka, M., Sadhanala, A., Moriarty, J., Szumilo, M., et al. Two-dimensional carrier distribution in top-gate polymer field-effect transistors: Correlation between width of density of localized states and urbach energy. Adv. Mater. 26, 728–733 (2014). 55. Nikolka, M. Hurhangee, M., Sadhanala, A.,Chen, H., Mcculloch, I., and Sirringhaus, H. Correlation of disorder and charge transport in a range of indacenodithiophene-based semiconducting polymers. Adv. Electron. Mater. 1700410, 1–7 (2017). 56. Xu, Y. Sun, H., Li, W., Lin, Y.F., Balestra, F., Ghibaudo, G., and Noh, Y.Y. Exploring the charge transport in conjugated polymers. Adv. Mater. 29, 1702729 (2017). 57. Schmidtke, J. P., Kim, J. S., Gierschner, J., Silva, C., and Friend, R. H. Optical spectroscopy of a polyfluorene copolymer at high pressure: Intra- and intermolecular interactions. Phys. Rev. Lett. 99, 6–9 (2007). 58. Liu, C. Li, G., Pietro, R.D., Huang, J., Noh, Y.Y., Liu, X., and Minari, T. Device physics of contact issues for the overestimation and underestimation of carrier mobility in field-effect transistors. Phys. Rev. Appl. 8, 034020 (2017).

Conjugated Polymer-Based OFET Devices


59. Bittle, E. G., Basham, J. I., Jackson, T. N., Jurchescu, O. D., and Gundlach, D. J. Mobility overestimation due to gated contacts in organic field-effect transistors. Nat. Commun. 7, 10908 (2016). 60. Okachi, T., Kashiki, T., and Ohya, K. Device operation mechanism of field-effect transistors with high mobility donor-acceptor polymer semiconductors. Proc. SPIE 9568, 95680I–1 (2015). 61. Choi, H. H., Cho, K., Frisbie, C. D., Sirringhaus, H., and Podzorov, V. Critical assessment of charge mobility extraction in FETs. Nat. Mater. 17, 2–7 (2018). 62. Takimiya, K., Takimiya, K., Yamamoto, T., Ebata, H., and Izawa, T. Design strategy for air-stable organic semiconductors applicable to high-performance field-effect transistors. Sci. Technol. Adv. Mater. 8, 273 (2007). 63. Nikolka, M., Nasrallah, I., Rose, B., Ravva, M.K., Broch, K., Harkin, D., Charmet, J., et al. High operational and environmental stability of high-mobility conjugated polymer field-effect transistors achieved through the use of molecular additives. Nat. Mater. 16, 356–362 (2017). 64. Sirringhaus, H. Reliability of organic field-effect transistors. Adv. Mater. 21, 3859–3873 (2009). 65. Mathijssen, S.G.J., Cölle, M., Gomes, H., Smits, E.C.P., Boer, B.D., McCulloch, I., Bobbert, P.A., et al. Dynamics of threshold voltage shifts in organic and amorphous silicon field-effect transistors. Adv. Mater. 19, 2785–2789 (2007). 66. Bobbert, P. A., Sharma, A., Mathijssen, S. G. J., Kemerink, M., and De Leeuw, D. M. Operational stability of organic field-effect transistors. Adv. Mater. 24, 1146–1158 (2012). 67. Wang, C., Fuentes-Hernandez, C., Yun, M., Singh, A., Dindar, A., Choi, S., Graham, S., et al. Organic field-effect transistors with a bilayer gate dielectric comprising an oxide nanolaminate grown by atomic layer deposition. ACS Appl. Mater. Interfaces 8, 29872-29876 (2016). 68. Jia, X., Fuentes-Hernandez, C., Wang, C., Park, Y., and Kippelen, B. Stable organic thin-film transistors. Sci. Adv. 4, 1–8 (2018). 69. Rivnay, J., Inal, S., Salleo, A., Berggren, M., and Malliaras, G. G. Organic electrochemical transistors. Nat. Rev. Mater. 3, 17086 (2018). 70. Torsi, L., Magliulo, M., Manoli, K., and Palazzo, G. Organic field-effect transistor sensors: A tutorial review. Chem. Soc. Rev. 42, 8612–8628 (2013). 71. Nketia-yawson, B., and Noh, Y. Organic thin film transistor with conjugated polymers for highly sensitive gas sensors. Macromol. Res. 25, 489–495 (2017). 72. Someya, T., Bao, Z., and Malliaras, G. G. The rise of plastic bioelectronics. Nature 540, 379–385 (2016). 73. Lanzani, G. Materials for bioelectronics: Organic electronics meets biology. Nat. Mater. 13, 775–776 (2014). 74. Feiner, R., and Dvir, T. Tissue–electronics interfaces: From implantable devices to engineered ­tissues. Nat. Rev. Mater. 3, 17076 (2017). 75. Knopfmacher, O., Hammock, M.L., Appleton, A.L., Schwartz, G., Mei, J., Lei, T., Pei, J., et al. Highly stable organic polymer field-effect transistor sensor for selective detection in the marine environment. Nat. Commun. 5, 2954 (2014). 76. Mulla, M. Y., Tuccori, E., Magliulo, M., Lattanzi, G., Palazzo, G., Persaud, K., Torsi, L. Capacitancemodulated transistor detects odorant binding protein chiral interactions. Nat. Commun. 6, 6010 (2015). 77. Stutzmann, N., Friend, R. H., and Sirringhaus, H. Self-aligned, vertical-channel, polymer fieldeffect transistors. Science 299, 1881–1885 (2003). 78. Lv, A., Pan, Y., and Chi, L. Gas sensors based on polymer field-effect transistors. Sensors 17, 213 (2017).

2 Electrical Doping of Organic Semiconductors with Molecular Oxidants and Reductants 2.1 Introduction......................................................................................... 21 2.2 Basics of Doping in Organic Materials............................................22

Stephen Barlow, Seth R. Marder, Xin Lin, Fengyu Zhang, and Antoine Kahn

Comparison to Doping of Inorganic Materials • Effects of Doping

2.3 2.4

Criteria for Dopant Choice.................................................................26 Survey of Dopants...............................................................................30


Device Examples.................................................................................. 35

p-Dopants • n-Dopants


2.6 Summary...............................................................................................38 Acknowledgments...........................................................................................38 References.........................................................................................................38

2.1 Introduction Electrical doping has long played a role in studies of conjugated polymers. Thus, for example, metallic levels of conductivity in polyacetylene were obtained by halogen doping in the 1970s [1], and PEDOT:PSS [2], in which a poly(​3,4-e​thyle​nedio​x ythi​ophen​e-1,5​-diyl​) main chain is partially oxidized (p-doped) and charged-compensated by an anionic polymer, is widely used as a hole-injection layer for organic lightemitting diodes (OLEDs). In contrast, until recently, the active layers of organic electronic devices, such as OLEDs, organic field-effect transistors (OFETs), and organic photovoltaics (OPVs), have generally been based on undoped organic semiconductors. However, in recent years, there has been a tremendous growth of interest in the controlled doping of organic semiconductors using molecular redox agents for the above-mentioned applications [3–31], as well as for thermoelectrics [32–36], and for use in hybrid devices such as lead halide perovskite solar cells [37–40], where doping can increase conductivity and improve carrier injection and collection at electrodes. In this chapter, we first examine the basic doping process in organic molecular and polymeric materials, we then discuss factors to be considered in selecting a dopant, survey some of the molecules that have been used as dopants, and, finally, discuss some case studies of the use of doping to improve device performance.



Conjugated Polymers

2.2 Basics of Doping in Organic Materials 2.2.1 Comparison to Doping of Inorganic Materials Doping in molecular or polymeric organic semiconductors is somewhat different from that in traditional crystalline inorganic semiconductors. In the latter, a few atoms from the parent lattice (e.g., silicon) are replaced by atoms that bring an extra valence electron (n-dopants, e.g., phosphorus) or one fewer valence electron (p-dopants, e.g., boron). This substitutional doping results in new donor or acceptor energy levels within the band gap of the parent semiconductor, close to the conduction or valence bands respectively (Figure 2.1). The magnitudes of the energy separations between these impurity levels and the relevant bands are often only a few times that of k BT, where k B is Boltzmann’s constant and T is absolute temperature, at room temperature (25.9 meV), allowing facile thermal excitation of carriers to the neighboring bands, in which they are extensively delocalized and highly mobile. From a statistical point of view, complete ionization of dopants is ensured at typical dopant:semiconductor atom ratios (often ca. 10–6) due to the large density-of-states at the edges of the neighboring bands. In the vast majority of organic semiconductors a band picture of carrier transport is inappropriate: intermolecular orbital overlap is generally weak compared to that within conjugated molecules or in inorganic semiconductors and, accordingly, in most cases charge carriers are confined to single (or sometimes a few) molecules or to relatively short portions of a polymer chain (polarons) and carrier transport involves hopping of the charge carrier from one site to another within a manifold of “transport” levels that are relatively close to one another both energetically and spatially. We note that the band terminology is nonetheless frequently used; however, the “bandwidths” found for organic

FIGURE 2.1  Schematic contrasting structural order (top) and electronic structure (bottom) in crystalline inorganic (left) and amorphous organic (right) semiconductors. Dark and light gray states are occupied and unoccupied, respectively, for the undoped semiconductors. n-Dopant states, which are occupied prior to electron transfer to the conduction band/electron traps/electron-transport states, are shown in black to exemplify the effects of doping (in the case of p-doping unoccupied dopant states would be located close to the valence band/hole-transport states).

Electrical Doping of Organic Semiconductors


semiconductors, particularly non-crystalline materials, generally result not from strong orbital overlap but from distributions of localized states due to differences in the local environment for different molecules or portions of polymers (Figure 2.1). The carrier localization in organic semiconductors means that their electrical properties are considerably less sensitive to impurities than those of materials such as single-crystal silicon. On the other hand, a distribution of site energies in non-crystalline organic materials, along with impurities originating in their chemical synthesis or introduced through processing, can lead to a significant density of states within the “transport gap” of the material that can act as traps for charge carriers. The differences in carrier localization, purities, and crystallinity between a typical amorphous organic and a crystalline inorganic semiconductor means that optimum doping levels in the former often correspond to dopant:semiconductor molecular ratios of 10–2 to 10–1, much higher than for the latter, where a dopant:semiconductor atom ratio of 10–3 is considered extremely heavy (and is also termed degenerate). The most conceptually simple way to achieve electrical p- or n-doping in organic materials is by introducing neutral one-electron oxidants or reductants. A molecule with an electron affinity (EA) that approaches the ionization energy (IE) of a semiconductor host can act as a p-dopant, accepting an electron from the host to form a dopant molecular anion and a host molecular cation (a radical cation for the usual case where the undoped host semiconductor is a closed-shell molecule), the position of the host-to-dopant electron-transfer equilibrium increasing with dopant EA. Similarly, a low IE is required for an n-dopant, ideally smaller than the EA of the host (as shown schematically in Figure 2.1). Coulombic attraction between the ions formed can also help drive the electron-transfer equilibrium; however, such interactions should ideally be minimized from the perspective of charge trapping (see Section 2.3). The relevant IE and EA values can be obtained using UV photoelectron-spectroscopy (UPS) and inverse photoelectron-spectroscopy (IPES) measurements on solid films (assuming that the IE/EA of the dopant are similar in the semiconductor host to the values obtained for the pure dopant), or can be estimated using electrochemical data (although solvation effects in electrochemical measurements may differ significantly from solidstate polarization effects) [41]. It is worth noting that undoped organic semiconductors are often classified as either hole- or electrontransporting materials (HTMs and ETMs) based on the function that they play in a particular device; although this requires that the relevant carrier mobility should be sufficiently large for the application in question, it also critically depends on how the IE and EA compare to those of other semiconductors and/or to the work function of common electrode materials of interest. The classification does not necessarily reflect the intrinsic ability of the material to transport either type of charge; i.e., materials usually thought of as HTMs may also exhibit large or moderate electron mobilities and reasonable hole mobilities may be obtained in ETMs. Similarly, although p-doping is most often applied to “HTMs” and n-doping to “ETMs”, p- and n-doping can in principle be applied to any organic semiconductors, leading to majority hole or electron conduction, respectively. Indeed, both n- and p-doping have been applied to several semiconductors. For example, p-i-n diodes have been demonstrated in both ZnPc [42] and CuPc [43], and field-effect transistors based on undoped and n-doped TIPS-pentacene with Au contacts function as p- and n-channel devices, respectively [16]. Some approaches to doping in crystalline organic materials have been developed by analogy with inorganic substitutional doping, in these an n-dopant molecule is designed to have a much lower IE than the host, but a molecular size and shape that are practically identical [44–48]. However, these approaches have shown limited success in improving material and device performance. The approach also requires synthesis of a new dopant specifically tailored for each host to ensure facile incorporation into its lattice. Moreover, most organic electronic devices are not based on single crystals, but rather on amorphous materials, semi-crystalline polymers (containing amorphous regimes), or microcrystalline materials, meaning that dopant molecules do not necessarily need to fit into precisely defined lattice sites. Accordingly, dopants are generally not matched to the host in size and shape but chosen based on their redox properties and other criteria as discussed below in Section 2.3.


Conjugated Polymers

2.2.2 Effects of Doping Enhancement of Conductivity The principal direct effects of doping are to increase the conductivity and to lower the barriers to chargecarrier injection and/or collection. The conductivity of a material depends on the charge-carrier mobility and on the carrier density. Both of these quantities can potentially be affected by doping. In the most simplistic picture of doping, increasing the concentration of p- or n-dopants will increase the number of host molecular cations or anions respectively, these ions acting as the charge carriers. However, as noted above, owing to the presence of impurities and/or disorder, organic semiconductors often contain appreciable densities of “deep traps”, i.e., filled states well above the filled transport levels or empty states well below the empty transport levels; charge carriers become trapped in these states such that the activation barrier required to return them to the relevant transport states is large relative to k BT. Thus, adding carriers to a semiconductor first results in the transfer of carriers to the deepest traps, then shallower traps, and, only at higher concentrations, the transport states. Even though “free” carriers are not contributed at relatively low doping levels and the highest conductivities are only obtained at high doping levels, low levels of doping can lead to large increases in conductivity by reducing the number of deep traps, thereby enhancing the mobility. Figure 2.2 shows an example where traps in C60 are filled using an n-dopant; the steepest increases of conductivity with doping level occurs in this trap-filling “ultralow doping” regime [49]. In a similar fashion, hole traps have been directly observed in CuPc using UPS and their energetic distribution has been found to be narrowed, and their number to be decreased, as the level of ultra-low p-doping is increased [50]. Note that, as discussed above, “low” doping levels in an organic context are much higher than what is considered low in crystalline inorganic semiconductors. As the dopant concentration is increased further, the increase in conductivity is less dramatic and, in this particular case (in Figure 2.2), is close to proportional to the doping level (indicated by the line of

FIGURE 2.2  Experimental data (symbols) and simulation results (lines) for the conductivity of a C60 layer as a function of (RuCp*mes)2 (see Figure 2.7 for structure) doping for three temperatures; the line of unit slope is shown as a guide for the dye. Reproduced with permission from Olthof et al., Phys. Rev. Lett. 109: 176601 (2012), copyright 2012 the American Physical Society.

Electrical Doping of Organic Semiconductors


unit gradient on the log-log plot), consistent with increased carrier density, rather than increased mobility, being the main contributor to the conductivity. In other cases, the mobility may still be doping-level dependent in this regime, albeit less so than in the trap-filling regime, either decreasing or increasing depending on the effects of the dopants on the ordering, and perhaps also the conformation, of the semiconductor [51]. Finally, at the highest doping levels in Figure 2.2, the conductivity (and inferred mobility) begins to fall again; this is attributed to disruption of charge-transport pathways through the host film by the high content of dopant ions. Lowering of Injection Barriers Lowering the effective barrier for charge-carrier injection and/or collection can be achieved via two distinct mechanisms. At a non-reactive interface between an organic semiconductor and a conducting electrode, the alignment of energy levels across the interface is dominated by the difference between the work functions of the two materials. p-Doping of the organic semiconductor drives its Fermi level toward the filled transport states, i.e., towards the highest occupied molecular orbital (HOMO) level, increasing the material’s work function (as measurable by techniques such as Kelvin probe and UPS, the latter being particularly useful as it yields both the work function and the IE, the difference between which decreases with increased levels of doping). The alignment of the electrode and semiconductor Fermi levels thus lowers the hole-injection barrier (Figure 2.3a,b). Conversely, when n-doping the semiconductor, the Fermi level moves towards the empty transport states, i.e., the lowest unoccupied molecular orbital (LUMO) level, leading to a lowering of the electron-injection barrier from the electrode Fermi level to the LUMO. This mechanism results in a true lowering of the injection barrier. A second mechanism can also contribute in the case of an “interactive” interface, where the energy difference between the relevant transport levels and the electrode Fermi level is more narrowly restricted by chemical interactions between the semiconductor and the electrode, such that it is largely independent of doping. Here the effective lowering of the barrier takes place through the formation of a depletion region, i.e., a region free of charge carriers, in the semiconductor, the width of which inversely depends

FIGURE 2.3  Changes in the electronic structure of the interface of an electrode with a hole-transport material upon p-doping for the cases of (b) a non-interactive interface and (c) an interactive interface with Fermi-level pinning. The effective hole-injection barrier is lowered in both cases.


Conjugated Polymers

on the dopant concentration. For the example of a p-doped organic semiconductor, the Fermi level will be closer to the HOMO in the bulk of the material, but can be closer to mid-gap at the interface due to the energy-level alignment imposed by the interface interaction. The difference between bulk and interface is compensated by a downward “band bending” with formation of a depletion region. If the dopant concentration is a few percent, which is not unusual in organic semiconductors, the depletion region width collapses down to a few nanometers, and electrons can be injected via tunneling through the barrier instead of thermionically over the barrier, thus considerably lowering the effective barrier (Figure 2.3a,c) To illustrate doping-enhanced charge-carrier injection at an organic semiconductor interface, we consider hole injection from a gold electrode into a small-molecule semiconductor, spiro-TAD (Figure 2.4; the structure of spiro-TAD is shown later in Figure 2.5) [52]. The Au electrode, previously exposed to ambient air, has a work function of ca. 4.7 eV, whereas spiro-TAD has an IE of 5.4 eV. The low-field current–voltage (I–V) characteristics of the spiro-TAD film (100 nm thick) measured between Au electrodes separated by an 80 µm gap are shown in Figure 2.4. The bulk of the film is undoped. In the first case, the undoped film is in direct contact with Au. In the second case, a thin (5 nm) “injection layer” of spiro-TAD p-doped with F6TCNNQ (the structure of which is shown later in Figure 2.6) is deposited in ultrahigh vacuum (UHV) onto the Au electrode prior to UHV deposition of the undoped film. The three-orders-of-magnitude increase in the current, coupled with the change in the I–V characteristics from injection-limited to an ohmic regime (slope equal to unity at low field), is formal proof of the effective lowering of the injection barrier.

2.3 Criteria for Dopant Choice As noted above, p- and n-dopants – most often simple one-electron oxidants and reductants, ­respectively – should, in general, be chosen to be sufficiently strong to transfer charge carriers (holes or

FIGURE 2.4  Log-log plot of current density vs. electric field (low bias, 300 K) for spiro-TAD planar diode structures with Au electrodes, and with and without a 5 nm injection layer (spiro-TAD doped with 20 mol% F6TCNNQ). The inset shows the structure of the device, in which inter-electrode gaps are 80 µm. Adapted with permission from Zhang and Kahn, Adv. Funct. Mater. 28: 1703780 (2018).

Electrical Doping of Organic Semiconductors


FIGURE 2.5  Structures of some organic semiconductors that have been doped and that are mentioned in the text. Note that NDI and PDI are generic structures and a variety of derivatives have been studied.

electrons) to the transport levels of the host material of interest. These electron-transfer reactions should occur cleanly, generating only the desired semiconductor ions alongside well-defined and stable dopantderived species. Moreover, stronger dopants, where the initial doping reaction is more exergonic, will render the doping more “irreversible”, which may be advantageous in cases where the neutral species formed by the endergonic back reaction can be lost from the doped film, for example, due to a high volatility. On the other hand, dopants that are insufficiently strong to contribute carriers to the transport levels may still be useful for passivating deep traps (see Section Although one-electron transfer is the most common mechanism for doping, in some systems other mechanisms play a role. Some combinations of planar semiconductors and planar dopants, such as pentacene doped with F4TCNQ (see Figure 2.5 and Figure 2.6, respectively, for the chemical structures), form ground-state charge-transfer (CT) complexes in which there is overlap of donor and acceptor π orbitals and partial charge transfer (sometimes referred to as “orbital hybridization” in the literature); these CT complexes can be thermally excited to release charge carriers [53–55]. Although doping through CT complexes can enhance conductivity, it is typically less effective than doping through integral electron-transfer reactions; the latter can be favored by using stronger dopants in order to increase the driving force for electron transfer and by selecting bulkier dopants whose frontier orbitals are less accessible for π-overlap with those of the semiconductor.

FIGURE 2.6  Chemical structures of some p-dopants referred to in the text and in Table 2.1; all of these examples function as simple one-electron oxidants.


Conjugated Polymers

In other cases, dopants are specifically chosen or designed to work through coupling of electrontransfer and other reactions (particularly those discussed in Section for the case of air-stable n-dopants). For some of these systems, the formation of side products is inevitable; whether their identity is clearly established, whether they are incorporated in the doped film, and whether they adversely affect carrier transport varies from system to system. Many applications of dopants in organic semiconductors require that the doping be spatially localized in a particular portion of the device. However, dopant ions – or neutral dopant molecules in equilibrium with dopant ions [56] – may diffuse and drift within a doped film, leading to a loss of this spatial localization. These issues are likely to be particularly acute in the case of small dopant ions and, indeed, the diffusion of lithium ions from the charge-injection region of organic light-emitting diodes to the emissive region has been identified as a mode of device degradation [57]. Planar dopants can also be rather prone to diffusion between other planar molecules [58]. Diffusion and drift are likely to be minimized when the dopant ions are relatively large and three-dimensional in nature. Indeed, several studies have shown diffusional stability for three-dimensional molecular dopants to be superior to that of planar analogues [9, 11, 59]. However, diffusional stability will also depend on the structure and rigidity of the host matrix; for example, Mo(tfd-CO2Me)3 (see Figure 2.6 for chemical structure) very rapidly diffuses from doped P3HT into undoped P3HT, yet does not diffuse from doped P3HT into an undoped P3HT:PCBM (see Figure 2.5 for chemical structures) bulk-heterojunction film [60]. Dopant-ion size is also expected to influence the Coulombic interaction between dopant ion and the charge-carrying species in the semiconductor [51]. Although this interaction can help favor the initial electron-transfer reaction when the dopant ions are small, it effectively creates a trap whereby the carrier prefers to reside on a molecule immediately adjacent to the dopant ion and so presumably release of the carrier is hindered, and the carrier mobility decreased, relative to what is found for larger dopants. On the other hand, the beneficial effects of large ion size on diffusional stability and charge trapping may, to some extent, be offset by greater disruption of the packing of the semiconductor, and by an increase of the average distance between semiconductor molecules, which could also reduce mobility (as seen in the decrease in conductivity at the highest dopant levels in Figure 2.2). This trade-off between beneficial and adverse effects will likely lead to different optimum doping levels for differently sized dopants and for different hosts; for example, more crystalline hosts might be more sensitive to disruption by large ions than amorphous materials. Another important consideration is that processing of the dopant should be compatible with processing of the semiconductor and of other portions of the desired device structure. In general, organic semiconductor devices are fabricated using either evaporation in UHV or are processed from solution using methods such as spin-coating, blade coating, and slot-die coating. The evaporation approach is particularly suitable for multilayer structures involving both doped and undoped layers. Vacuumprocessed films are usually doped by coevaporation of the dopant and the semiconductor. The dopant should be volatile at temperatures below that at which it decomposes, yet not sufficiently volatile than uncontrolled evaporation takes place in the vacuum chamber. For example, the n-dopant CoCp2 (Figure 2.7) is rather too volatile for use in standard crucibles used in evaporation chambers, although this issue can be circumvented by admission of the dopant vapor to the UHV deposition chamber via a leak valve [61]. The widely used p-dopant F4TCNQ is also rather volatile; as well as leading to complications with vacuum-processing doped films, this volatility means the dopant can be gradually lost from films through back-electron transfer to the host, followed by sublimation of the neutral species [62, 63]. Perhaps the most commonly used approach to doping solution-processible materials is to mix dopant and semiconductor solutions together prior to film deposition. This requires that a solvent, or solvent mixture, can be identified that dissolves the dopant, the semiconductor, and the dopant:semiconductor ion pair (or partially oxidized or reduced semiconductor polymer chain in combination with dopantderived counterions). However, other approaches have been developed whereby an undoped semiconductor film is dipped in a dopant solution [17, 30], where a dopant solution is spin-coated onto a semiconductor film [64, 65], or where a dopant is evaporated onto a semiconductor film [34, 66]. For the

Electrical Doping of Organic Semiconductors


first two of these approaches the dopant should ideally be soluble in a solvent that does not dissolve the semiconductor. These approaches do require that the dopant be able to diffuse into the semiconductor, at least to some extent, either during the dipping or coating process, or during subsequent thermal annealing. Accordingly, their effectiveness will depend on the specific semiconductor:dopant combination. For example, spin-coating and annealing of (RuCp*mes)2 (shown later in Figure 2.7) onto P(NDI2OD-T2) (also known as N2200, shown in Figure 2.5) is unsuccessful, but works well for P(BTP-DPP) (Figure 2.5), a much less ordered polymer with comparable EA [65]. Sequential deposition is likely limited to relatively thin films; thicker (mm-thick) films can, however, be successfully doped by incorporating small NaCl particles into the film, treating with water to remove the salt, and infiltrating dopant into the voids in the resulting polymer “foams” [35]. Several studies have also pointed out the importance of dopant miscibility with host semiconductors [67], and have demonstrated improved miscibility and doping effectiveness by modifying host structures, in particular using oligoether side chains [36, 68]. Alternatively, rather than adding a molecular oxidant or reductant to the semiconductor, the radical cation or anion of the semiconductor can potentially be synthesized using any convenient oxidant or reductant of suitable strength and purified as a salt with a redox-innocent counterion, which can then be added to the processing solution alongside the neutral semiconductor. The approach has been rarely used and will generally be unsuitable for n-doping of molecules with low EA, where the corresponding radical anions will likely be highly air and moisture sensitive, complicating their purification, isolation, storage, and use, or p-doping molecules with very high IE, where the radical cation may be highly sensitive to reduction by traces of water. However, the approach has been used for some molecules with moderate redox potentials, especially diamine-based hole-transport materials; for example, TMTPD (Figure 2.5) doped with TMTPD+SbF6– has been used to study the doping-level dependence of conductivity and activation barrier for transport [69], and spiro-OMeTAD2+(TFSI–)2 (synthesized by oxidation of the neutral molecule using Ag+TFSI–; TFSI– = N(SO2CF3)2–) has been used to dope the spiro-OMeTAD (Figure 2.5) hole-extraction layers of haloplumbate perovskite ((MeNH3)PbI3) solar cells [70]. The use of zwitterionic molecules in which radical anions of perylene diimides (PDIs) are tethered to tetraalkylammonium cations as dopants for neutral PDIs [45] can be regarded as a variation on this approach. Finally, the chemical stability of dopants can be an important consideration. Strong n-dopants that operate via one-electron transfer necessarily have low IE and so tend to be sensitive to oxidation by oxygen and/or water. This issue has inspired the development of several classes of dopants in which the electron transfer is coupled to other chemical reactions in order to obtain moderate air stability to facilitate,

FIGURE 2.7  Chemical structures of some n-dopants and related species referred to in the text and in Table 2.2.


Conjugated Polymers

for example, transfer of the dopant between a glove-box and a vacuum chamber, or accurate weighing in air, while retaining adequate dopant strength (see Section On the other hand, strong oxidants can be sensitive to reduction by adventitious water, especially in basic solvents. Moreover, the dopant should ideally not exhibit other undesirable reactivity, for example, towards common organic functional groups, or common solvents. We note that the air and water stability of the doped film and casting solutions of doped material will largely be determined by the stability of the semiconductor ions that are formed, although the dopant counterion may also limit stability in some cases, for example, if it is particularly hygroscopic.

2.4 Survey of Dopants This section surveys some of the classes of molecules that have been used as p- and n-dopants. Dopant strength is quantified where possible by redox potentials in non-aqueous solvent, quoted vs. the ferrocenium/ferrocene couple (FeCp2+/0, in some cases converted from values quoted against other references using values from refs. [71] and [72]) and by solid-state EA or IE, as directly measured using IPES or UPS. Figure 2.5 shows the chemical structures of some of the molecular and polymeric organic semiconductors that have been p- and n-doped and that are discussed in this section, or in other sections of this chapter. Redox potentials and directly measured IE or EA values, where available, are given in the text. Values of dopant and semiconductor redox potentials from different sources should, however, be compared with caution, since many different reference couples have been used in the literature, it is not always clear whether the original data are correctly referenced, and different studies variously quote onset, half-wave, and peak potentials. Potentials are also, in principle, solvent-dependent, although this effect is expected to be less dramatic for typical organic semiconductors and for most molecular dopants than for redox couples involving small ions, such as NO+. We have, in general, refrained from quoting electrochemically based estimates of IE and EA (often referred to as “HOMO and LUMO energies”, respectively) from the literature as these values are plagued by additional confusion arising from the use of several different methods for estimating these quantities from redox potentials [41]; however, readers can, of course, estimate these values from the potentials quoted using their preferred approach.

2.4.1 p-Dopants Inorganic p-Dopants In the early days of conducting polymer research, molecular halogens, particularly iodine, were commonly used as p-dopants. However, they exhibit a number of drawbacks. Iodine is much more easily handled than chlorine or bromine, but is not a particularly strong oxidant (E(I2/2I–) = –0.14 V in MeCN [71]) and is still rather volatile; accordingly, doping with iodine (often by exposure of a polymer film to I2 vapor) is often found to be reversible, as molecular iodine slowly sublimes out of the film (as seen, for example, in the case of iodine-doped poly(​4,4'-​dioct​ylcyc​lopen​tadit​hioph​ene-2​,6-di​yl), the conductivity of which decreases in vacuum, whereas that of the same polymer doped with DDQ does not [73]). High-valent halides including FeCl3, AsCl5, and SbCl5 have also been used [3, 5, 74–78]. SbCl5 exemplifies some of the pros and cons of this approach; at least in some cases, it can act as an effective dopant [3] and it forms the relatively stable SbCl6– counterion, but it can exhibit competing Lewis-acid reactivity, is highly moisture sensitive, liberating potentially mobile Cl– ions, and SbIII side products are inevitably formed through reactions of the type shown in Eq. 2.1 (where SC = semiconductor, represented as a closed-shell molecule, although similar equations can be written for open-shell molecules and for polymeric semiconductors).

3SbCl 5 + 2SC → 2SC •+ + 2SbCl 6 − + SbCl 3 (2.1)


Electrical Doping of Organic Semiconductors

Molecular oxygen can also act as a p-dopant; indeed, low-IE semiconductors are sometimes found to be unintentionally p-doped by adventitious exposure to O2 [79, 80], although the identity of the O2-derived anions is not always clear. However, O2-doping can be promoted by the addition of other non-redoxactive species that help stabilize the oxygen-based products. For example, Li+TFSI–, in combination with exposure to air, has been widely used to dope spiro-OMeTAD (Figure 2.5, Eox = +0.01 V [81]; IE = 5.0 eV [82]) for use as a HTM in dye-sensitized and haloplumbate perovskite solar cells [83, 84], the lithium ions presumably stabilizing O2-derived anions. The combination of O2 and a Brønsted acid (HX, either solution- or vacuum-processed) can be particularly effective; H+ increases the oxidant strength by leading to the formation of water as a stable byproduct (Eq. 2.2), while the charge-balancing counterion in the film can be varied through the choice of acid. 4HX + O2 + 4SC → 4SC •+ + 4 X − + 2H2O (2.2)

H+ itself, especially when poorly solvated, can potentially also act as an oxidant for lower IE semiconductors, forming H2 as a side-product; for example, non-aqueous Brønsted acids have been used in the inert-atmosphere doping of spiro-OMeTAD [37]. Organic Lewis acids, such as B(C6F5)3, have also been used as p-dopants [6, 40, 63] and may operate through similar mechanisms if trace water and/or oxygen are present, but in other cases may coordinate to Lewis basic sites in the semiconductor [85]. Nitrosonium salts, NO+X–, are highly oxidizing (see Table 2.1) solution-processible species, generating neutral NO• as a gaseous side product and leaving the counter anion to balance the charge introduced to the semiconductor [24, 86, 87]. On the other hand, they are highly moisture sensitive. High-valent oxides, notably MoO3, can be used as vacuum-processible p-dopants. Bulk MoO3 is an extended solid, consisting of vertex- and edge-sharing MoO6 octahedra, that has very deep-lying valence and conduction bands and that acts as a n-type semiconductor with a high work function TABLE 2.1  Redox potentials and solid-state electron affinities for selected one-electron oxidants p-dopant

Ereda/V vs. FeCp2+/0


C60F36 F3TCNQ-Ad1 DDQ F4TCNQ F6TCNNQ Mo(tfd-CO2Me)3 Mo(tfd)3 Mo(tfd-COCF3)3 C60F48 N(C6H4-p-Br)3•+X– CN6CP NO+X–

–0.14 [100]c +0.05 [97]d +0.13 [71] +0.18 [99] +0.24 [99] +0.12 [101] +0.28 [101] +0.39 [101] +0.48 [103] +0.70 [71] +0.78 [99] +0.56 to +1.00 [71]e

– – – 5.2 [58] 5.6 [52] 5.0 [27] 5.6 [102] – – – – –

Electrochemical potentials in non-aqueous solvents. Adiabatic values in the solid state measured by inverse photoelectron spectroscopy. c Value re-referenced to FeCp +/0 by comparing potentials 2 reported for C60F36 and C60F48 in ref. [100] to a value for the C60F48 potential vs. FeCp2+/0 obtained from ref. [103]. d Value re-referenced to FeCp +/0 by comparing potentials 2 reported for F3TCNQ-Ad1 and F4TCNQ in ref. [97] to a value for the F4TCNQ potential vs. FeCp2+/0 obtained from ref. [99]. e Strongly solvent dependent. a



Conjugated Polymers

(ca. 6.9 eV) appropriate for hole injection into high-IE organic semiconductors, such as OLED materials [88]. Similar electronic properties and high work functions are found for other high-valent oxides including WO3 [89] and V2O5 [90]. However, MoO3 sublimes as discrete trimers in which MoO4 tetrahedra share vertices, so can potentially act as a molecular p-dopant when cosublimed with organic semiconductors; indeed, it is one of the strongest p-dopants to have been studied, effectively doping CBP (IE = 6.2 eV) [91]. However, relatively poor efficiencies of doping using MoO3 and other high-valent oxides have been attributed to the tendency of the oxide molecules to form aggregates and phase separate from organic host materials [92, 93]. Another MoVI oxide derivative is the large cluster molecule phosphomolybdic acid (H3PMo12O40), which, in contrast to MoO3, is soluble but not vacuum-processible; H3PMo12O40 has been used in the immersion method to effectively dope and insolubilize polymers with estimated IEs of up to at least 5.3 eV [29, 30]. In particular, it is capable of infiltrating films of organic solar-cell blends, but only to a certain extent, thus providing a means of doping only in the vicinity the hole-extraction interface of the device [30]. The mechanism of doping is not clearly established; it can potentially act as a one-electron oxidant and, although reported redox potentials are moderate (e.g., ca. +0.03 V in CH2Cl2/MeCN [94]), its dopant strength is pH and solvent dependent. Organic and Metal–Organic p-Dopants In contrast to the purely inorganic p-dopants described above, much of the recent work on p-dopants has focused on well-defined organic or metal–organic molecular species, particularly examples that can cleanly accept an electron to form a stable ion. Examples are shown in Figure 2.6 and redox potentials and/ or EA values are given in Table 2.1. The planar F4TCNQ molecule is one of the most widely investigated molecular p-dopants [4, 58] and can be both solution and vacuum processed, although its volatility and its tendency to diffuse within host materials are rather high [9, 11, 59]. Planar acceptors such as F4TCNQ are also particularly prone to form CT complexes, rather than undergoing integral electron-transfer reactions, especially when combined with planar host materials, such as pentacene and oligothiophenes [54, 55]. Various related molecules have been developed to address some or all of these drawbacks. F6TCNNQ is less volatile and a stronger oxidant, although it is also less soluble [8, 95, 96]. The bulky or long-chain hydrocarbon moieties of molecules such as F3TCNQ-Ad1 and F4OCTCNQ improve solubility, miscibility with host materials, and, in the former case, may help limit diffusion [97, 98], although both are slightly weaker dopants than F4TCNQ. BAPD is also a weaker dopant, but is significantly non-planar, leading to a lower tendency to diffuse, and can be sublimed in a more controlled fashion than F4TCNQ [9]. CN6CP, on the other hand, is the strongest organic neutral p-dopant reported to date [99]. In common with many other highly oxidizing species, however, it is unstable in air and is reduced by many common solvents. It is poorly soluble in non-reactive solvents but can be “solution-processed” as a suspension, p-doping the high-IE polymer PDPP(6-DO)2TT (Eox = ca. +0.4 V) to obtain very high conductivities of 30–70 S cm–1, and can be sublimed, so can likely be used as a vacuum-processible dopant. Three-dimensional and less volatile molecular dopants that can be both solution and vacuum processed include fluorinated fullerenes, such as C60F36 [7, 59, 95, 104] and C60F48 [20, 105]. For example, C60F36 has been shown to diffuse much more slowly into MeO-TPD (Eox = +0.09 V [106], IE = 5.1 eV [95]) than F4TCNQ [59]. On the other hand, fluorofullerene dianions can lose fluoride ions [100, 103], meaning that mixtures of species may, in some cases, be present in doped films. Molybdenum tris(dithiolene)s, such as Mo(tfd)3 and more soluble substituted derivatives (Figure 2.6), are also more three-dimensional and less volatile than F4TCNQ, form stable monoanions, can be processed from solution or by evaporation, and have been used to p-dope materials such as NPB (Eox = +0.30 V, IE = 5.4 eV) [102], P3HT (Eox = +0.18 V [80], 4.7 eV) [27], and TIPS-pentacene (Eox = +0.4 V [107], 5.2 eV) [19]. “Magic blue”, N(C6H4-p-Br)3•+SbCl6–, is an ionic oxidant that has occasionally been used as a solution-processible p-dopant for organic semiconductors [3, 108–110]; here the side product is the neutral triarylamine, which may or may not be incorporated in the film or removed with a low-polarity organic solvent, and the SbCl6– anion will balance the holes that are created. Other salts of the same oxidizing cation could in principle be used to introduce other counterions of varying sizes and shapes.


Electrical Doping of Organic Semiconductors

2.4.2 n-Dopants One-Electron Reductants Alkali metals are highly reducing (e.g., ca. –3.0 V for Na in THF [71]); however, they are not easily handled due to their air-sensitivity and volatility and they can exhibit reactivity that competes with the desired electron-transfer chemistry, for example, they can cleave aryl-halogen bonds. As noted in Section 2.3, the small monoatomic ions can be very mobile within doped films and can act as electrostatic traps for the contributed charge carriers [57, 111]. More recent work on n-doping has, therefore, utilized reducing neutral organic or, more commonly, metal–organic molecules that form stable molecular cations (Figure 2.7, Table 2.2). First we will consider those that act as simple one-electron reductants, D, which include both closed-shell molecules that form open-shell cations on doping – e.g., TDAE and W2(hpp)4 – and open-shell molecules that form closed-shell cations – e.g., CoCp2 – and which react as shown in Eq. 2.3 (SC is represented here as a closed-shell molecule forming a radical anion since most ETMs are indeed closed-shell species, although a few open-shell semiconductors, such as CuPc [112], have also been n-doped with one-electron reductants). D + SC → D+ + SC •− (2.3)

TDAE has been used as a one-electron reductant for conjugated polymers including P(NDI2OD-T2) (Ered = –1.03 V [113], EA = 3.9 eV [114]) and the ladder polymer BBL (Ered = –0.96 V [115]); TDAE is a volatile TABLE 2.2  Redox potentials and solid-state ionization energies for selected reductants n-dopant

Eoxa/V vs. FeCp2+/0

TDAE PyB• CoCp2 (RhCp2)2 Ru(terpy)2 (Cyc-DMBI)2 CoCp*2 (RhCp*Cp)2 (RuCp*mes)2 MeO-DMBI•c W2(hpp)4

–0.99 [124]c –1.06 [125]d –1.33 [71] –1.72 [127]f –1.93 [128]dg ca. –1.9 [129]f –1.94 [71] –1.97 [127]f –2.04 [127]f –2.22 [130] –2.25 [123]

IEb/eV – 4.3 [126] 4.1 [61] 4.0 [43]h – 3.3 [112] – – – 2.4 [121]

Oxidation potential in non-aqueous solution. Adiabatic solid-state ionization energy measured by UV photoelectron spectroscopy. c Assuming the D• radical is generated on sublimation of the corresponding D+X– salt and so is the active reductant when vacuum processed. d Re-referenced to FeCp +/0 using a value for FeCp +/0 vs. SCE in 2 2 the relevant solvent from ref. [71]. e Data for a close analogue in which Et groups are replaced by Me. f Effective potential for 2D+/D couple estimated using the D+/D• 2 potential and DFT estimate of the free energy of dissociation for D2. g Re-referenced also assuming SSCE is at –0.04 V vs. SCE [72]. h IE of dimer does not measure overall dopant strength since subsequent bond cleavage and ionization of a neutral molecule takes place in the doping reaction [43]. a



Conjugated Polymers

liquid and doping was accomplished by exposure of the polymer film to TDAE vapor [116]. It has also been used to “de-dope” adventitiously p-doped materials [117]. Cobaltocene, CoCp2, is a somewhat stronger dopant and is both solution and vacuum processible (although rather too volatile for use in a conventional evaporator) [61, 118, 119]. Stronger dopants still include CoCp*2 (which is also rather volatile) [112, 120], Ru(terpy)2 [42], and even W2(hpp)4 [18, 21, 28, 32, 121], which has the lowest reported gas-phase IE of any stable molecule (3.5 eV) [122] and is also very easily ionized in both films and solutions (see Table 2.2) [121, 123]. The low IEs of these compounds inevitably lead to air-sensitivity, with the strongest dopants being particularly sensitive. Air-Stable n-Dopants As noted above, one-electron reductants sufficiently strong to reduce many semiconductors of interest must have low IEs and so are air-sensitive. Here we discuss how air-stable, or at least less air-sensitive, yet useful, n-dopants can be obtained by identifying reductants in which the electron-transfer process leading to doping is coupled to a chemical reaction. In some cases, an air-sensitive one-electron reductant can be liberated from an air-stable precursor. Halide salts of some stable organic cations, such as MeO-DMBI+ [131, 132] and of dyes such as pyronine B (PyB+) [22, 126, 133–135], constitute one such class of air-stable precursors that can be used in vacuum deposition of doped films. On heating, these materials release the corresponding highly reducing organic radicals (MeO-DMBI• and PyB•, respectively), presumably accompanied by molecular halogens according to Eq. 2.4, although in some cases other organic species are apparently also formed [126, 132].

2D+ X − = 2D• + X 2 (2.4)

Simple very stable salts consisting of redox-inert cations, such as tetraalkylammonium ions, and inorganic anions, such as halides, have also been used as solution-processible n-dopants for moderate-EA materials. Despite the non-reducing character of either the cation or anion, Bu4NF has been found to generate the radical anions of naphthalene diimides (NDIs, Ered = ca. –1.1 V [136]) in solution [137] and has been used as an n-dopant for fullerenes and other materials with comparable EA [31, 138, 139]. Recent work indicates that the reaction of fluoride and NDI is not clean and it has been suggested that the active reducing agent is the conjugate base of the solvent, which is formed through deprotonation of the solvent by the fluoride ion [140], while in the case of C60 reduction, nucleophilic attack on the fullerene may be the first step [141]. Regardless of the doping mechanism, it is clear that, in common with some other approaches to air-stable n-dopants, side-products are inevitable with this approach; whether these adversely affect the electronic properties or compromise device stability, relative to what can be achieved with a simple one-electron reductant, will likely depend on the specific dopant, semiconductor, and properties of interest. Also, in common with the use of amines and hydride donors (discussed below), a doping strength for nucleophilic and/or basic anions that is independent of the semiconductor properties and doping conditions cannot be defined. Simple alkylamines such as NEt3 have long been used for photoreductants for molecules such as perylene diimides (PDIs) (Ered = ca. –1.0 V [142]) [143]; although electron transfer from NEt 3 to groundstate PDIs is highly endergonic, electron transfer to the excited state is feasible and is followed by rapid decomposition of the amine radical cation to more stable species [144], rendering the process irreversible. Examples of “self-doped” ETMs have been studied in which trialkylamines are covalently tethered to PDIs or other semiconducting moieties (species that, in some cases, have been generated by thermal treatment of the corresponding tetraalkylammonium hydroxides) [145], and there is evidence for irreversible chemistry following amine-to-NDI electron transfer in other tethered systems [146]. Similar reactions also likely take place where moderate-EA ETMs, such as fullerenes and P(NDI2OD-T2), are in contact with poly(ethyleneimine) materials [147], which have been used as electrode modifiers in various organic electronic devices [148]. Molecules such as leuco-crystal violet (CV-H) and DMBI-H derivatives can be regarded as hydride reduction products (DH) of stable organic cations and are rather stable, yet can be used for both solution

Electrical Doping of Organic Semiconductors


and vacuum doping of semiconductors [13, 16, 33, 36, 38, 67, 149–153], which results in the formation of the stable D+ cations (CV+ and DMBI+). N-DMBI-H has been particularly widely used. In general, the ability of semiconductors to be doped by these molecules will depend on their thermodynamic and/or kinetic abilities to accept hydride ions or hydrogen radicals, in addition to their EA. At least in the case of fullerenes, for which these dopants are particularly effective, partially hydrogenated semiconductor molecules [130, 149] are formed alongside the cations and semiconductor radical anions according to Eq. 2.5.

DH + (1 + 1 / x ) PCBM → D+ + PCBM•− + (1 / x ) PCBMH x (2.5)

Finally, dimers (D2) formed by certain highly reducing odd-electron organometallics, such as (RhCp*Cp)2 and (RuCp*mes)2 [43, 127], or by certain organic radicals, such as (Cyc-DMBI)2 [129, 154], have been found to be very effective vacuum- and solution-processible n-dopants that exhibit moderate air stability. The dimers react with the semiconductors to form semiconductor radical anions and stable monomeric organometallic or organic cations (with at least two mechanisms possible [129, 155]) as shown in Eq. 2.6; in contrast to many of the other air-stable dopants discussed above, there is no inevitable formation of additional unwanted side-products.

D2 + 2SC → 2D+ + 2SC •− (2.6)

The ability of a semiconductor to be doped by many of the other n-dopants discussed in this section, such as halide ions as DMBI-H derivatives, may depend both on its EA and its ability to participate in other reactions; however, the doping strength of a dimer depends only on the redox potential of the corresponding monomeric molecule and the free energy of dissociation of the dimer, and so the ability for a semiconductor to be doped depends only on its EA [127, 129]. Moreover, the dimer classes examined to date are also rather strong dopants, n-doping materials such as CuPc (EA = 3.1 eV), TIPS-pentacene (Ered = –1.45 V [155]; EA = 3.0 V), and pentacene (EA = 2.8 eV) [43, 129]. Doping of an OLED transport material (POPy2, Ered = –2.24 V; EA = 2.2) has also been achieved using photoirradiation (see Section 2.5.1 for further discussion) [10]. On the other hand, although much more air stable that simple one-electron reductants with comparable doping strengths, these dopants can only be handled relatively briefly in air and still require long-term storage in inert atmosphere.

2.5 Device Examples Since the late 1990s, doping has been increasingly used to improve the performance of organic devices. Early work with alkali metals to n-dope electron transport layers (ETLs) [111] or molecular oxidants to p-dope hole transport layers (HTLs) [4] in OLEDs led to considerable improvements in device performance. Doping has since been extensively applied to OLEDs and OFETs, and, to a lesser degree, to OPV cells, mostly, but not exclusively, in the context of lowering contact barriers. Many of the molecular dopants discussed above have been used in such devices. We illustrate their use here with discussion of a few recent examples.

2.5.1 OLEDs In OLEDs, n- and p-doping has been primarily directed to reducing electron- or hole-injection barriers and increasing the conductivity of ETMs or HTMs in order to reduce the driving voltage necessary to obtain suitably intense electroluminescence. Doping these layers also allows the use of thicker layers that can offer more protection of the active emissive layer during metallization and final processing without paying a price in operating voltage. However, the low-EA ETMs required for green- and blueemitting OLEDs are particularly challenging to n-dope; although alkali metals are sufficiently strong reductants, the mobile metal ions can lead to device instabilities, and most molecular dopants are


Conjugated Polymers

insufficiently strong reductants. Recently, however, the class of cleavable dimers mentioned above has opened new pathways for n-doping OLED ETMs. The RuCp*mes monomer would be highly reducing (Eox = –2.67 V [127]), but the compound exists as a dimer (RuCp*mes)2 (Figure 2.7), the effective thermodynamic reducing power of which (E(D+/0.5D2) = –2.04 V, Table 2.2 [127]) falls short of that required to dope POPy2 (Figure 2.5; Ered = –2.24 V; EA = 2.2). However, photo-activation of a (RuCp*mes)2:POPy2 film leads to efficient and stable n-doping [10]; diffusion and/or reorientation of monomer cations away from one another following bond cleavage, along with two endergonic reaction steps for the reverse reaction, is thought to provide a sufficient kinetic barrier to spontaneous de-doping. The doped layer was used as the ETL in an Alq3-based OLED. The structure and energy diagram of the OLED are shown in Figure 2.8a,b. PEDOT:PSS (work function = 5.0 eV) is more usually used as a hole-injecting electrode, but was chosen here as the electron-injecting contact to demonstrate the power of doping to eliminate constraints on contact work function. The beneficial impact of doping is clearly demonstrated through the current density, luminance, and external quantum efficiency (EQE) of the OLED in Figure 2.8c. The undoped ETL allows negligible electron current injection and, therefore, there is no measurable emission from the OLED, whereas the doped ETL yields a significantly larger current and strong emission. The maximum EQE value for the OLED of over 1% is standard for the simple fluorescent Alq3 emitter. Green-emitting phosphorescent OLEDs, with a much higher EQE of 18%, were also realized with the same doped ETL using 2,2′,​2”-(1​,3,5-​benze​netri​yl)-t​ris(1​-phen​yl-1H-benzimidazole tris(2-phenylpyridinato)iridium (TPBI:Ir(ppy)3) as the emissive layer [10].

FIGURE 2.8  (a) Structure and (b) estimated energy-level diagram of an Alq3 OLED using POPy2 as the ETL (Alq3 = tris(8-hydroxyquinolinato)aluminum; for NPB see Figure 2.5). (c) Current density and luminance, and (d) EQE of the OLED with an undoped or (RuCp*mes)2-doped ETL. Ru/O ~ 1.4 corresponds to a 10% molar doping concentration. The dashed line in (b) shows the optical bandgap of Alq3. Adapted from Lin et al., Nat. Mater. 16: 1209 (2017).

Electrical Doping of Organic Semiconductors


2.5.2 OFETs Doping has also found important uses in OFETs. A natural application is in the reduction of metal– organic injection and collection barriers (as described in Section in pursuit of lower contact resistance, RC. Considerable reduction of RC has been demonstrated for both p-channel (hole-conducting) and n-channel (electron-conducting) OFETs, for example using the p-dopant Mo(tfd)3 (Figure 2.6) at Au:pentacene contacts [12], and the n-dopant (RhCp2)2 (Figure 2.7) at Al:C60 contacts [15], respectively, leading in both cases to increases in measured charge-carrier mobility. Another application of doping to OFETs focuses on the lowering of the threshold voltage, Vth, by the incorporation of very low concentrations of dopants into the channel of the device. The mechanism relies on the (partial) filling of deep gap states by charges originating from the dopants, which is equivalent to the action of turning on the gate voltage in an undoped OFET. Filling these trap states also aids in carrier transport as mentioned in Section For example, low concentrations (molar ratio of ca. 10–3) of the n-dopant (RuCp*mes)2 were found to significantly reduce Vth for C60 OFETs (Figure 2.9) without significantly compromising the current on/off ratio of the device [14]. Similar results have been found for p-doped p-channel OFETs [19]. Another mechanism at play in this process is believed to be the effective lowering of the contact barriers, allowing carrier accumulation in the channel at lower gate voltages. Finally, whereas the overwhelming majority of OFETs currently operate in accumulation mode without any dopant present in the channel, doping has also enabled the realization of devices working in inversion mode. These devices rely on the insertion of a thin (few nm) layer doped with one type of dopant, e.g., n-type, between the undoped organic channel and the gate dielectric, and a thicker layer doped with the opposite type of dopant at the injecting contact. These transistors display highly reproducible characteristics such as Vth, current on/off ratio, sub-threshold swing, and carrier mobility [18]. The devices made possible by this approach can potentially be used in the realization of organic circuits.

FIGURE 2.9  Reduction in threshold voltage for OFET devices made with as-received and purified C60, as a function of doping ratio. The lines are guides for the eye. Reprinted from Olthof et al., Appl. Phys. Lett. 101: 253303/1 (2012), with the permission of AIP Publishing.


Conjugated Polymers

2.5.3 OPVs Doping in OPV cells has not yet played the critical roles it has in other organic devices. So far, the most notable applications have been the introduction of doped extraction layers to circumvent the use of high- or low-work function-contact materials that have proven somewhat problematic, and to ensure the presence of a large built-in voltage [156]. One example on the hole-extraction side of a polymer-based bulk heterojunction cell is the replacement of PEDOT:PSS, which is widely used as a high-work-function (~5.0 eV) hole-extraction layer, with a thin (ca. 20 nm) heavily p-doped layer of the donor polymer, creating a low barrier contact [27]. Device characteristics show nearly equivalent performance, thus greatly widening the choice of materials for these structures. Similar attempts have been made on the electronextraction side, with thin C60 layers n-doped with CoCp*2 placed between the active layer and a Ag contact [23]. Devices in which both hole- and electron-extraction layers are p- ad n-doped, respectively, fall in the general category of p-i-n solar cells; for example, F4-TCNQ-doped MeO-TPD and C60 doped with rhodamine B have been used as hole- and electron-injection layers, respectively, for evaporated ZnPc:C60 OPVs [22]. A particularly simple approach to solution-processed p-i-n solar cells has recently been reported; spin-coating of a solution containing ethoxylated poly(ethyleneimine), P3HT, and C60 bis(indene) adduct (ICBA) onto a bottom electrode, followed by immersion in H3PMo12O40 solution, and finally top-electrode deposition, results in a P3HT:ICBA bulk-heterojunction device n- and p-doped in the vicinity of the bottom and top electrodes, respectively [30]. Doping of the active layer has been performed with moderate success: small amounts of F4TCNQ have been used to p-dope polymer:fullerene bulk-heterojunctions [25]. For example doping of PCPDTBT:PCBM cells with a molar ratio of ca. 0.5% [25, 26] resulted in small increases in the shortcircuit current, the open-circuit voltage, and the power-conversion efficiency, which were attributed to improved collection of photogenerated carriers, likely due to the dopant-induced filling of hole traps.

2.6 Summary Although challenges remain – especially regarding obtaining spatially stable doping in the highest IE hole-transport materials and lowest-EA electron-transport materials – considerable recent advances have been made in the electrical p- and n-doping of organic semiconductors using organic or metalorganic molecular oxidants and reductants, respectively. A wide range of dopants have been identified and/or developed, exhibiting a range of doping strengths, varied processibility, different dopant-ion shapes and sizes, and, in some cases, excellent chemical and diffusional stability. Although many studies rely on codepositing dopant and semiconductor from the gas phase or on spin-coating of dopant:semiconductor solutions, increasingly, different methods of dopant processing are also being explored. Furthermore, doped organic semiconductors are now increasingly being used in a wide variety of organic and organic-inorganic hybrid electronic devices and this trend seems set to continue for the foreseeable future.

Acknowledgments Work on molecular doping at Georgia Tech and Princeton was partly supported by the National Science Foundation (DMR-1305247, DMR-1506097, and DMR-1807797).

References 1. H. Shirikawa, E. J. Louis, A. G. MacDiarmid, C. K. Chiang and A. J. Heeger, J. Chem. Soc., Chem. Commun. 16: 578 (1977). 2. F. Jonas, W. Krafft and B. Muys, Macromol. Symp. 100: 169 (1995). 3. C. Ganzorig and M. Fujihira, Appl. Phys. Lett. 77: 4211 (2000).

Electrical Doping of Organic Semiconductors


4. X. Zhou, M. Pfeiffer, J. Blochwitz, A. Werner, A. Nollau, T. Fritz and K. Leo, Appl. Phys. Lett. 78: 410 (2001). 5. J. Endo, T. Matsumoto and J. Kido, Jpn. J. Appl. Phys. 41: L358 (2002). 6. K. Luan, T. Dao and J. Kido, J. Photopolym. Sci. Tech. 15: 261 (2002). 7. Y.-J. Yu, O. Solomeshch, H. Chechik, A. A. Goryunkov, R. F. Tuktarov, D. H. Choi, J.-I. Jin, Y. Eichen and N. Tessler, J. Appl. Phys. 104: 124505 (2008). 8. P. K. Koech, A. B. Padmaperuma, L. Wang, J. S. Swensen, E. Polikarpov, J. T. Darsell, J. E. Rainbolt and D. J. Gaspar, Chem. Mater. 22: 3926 (2010). 9. I. Bruder, S. Watanabe, J. Qu, I. B. Mueller, R. Kopecek, J. Hwang, J. Weis and N. Langer, Org. Electron. 11: 589 (2010). 10. X. Lin, B. Wegner, K. M. Lee, M. A. Fusella, F. Zhang, K. Moudgil, B. P. Rand, S. Barlow, S. R. Marder, N. Koch and A. Kahn, Nat. Mater. 16: 1209 (2017). 11. M.-C. Jung, H. Kojima, I. Matsumura, H. Benten and M. Nakamura, Org. Electron. 52: 17 (2018). 12. S. P. Tiwari, W. J. Potscavage, T. Sajoto, S. Barlow, S. R. Marder and B. Kippelen, Org. Electron. 11: 860 (2010). 13. P. Wei, J. H. Oh, G. Dong and Z. Bao, J. Am. Chem. Soc. 132: 8852 (2010). 14. S. Olthof, S. Singh, S. K. Mohapatra, S. Barlow, S. R. Marder, B. Kippelen and A. Kahn, Appl. Phys. Lett. 101: 253303 (2012). 15. S. Singh, S. K. Mohapatra, A. Sharma, C. Fuentes-Hernandez, S. Barlow, S. R. Marder and B. Kippelen, Appl. Phys. Lett. 102: 153303 (2013). 16. B. D. Naab, S. Himmelberger, Y. Diao, K. Vandewal, P. Wei, B. Lussem, A. Salleo and Z. Bao, Adv. Mater. 25: 4663 (2013). 17. I. D. V. Ingram, D. J. Tate, A. V. S. Parry, R. S. Sprick and M. L. Turner, Appl. Phys. Lett. 104: 153304 (2014). 18. B. Lüssem, M. L. Tietze, C. H. Hans Kleemann1, J. W. Bartha, A. Zakhidov and K. Leo, Nat. Commun. 4: 2775 (2013). 19. J. Belasco, S. K. Mohapatra, Y. Zhang, S. Barlow, S. R. Marder and A. Kahn, Appl. Phys. Lett. 105: 063301 (2014). 20. I. Isakov, A. F. Paterson, O. Solomeshch, N. Tessler, Q. Zhang, J. Li, X. Zhang, Z. Fei, M. Heeney and T. D. Anthopoulos, Appl. Phys. Lett. 109: 263301 (2016). 21. A. Al-Shadeedi, S. Liu, C.-M. Keum, D. Kasemann, C. Hoßbach, J. Bartha, S. D. Bunge and B. Lu s̈ sem, ACS Appl. Mater. Interf. 8: 32432 (2016). 22. B. Maennig, J. Dreschel, D. Gebeyehu, P. Simon, F. Kozlowski, A. Werner, A. Li, S. Grundmann, S. Sonntag, M. Koch, K. Leo, M. Pfeiffer, H. Hoppe, D. Meissner, N. S. Sariciftci, I. Riedel, V. Dyakonov and J. Parisi, Appl. Phys. A 79: 1 (2004). 23. C. K. Chan, W. Zhao, A. Kahn and I. G. Hill, Appl. Phys. Lett. 94: 203306 (2009). 24. Y. Sun, S.-C. Chien, H.-L. Yip, Y. Zhang, K.-S. Chen, D. F. Zeigler, F.-C. Chen, B. Lin and A. K.-Y. Jen, Chem. Mater. 23: 5006 (2011). 25. A. V. Tunc, A. D. Sio, D. Riedel, F. Deschler, E. Da Como, J. Parisi and E. von Hauff, Org. Electron. 13: 290 (2012). 26. Y. Zhang, H. Zhou, J. Seifter, L. Ying, A. Mikhailovsky, A. J. Heeger, G. C. Bazan and T.-Q. Nguyen, Adv. Mater. 25: 7038 (2013). 27. A. Dai, Y. Zhou, A. L. Shu, S. K. Mohapatra, H. Wang, C. Fuentes-Hernandez, Y. Zhang, S. Barlow, Y.-L. Loo, S. R. Marder, B. Kippelen and A. Kahn, Adv. Funct. Mater. 24: 2197 (2014). 28. F. Selzer, C. Falkenberg, M. Hamburger, M. Baumgarten, K. Müllen, K. Leo and M. Riede, J. Appl. Phys. 115: 054515 (2014). 29. N. Aizawa, C. Fuentes-Hernandez, V. A. Kolesov, T. M. Khan, J. Kido and B. Kippelen, Chem. Commun. 52: 3825 (2016). 30. V. A. Kolesov, C. Fuentes-Hernandez, W.-F. Chou, N. Aizawa, F. A. Larrain, MingWang, A. Perrotta, S. Choi, S. Graham, G. C. Bazan, T.-Q. Nguyen, S. R. Marder and B. Kippelen, Nat. Mater. 16: 474 (2017).


Conjugated Polymers

31. Y. Xu, J. Yuan, J. Sun, Y. Zhang, X. Ling, H. Wu, G. Zhang, J. Chen, Y. Wang and W. Ma, ACS Appl. Mater. Interf. 10: 2776-2784 (2018). 32. T. Menke, D. Ray, J. Meiss, K. Leo and M. Riede, Appl. Phys. Lett. 100: 093304 (2012). 33. K. Shi, F. Zhang, C.-A. Di, T.-W. Yan, Y. Zou, X. Zhou, D. Zhu, J.-Y. Wang and J. Pei, J. Am. Chem. Soc. 137: 6979 (2015). 34. S. N. Patel, A. M. Glaudell, D. Kiefer and M. L. Chabinyc, ACS Macro Lett. 5: 268 (2016). 35. R. Kroon, J. D. Ryan, D. Kiefer, L. Yu, J. Hynynen, E. Olsson and C. Mu ̈ller, Adv. Funct. Mater. 27: 1704183 (2017). 36. D. Kiefer, A. Giovannitti, H. Sun, T. Biskup, A. Hofmann, M. Koopmans, C. Cendra, S. Weber, L. J. A. Koster, E. Olsson, J. Rivnay, S. Fabiano, I. McCulloch and C. Müller, ACS Energy Lett. 3: 278 (2018). 37. A. Abate, D. J. Hollman, J. l. Teuscher, S. Pathak, R. Avolio, G. D’Errico, G. Vitiello, S. Fantacci and H. J. Snaith, J. Am. Chem. Soc. 135: 13538 (2013). 38. Z. Wang, D. P. McMeekin, N. Sakai, S. van Reenen, K. Wojciechowski, J. B. Patel, M. B. Johnston and H. J. Snaith, Adv. Mater. 29: 1604186 (2017). 39. A. Pellaroque, N. K. Noel, S. N. Habisreutinger, Y. Zhang, S. Barlow, S. Marder and H. J. Snaith, ACS Energy Lett. 2: 2044 (2017). 40. T. Ye, J. Wang, W. Chen, Y. Yang and D. He, ACS Appl. Mater. Interf. 9: 17923 (2017). 41. J. Sworakowski, J. Lipinski and K. Janus, Org. Electron. 33: 300 (2016). 42. K. Harada, A. G. Werner, M. Pfeiffer, C. J. Bloom, C. M. Elliott and K. Leo, Phys. Rev. Lett. 94: 036601 (2005). 43. S. Guo, S. B. Kim, S. K. Mohapatra, Y. Qi, T. Sajoto, A. Kahn, S. R. Marder and S. Barlow, Adv. Mater. 24: 699 (2012). 44. T. P. Vaid, A. K. Lytton-Jean and B. C. Barnes, Chem. Mater. 15: 4292 (2003). 45. B. A. Gregg, S.-G. Chen and H. M. Branz, Appl. Phys. Lett. 84: 1707 (2004). 46. W. M. Porter, T. P. Vaid and A. L. Rheingold, J. Am. Chem. Soc. 127: 16559 (2005). 47. W. W. Porter and T. P. Vaid, J. Mater. Chem. 17: 469 (2007). 48. P. Bag, M. E. Itkis, D. Stekovic, S. K. Pal, F. S. Tham and R. C. Haddon, J. Am. Chem. Soc. 137: 10000 (2015). 49. S. Olthof, S. K. Mohapatra, S. Barlow, S. Mehraeen, V. Coropceanu, J.-L. Brédas, S. R. Marder and A. Kahn, Phys. Rev. Lett. 109: 176601 (2012). 50. X. Lin, G. E. Purdum, Y. Zhang, S. Barlow, S. R. Marder, Y.-L. Loo and A. Kahn, Chem. Mater. 28: 2677 (2016). 51. S.-J. Yoo and J.-J. Kim, Macromol. Rapid Commun. 36: 984 (2015). 52. F. Zhang and A. Kahn, Adv. Funct. Mater. 28: 1703780 (2018). 53. I. Salzmann, G. Heimel, S. Duhm, M. Oehzelt, P. Pingel, B. M. George, A. Schnegg, K. Lips, RalfPeter Blum, A. Vollmer and N. Koch, Phys. Rev. Lett. 108: 035502 (2012). 54. H. Méndez, G. Heimel, A. Opitz, K. Sauer, P. Barkowski, M. Oehzelt, J. Soeda, T. Okamoto, J. Takeya, J.-B. Arlin, J.-Y. Balandier, Y. Geerts, N. Koch and I. Salzmann, Angew. Chem. Int. Ed. 52: 7751 (2013). 55. I. Salzmann, G. Heimel, M. Oehzelt, S. Winkler and N. Koch, Acc. Chem. Res. 49: 370 (2016). 56. J. Li, C. Koshnick, S. O. Diallo, S. Ackling, D. M. Huang, I. E. Jacobs, T. F. Harrelson, K. Hong, G. Zhang, J. Beckett, M. Mascal and A. J. Moulé, Macromolecules 50: 5476 (2017). 57. G. Parthasarathy, C. Shen, A. Kahn S. R. Forrest, J. Appl. Phys. 89: 4986 (2001). 58. W. Gao and A. Kahn, Appl. Phys. Lett. 79: 4040 (2001). 59. J. Li, C. W. Rochester, I. E. Jacobs, S. Friedrich, P. Stroeve, M. Riede and A. J. Moule, ACS Appl. Mater. Interf. 7: 28420 (2015). 60. A. Dai, A. Wan, C. Magee, Y. Zhang, S. Barlow, S. R. Marder and A. Kahn, Org. Electron. 23: 151 (2015). 61. C. K. Chan, F. Amy, Q. Zhang, S. Barlow, S. R. Marder and A. Kahn, Chem. Phys. Lett. 431: 67 (2006).

Electrical Doping of Organic Semiconductors


62. J. Li, C. W. Rochester, I. E. Jacobs, E. W. Aasen, S. Friedrich, P. Stroeve and A. J. Moulé, Org. Electron. 33: 23 (2016). 63. P. Pingel, M. Arvind, L. Kölln, R. Steyrleuthner, F. Kraffert, J. Behrends, S. Janietz and D. Neher, Adv. Electron. Mater. 2: 1600204 (2016). 64. I. E. Jacobs, E. W. Aasen, J. L. Oliveira, T. N. Fonseca, J. D. Roehling, J. Li, G. Zhang, M. P. Augustine, M. Mascal and A. J. Moulé, J. Mater. Chem. C 4: 3454 (2016). 65. E. Perry, C.-Y. Chiu, K. Moudgil, R. Schlitz, C. Takacs, K. O’Hara, J. Labram, A. Glaudell, J. Sherman, S. Barlow, C. J. Hawker, S. R. Marder and M. L. Chabinyc, Chem. Mater. 29: 9742 (2017). 66. C. Y. Kao, B. Lee, L. S. Wielunski, M. Heeney, I. McCulloch, E. Garfunkel, L. C. Feldman and V. Podzorov, Adv. Funct. Mater 19: 1906 (2009). 67. R. A. Schlitz, F. G. Brunetti, A. M. Glaudell, P. L. Miller, M. A. Brady, C. J. Takacs, C. J. Hawker and M. L. Chabinyc, Adv. Mater. 26: 2825 (2014). 68. R. Kroon, D. Kiefer, D. Stegerer, L. Yu, M. Sommer and C. Müller, Adv. Mater. 29: 1700930 (2017). 69. Y. Shen, K. Diest, M. H. Wong, B. R. Hsieh, D. H. Dunlap and G. G. Malliaras, Phys. Rev. B 68: 081204 (2003). 70. W. H. Nguyen, C. D. Bailie, E. L. Unger and M. D. McGehee, J. Am. Chem. Soc. 136: 10996 (2014). 71. N. G. Connelly and W. E. Geiger, Chem. Rev. 96: 877 (1996). 72. P. Vanýsek, “Electrochemical Series”, in J. R. Rumble (Ed.), Handbook of Chemistry and Physics, 99th ed., CRC Press, Boca Raton, FL, 2018. 73. P. Coppo, R. Schroeder, M. Grell and M. L. Turner, Synth. Met. 143: 203 (2004). 74. D. M. Ivory, G. G. Miller, J. M. Sowa, L. W. Schacklette, R. R. Chance and R. H. Baughman, J. Chem. Phys. 71: 1506 (1979). 75. P. Kovacic and J. W. Timberlake, Polymer J. 20: 819 (1988). 76. J. Mort, S. Grammatica, D. J. Sandman and A. Troup, J. Electron. Mater. 9: 411 (1980). 77. A. Troup, J. Mort, S. Grammatica and D. L. Sandman, J. Non-Cryst. Solids 35: 151 (1980). 78. D. Hurum, B. Bovenzi, R. W. Kreilick and D. Weiss, J. Phys. Chem. B 102: 1071 (1998). 79. J. J. Andre, J. Simon, R. Even, B. Boudjema, G. Guillaud and M. Maitrot, Synth. Met. 18: 683 (1987). 80. J. Y. Liu, R. Zhang, G. Sauve, T. Kowalewski and R. D. McCullough, J. Am. Chem. Soc. 130: 13167 (2008). 81. A. Kolesnicenko, T. Malinauskas, E. Kasparavicius, R. Send, V. Gaidelis, V. Jankauskas, H. Wonneberger, I. Bruder and V. Getautis, Tetrahedron 71: 8162 (2015). 82. P. Schulz, E. Edri, S. Kirmayer, G. Hodes, D. Cahen and A. Kahn, Energy Environ. Sci. 7: 1377 (2014). 83. U. B. Cappel, T. Daeneke and U. Bach, Nano Lett. 12: 4925 (2012). 84. A. Abate, T. Leijtens, S. Pathak, J. Teuscher, R. Avolio, M. E. Errico, J. Kirkpatrik, J. M. Ball, P. Docampo, I. McPherson and H. J. Snaith, Phys.Chem. Chem. Phys. 15: 2572 (2013). 85. P. Zalar, M. Kuik, Z. B. Henson, C. Woellner, Y. Zhang, A. Sharenko, G. C. Bazan and T.-Q. Nguyen, Adv. Mater. 26: 724 (2014). 86. L. Tan, M. David Curtis and A. H. Francis, Chem. Mater. 15: 2272 (2003). 87. H. Li, M. E. DeCoster, R. M. Ireland, J. Song, P. E. Hopkins and H. E. Katz, J. Am. Chem. Soc. 139: 11149 (2017). 88. M. Kröger, S. Hamwi, J. Meyer, T. Riedl, W. Kowalsky and A. Kahn, Appl. Phys. Lett. 95: 123301 (2009). 89. J. Meyer, M. Kröger, S. Hamwi, F. Gnam, T. Riedl, W. Kowalsky and A. Kahn, Appl. Phys. Lett. 96: 193302 (2010). 90. K. Zilberberg, S. Trost, J. Meyer, A. Kahn, A. Behrendt, D. Lützenkirchen-Hecht, R. Frahm and T. Riedl, Adv. Funct. Mater. 21: 4776 (2011). 91. M. Kröger, S. Hamwi, J. Meyer, T. Riedl, W. Kowalsky and A. Kahn, Org. Electron. 10: 932 (2009). 92. J.-H. Lee, H.-M. Kim, K.-B. Kim and J.-J. Kim, Org. Electron. 12: 950 (2011).


Conjugated Polymers

93. J.-H. Lee, H.-M. Kim, K.-B. Kim, R. Kabe, P. Anzenbacher and J.-J. Kim, Appl. Phys. Lett. 98: 173303 (2011). 94. S. Zhang, X. Lu, J. Sun, Y. Zhao and X. Shao, Cryst. Eng. Comm. 17: 4110 (2015). 95. M. L. Tietze, L. Burtone, M. Riede, B. Lüssem and K. Leo, Phys. Rev. B 86: 035320 (2012). 96. Y. Karpov, T. Erdmann, M. Stamm, U. Lappan, O. Guskova, M. Malanin, I. Raguzin, T. Beryozkina, V. Bakulev, F. Günther, S. Gemming, G. Seifert, M. Hambsch, S. Mannsfeld, B. Voit and A. Kiriy, Macromolecules 50: 914 (2017). 97. J. E. Rainbolt, P. K. Koech, E. Polikarpov, J. S. Swensen, L. Cosimbescu, A. Von Ruden, L. Wang, L. S. Sapochak, A. B. Padmaperuma and D. J. Gaspar, J. Mater. Chem. C 1: 1876 (2013). 98. J. Li, G. Zhang, D. M. Holm, I. E. Jacobs, B. Yin, P. Stroeve, M. Mascal and A. J. Moule, Chem. Mater. 27: 5765 (2015). 99. Y. Karpov, T. Erdmann, I. Raguzin, M. Al-Hussein, M. Binner, U. Lappan, M. Stamm, K. L. Gerasimov, T. Beryozkina, V. Bakulev, D. V. Anokhin, D. A. Ivanov, F. Gu n ̈ ther, S. Gemming, G. Seifert, B. Voit, R. D. Pietro and A. Kiriy, Adv. Mater. 28: 6003 (2016). 100. N. Liu, Y. Morio, F. Okino, H. Touhara, O. V. Boltalina and V. K. Pavlovich, Synth. Met. 86: 2289 (1997). 101. S. K. Mohapatra, Y. Zhang, B. Sandhu, M. S. Fonari, T. V. Timofeeva, S. R. Marder and S. Barlow, Polyhedron 116: 88 (2016). 102. Y. Qi, T. Sajoto, S. Barlow, E.-G. Kim, J.-L. Brédas, S. R. Marder and A. Kahn, J. Am. Chem. Soc. 131: 12530 (2009). 103. F. Zhou, G. J. Van Berkel and J. B. T. Donovan, J. Am. Chem. Soc. 116: 5485 (1994). 104. R. Meerheim, S. Olthof, M. Hermenau, S. Scholz, A. Petrich, N. Tessler, O. Solomeshch, B. Lüssem, M. Riede and K. Leo, J. Appl. Phys. 109: 103102 (2011). 105. A. F. Paterson, N. D. Treat, W. Zhang, Z. Fei, G. Wyatt-Moon, H. Faber, G. Vourlias, P. A. Patsalas, O. Solomeshch, N. Tessler, M. Heeney and T. D. Anthopoulos, Adv. Mater. 28: 7791 (2016). 106. C. Lambert and G. Nöll, J. Am. Chem. Soc. 121: 8434 (1999). 107. O. L. Griffith, A. G. Jones, J. E. Anthony and D. L. Lichtenberger, J. Phys. Chem. C 114: 13838 (2010). 108. R. P. Veregin and J. R. Harbour, J. Phys. Chem. 94: 6231 (1990). 109. A. Yamamori, C. Adachi, T. Koyama and Y. Taniguchi, Appl. Phys. Lett. 72: 2147 (1998). 110. Y. Sato, T. Ogata and J. Kido, Proc. S.P.I.E., Int. Soc. Opt. Eng. 4105: 134 (2000). 111. J. Kido and T. Matsumoto, Appl. Phys. Lett. 73: 2866 (1998). 112. C. K. Chan, W. Zhao, S. Barlow, S. R. Marder and A. Kahn, Org. Electron. 9: 575 (2008). 113. H. Yan, Z.-H. Chen, Y. Zheng, C. Newman, J.-R. Quinn, F. Dötz, M. Kastler and A. Facchetti, Nature 457: 679 (2009). 114. Y. Qi, S. K. Mohapatra, S. B. Kim, S. Barlow, S. R. Marder, and A. Kahn, Appl. Phys. Lett., 100: 083305 (2012). 115. M. M. Alam and S. A. Jenekhe, J. Phys. Chem. B 106: 11172 (2002). 116. S. Wang, H. Sun, U. Ail, M. Vagin, P. O. Å. Persson, J. W. Andreasen, W. Thiel, M. Berggren, Xavier Crispin, D. Fazzi and S. Fabiano, Adv. Mater. 28: 10764 (2016). 117. D. M. Russell, T. Kugler, C. J. Newsome, S. P. Li, M. Ishida and T. Shimoda, Synth. Met. 156: 769 (2006). 118. C. K. Chan, A. Kahn, Q. Zhang, S. Barlow and S. R. Marder, J. Appl. Phys. 102: 014906 (2007). 119. C. S. Kim, S. Lee, L. L. Tinker, S. Bernhard and Y.-L. Loo, Chem. Mater. 21: 4583 (2009). 120. C. K. Chan and A. Kahn, Appl. Phys. A 95: 7 (2009). 121. M. L. Tietze, B. D. Rose, M. Schwarze, A. Fischer, S. Runge, J. Blochwitz-Nimoth, B. Lüssem, K. Leo and J.-L. Brédas, Adv. Funct. Mater. 26: 3730 (2016). 122. F. A. Cotton, N. E. Gruhn, J. Gu, P. Huang, D. L. Lichtenberger, C. A. Murillo, L. O. Van Dorn and C. C. Wilkinson, Science 298: 1971 (2002). 123. F. A. Cotton, J. P. Donahue, N. E. Gruhn, D. L. Lichtenberger, C. A. Murillo, D. J. Timmons, L. O. Van Dorn, D. Villagrán and X. Wang, Inorg. Chem. 45: 201 (2006).

Electrical Doping of Organic Semiconductors


124. C. Burkholder, W. R. Dolbier and M. Médebielle, J. Org. Chem. 63: 5385 (1998). 125. M. Fu, Y. Xiao, X. Qian, D. Zhao and Y. Xu, Chem. Commun. 15: 1780 (2008). 126. C. K. Chan, E.-G. Kim, J.-L. Brédas and A. Kahn, Adv. Funct. Mater. 16: 831 (2006). 127. S. K. Mohapatra, A. Fonari, C. Risko, K. Yesudas, K. Moudgil, J. H. Delcamp, T. V. Timofeeva, J.-L. Brédas, S. R. Marder and S. Barlow, Chem. Eur. J. 20: 15385 (2014). 128. D. E. Morris, K. W. Hanck and M. K. DeArmond, J. Electroanal. Chem. 149: 115 (1983). 129. S. Zhang, B. D. Naab, E. V. Jucov, S. Parkin, E. G. B. Evans, G. L. Millhauser, T. V. Timofeeva, C. Risko, J.-L. Brédas, Z. Bao, S. Barlow and S. R. Marder, Chem. Eur. J. 21: 10878 (2015). 130. B. D. Naab, S. Guo, S. Olthof, E. G. B. Evans, P. Wei, G. L. Millhauser, A. Kahn, S. Barlow, S. R. Marder and Z. Bao, J. Am. Chem. Soc. 135: 15018 (2013). 131. P. Wei, T. Menke, B. D. Naab, K. Leo, M. Riede and Z. Bao, J. Am. Chem. Soc. 134: 3999 (2012). 132. M. Schwarze, B. D. Naab, M. L. Tietze, R. Scholz, P. Pahner, F. Bussolotti, S. Kera, D. Kasemann, Z. Bao and K. Leo, ACS Appl. Mater. Interf. 10: 13401 (2018). 133. A. G. Werner, F. Li, K. Harada, M. Pfeiffer, T. Fritz and K. Leo, Appl. Phys. Lett. 82: 4495 (2003). 134. A. G. Werner, F. Li, K. Harada, M. Pfeiffer, T. Fritz, K. Leo and S. Machill, Adv. Funct. Mater. 14: 255 (2004). 135. D. Gebeyehu, B. Maennig, J. Dreschel, K. Leo and M. Pfeiffer, Sol. Energy Mater. Sol. Cells 79: 81 (2003). 136. S.-L. Suraru and F. Würthner, Angew. Chem. Int. Ed. 53: 7428 (2014). 137. G. Samit and S. Sourav, J. Am. Chem. Soc. 132: 17674 (2010). 138. C.-Z. Li, C.-C. Chueh, F. Ding, H.-L. Yip, P.-W. Liang, X. Li and A. K.-Y. Jen, Adv. Mater. 25: 4425 (2013). 139. C.-C. Chueh, C.-Z. Li, F. Ding, Z. Li, N. Cernetic, X. Li and A. K.-Y. Jen, Appl. Mater. Interf. 9: 1136 (2017). 140. G. Bélanger-Chabot, A. Ali and F. P. Gabbaï, Angew. Chem. Int. Ed. 56: 9958 (2017). 141. C. D. Weber, C. Bradley and M. C. Lonergan, J. Mater. Chem. A 2: 303 (2014). 142. C. Huang, S. Barlow and S. R. Marder, J. Org. Chem. 76: 2386 (2011). 143. M. J. Tauber, R. F. Kelley, J. M. Giaimo, B. Rybtchinski and M. R. Wasielewski, J. Am. Chem. Soc. 128: 1782 (2006). 144. P. J. Smith and C. K. Mann, J. Am. Chem. Soc. 34: 1821 (1969). 145. B. Russ, M. J. Robb, B. C. Popere, E. E. Perry, C.-K. Mai, S. L. Fronk, S. N. Patel, T. E. Mates, G. C. Bazan, J. J. Urban, M. L. Chabinyc, C. J. Hawker and R. A. Segalman, Chem. Sci. 7: 1914 (2016). 146. Y. Matsunaga, K. Goto, K. Kubono, K. Sako and T. Shinmyozu, Chem. Eur. J. 20: 7309 (2014). 147. S. Fabiano, S. Braun, X. Liu, E. Weverberghs, P. Gerbaux, M. Fahlman, M. Berggren and X. Crispin, Adv. Mater. 26: 6000 (2014). 148. Y. Zhou, C. Fuentes-Hernandez, J. Shim, J. Meyer, A. J. Giordano, H. Li, P. Winget, T. Papadopoulos, H. Cheun, J. Kim, M. Fenoll, A. Dindar, W. Haske, E. Najafabadi, T. M. Khan, H. Sojoudi, S. Barlow, S. Graham, J.-L. Brédas, S. R. Marder, A. Kahn B. Kippelen, Science 336: 327 (2012). 149. F. Li, A. Werner, M. Pfeiffer and K. Leo, J. Phys. Chem. B 108: 17076 (2004). 150. J. H. Oh, P. Wei and Z. Bao, Appl. Phys. Lett. 97: 243305 (2010). 151. N. Cho, H.-L. Yip, J. A. Davies, P. D. Kazarinoff, D. F. Zeigler, M. M. Durban, Y. Segawa, K. M. O’Malley, C. K. Luscombe and A. K.-Y. Jen, Adv. Energy Mater. 1: 1148 (2011). 152. S. Rossbauer, C. Müller and T. D. Anthopoulos, Adv. Funct. Mater. 24: 7116 (2014). 153. M. Lu, H. T. Nicolai, G.-J. A. H. Wetzelaer and P. W. M. Blom, Appl. Phys. Lett. 99: 173302 (2011). 154. B. D. Naab, S. Zhang, K. Vandewal, A. Salleo, S. Barlow, S. Marder and Z. Bao, Adv. Mater. 26: 4268 (2014). 155. S. K. Mohapatra, A. Fonari, C. Risko, K. Yesudas, K. Moudgil, J. H. Delcamp, T. V. Timofeeva, J.-L. Brédas, S. R. Marder and S. Barlow, Chem. Eur. J., 20, 15385 (2014). 156. R. Steim, F. R. Kogler and C. J. Brabec, J. Mater. Chem. 20: 2499 (2010).

3 Electric Transport Properties in PEDOT Thin Films 3.1 Introduction.........................................................................................46 3.2 Chemistry of PEDOT..........................................................................47



3.5 3.6


Nara Kim, Ioannis Petsagkourakis, Shangzhi Chen, Magnus Berggren, Xavier Crispin, Magnus P. Jonsson, and Igor Zozoulenko


Chemical vs. Electrochemical Polymerization of PEDOT:X • Chemical Water Dispersion: PEDOT:PSS • PEDOT:Biopolymer Dispersion Polymerization • Tuning the Oxidation/Doping Level Chemically vs. Electrochemically

Electronic Structure of PEDOT: From a Single Chain to a Thin Film............................................................................................... 55 Nature of Charge Carriers and Electronic Structure of PEDOT Chains • Density of States of PEDOT: From a Single Chain to a Thin Film • Band Gap and Optical Transitions in PEDOT

Morphology of PEDOT...................................................................... 61 Brief Review of Experimental Data for PEDOT:X and PEDOT:PSS (GIWAXS, TEM, AFM) • Morphology of PEDOT: A Theoretical Perspective

Electrical Conductivity.......................................................................67 Basic Thermodynamics of Thermoelectrical Processes • Temperature Dependence • Secondary Doping • Acid-Base Effect

Optical Conductivity.......................................................................... 81 Basic Definitions and Relations • Methodologies for Measuring the Dielectric Function • Optical Conductivity and Permittivity of PEDOT • Concluding Remarks on PEDOT Optical Conductivity

Transport Properties of PEDOT: A Theoretical Perspective........99 Basics of the Hopping Transport: Semi-Analytical Approach and Kinetic Monte Carlo • Boltzmann Approach to Conductivity Based on the Model of an Ideal Crystal • Multi-Scale Modelling Based on the Realistic Morphology

Mixed Electron-Ion Transport in PEDOT.................................... 107 Devices Utilizing Mixed Electron and Ion Conductivity  •  Experimental Results • Modelling of Mixed Electron-Ion Transport in PEDOT • Calculation of Ion Diffusion in PEDOT

3.9 Conclusions and Outlook..................................................................115 Acknowledgments..........................................................................................116 References........................................................................................................116



Conjugated Polymers

3.1 Introduction Poly(3,4-ethylenedioxythiophene) (PEDOT) is a bicyclic derivative of polythiophene that emerged on the scientific scene at the end of the 1980s.1 There are more than 10000 published manuscript reports on the synthesis, fundamentals, and applications of PEDOT, thus making this material the most ­studied and explored conducting polymer of all kinds today. When charge-compensated with molecular or polymeric anions, it can form a highly conducting solid (i.e. the highest value of 8800 S/cm up to date)2 that exhibits great stability in its positively charged (p-doped) state. Several synthesis protocols have been developed for the material to enable the formation of thin and thick films on a vast array of substrates and carriers, such as on large area flexible foils and as conformal claddings on fibers. Intrinsically, PEDOT is typically insoluble; however, when dispersion-polymerized with for instance poly(styrenesulfonate) (PSS) a processable emulsion is achieved.3 PEDOT:PSS has been further developed into various ink and coating formulations making the production of conducting patterns and areal electrodes possible using standard printing and coating techniques. Many of these techniques, successfully utilized today for PEDOT:PSS, were developed for the graphic art and industrial printing industry during the past centuries.4 In addition to high electrical conductivity, unprecedented air and thermal stability are attributed to the cyclic oxygen-bearing substituents that stabilize free radical and positive charge in a conjugated backbone (Figure 3.1).5 The electron-donating oxygen substituents also lead to a low band-gap of 1.5 eV, which is 0.5 eV lower than that of polythiophene, and thereby favorable electro-optical properties: the thin film of doped PEDOT is almost transparent in the visible region, while the undoped, neutral PEDOT is blue-black with an absorption maximum in the middle of visible range.6 The transparency of the high conducting state together with a Fermi level of about 4.5–4.9 eV makes the material suitable as a transparent hole-injecting/-accepting electrode in organic light emitting and photovoltaic devices, respectively. Besides a high electronic conductivity, many of the solid PEDOT formulations also exhibit a relatively high ion mobility, especially when the material is hydrated. There is a strong coupling between ionic and electronic charge carriers in PEDOT, which is manifested in the electrochemical modulation of color and conductivity. PEDOT possesses a low oxidation potential and a broad electrochemical window, which allow facile, long-term, and wide-range electrochemical switching.7 These phenomena define the principle of operation in, for instance, PEDOT-based paper displays and tunable windows, and in electrochemical transistors and switches, respectively. The coupling between ion exchange and charge accumulation has also been explored in various bioelectronic applications, then targeting to translate signals and to regulate functions across the technology–biology interface. Various bioelectronic sensors and actuators are today explored to record and regulate physiology and processes in cells and organs, in vitro and in vivo. Finally, due to its high electrical conductivity and low thermal conductivity, PEDOT is considered as potential material in thermoelectric generators; while the absence of spin in its highly conducting state

FIGURE 3.1  Chemical structures of (a) EDOT monomer, (b) oxidatively doped PEDOT, and (c) neutral PEDOT. Reproduced from Mitraka et al., Journal of Materials Chemistry A 2017, 5, 4404–4412 with permission of The Royal Society of Chemistry.

Charge Transport in PEDOT Thin Films


and low spin-orbit coupling promote long spin lifetime which is of interest in spintronic applications. To optimize the performance of the vast array of different PEDOT-based devices, currently being developed and explored, it is crucial to understand the material and its functionality, from the atomistic level all the way to its macroscopic dimensions. In this chapter, we summarize our present understanding of PEDOT, with respect to its chemical and physical fundamentals. By combining the results from microscopy, spectroscopy and crystallographic techniques with those from molecular simulations, we conclude upon the structure of several PEDOT systems, from the ångström level and up, and the impact on both electronic and ionic transport.

3.2 Chemistry of PEDOT 3.2.1 Chemical vs. Electrochemical Polymerization of PEDOT:X The conducting polymer, PEDOT, is the product of the polymerization of the 3,4-ethylenedioxythiophene (EDOT) monomers. The oxidant oxidizes EDOT, forming its cationic radicals, which consecutively form a dimer, releasing a proton in the polymerization media. Similarly, the dimers are being oxidized and polymerize in PEDOT (Figure 3.2).8–10 This basic principle applies in all polymerization methods of PEDOT.10 However, a different method can result in PEDOT with different counter-ions, (PEDOT:X), different morphologies and different properties. The most significant methods are the insitu chemical polymerization of PEDOT:X thin films (ICP), the vapor-phase chemical polymerization of PEDOT:X thin films (VPP), the chemical vapor deposition of PEDOT:X films (CVD), the electropolymerization of PEDOT:X films (EP) and the wet polymerization of PEDOT:X (WP), which forms dispersions of PEDOT inside a solvent. The ICP of PEDOT thin films was one of the first reported ways of producing PEDOT.11 Highly conducting PEDOT films can be synthesized by this method reaching values as high as 1000–2000 S/cm. This approach can produce PEDOT:X with various counter-ions like p-toluenesulfonate (i.e. tosylate, Tos–), chloride (Cl–) and trifluoromethanesulfonate (OTf–). Usually, the counter-ion originates from

FIGURE 3.2  The oxidative polymerization of the EDOT to PEDOT from iron(III) tosylate (Fe(Tos)3) (the Tos anion is referred as OTs in the figure). B is an organic base and Im is imidazole (an example of an organic base), while H+ are the free protons. Reproduced with permission from Ha et al., Advanced Functional Materials 2004, 14, 615–622, Copyright (2004) WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.


Conjugated Polymers

the oxidant complex, as the same oxidizing agent can also dope the system.11 Let’s have a closer look at the reactions going on during the polymerization. As shown in Figure 3.2 during the oxidative polymerization of EDOT, protons are being formed inside the polymerization media. Those free protons can further accelerate the rate of EDOT polymerization, as they can also oxidize EDOT. This is summarized in Eq. (3.1) for the polymerization rate,

2 2 r = k1CEDOT CFe ( III) + k2CH+ CEDOTCFe ( III) (3.1)

where r is the reaction rate, k1 and k2 the rate constants for the oxidation of EDOT by the catalyst and the accelerating influence of protons, respectively, and Cx the concentration of the different species. The rate of this oxidative polymerization can be reduced by adding a proton scavenger to the polymerization media, like pyridine, imidazole or N-methyl-2-pyrrolidone (NMP). In fact, as shown in Figure 3.2, those bases attract the free protons that are formed during the polymerization, thus reducing the CH+ part of Eq. (3.1). This reduction in the rate allows the connection of more oligomers of EDOT, forming PEDOT of higher molecular weight. From an electronic point of view, this is quite beneficial for the material properties, as charges transport along longer chains, forming 2D conducting networks.8, 11–12 As a result, films of high electrical conductivity are produced. The VPP polymerization is based on a similar concept with ICP, but instead of mixing EDOT to the oxidant solution, the oxidant films are exposed to EDOT vapors, as seen in Figure 3.3. The conductivity of VPP PEDOT:X can also be enhanced by the addition of a base, acting as a polymerization inhibitor like in ICP, or through a secondary dopant. The addition of a copolymer like poly(ethylene glycol)poly(propylene glycol)-poly(ethylene glycol) (PEG-PPG-PEG) to the oxidant enables a homogeneous distribution of small oxidant crystals. The copolymer makes the film waxy and acts as a template for growth, thus promoting good packing between PEDOT chains. Particularly, the presence of such a copolymer in the system is forcing a bottom-up synthesis of the PEDOT:X (Figure 3.3b).13–17 Additionally, the oxidant is oriented around this relatively long copolymer, allowing the polymerization of PEDOT films of higher thin film crystallinity (Figure 3.4i). In fact, it has been shown that the addition of longer copolymers can result in the production of more crystalline PEDOT:Tos (Figure 3.4ii). A higher crystallinity is beneficial for the electrical conductivity of the conducting polymer, as the charge

FIGURE 3.3  The steps for the VPP process. First (a) the oxidant solution is deposited on the substrate. Then, (b) the oxidant films are exposed to monomer vapor at a given temperature (T) and pressure (P), where ­oxidant/ monomer is transported/condensed at the interface to initiate polymerization. Finally, (c) the excess oxidant and monomer are washed away by bathing the films in a solvent, like ethanol, to yield an VPP thin film. Reprinted from Brooke et al., Progress in Materials Science 2017, 86, 127–146, Copyright (2017), with permission from Elsevier.

Charge Transport in PEDOT Thin Films


FIGURE 3.4  (i) The iron catalyst is oriented around the block copolymer (left), which acts as a template resulting in more ordered PEDOT (right). (ii) XRD diffraction patterns of VPP PEDOT. The peak observed at 2θ = 7.38° can be used to qualitatively compare the crystallinity of the films. The film with no block copolymer appears to be the less crystalline (as the (100) peak has the lowest intensity, red), with the degree of crystallization to be increased with the addition of the copolymer in (blue) Mw = 2900 g/mol of the copolymer and (green) Mw = 5800 g/mol of the copolymer with ethanol as the solvent. AFM images and schematic are included to highlight the change in the thin film morphology. Reprinted with permission from Fabretto et al., Chemistry of Materials 2012, 24, 3998– 4003, Copyright (2012) American Chemical Society. (iii) Schematic Images of the VPP PEDOT:Tos with different nanostructures of PEO-b-PPO-b-PEO block copolymer/iron(III) tosylate (Fe(Tos)3) mixture films from various evaporation-induced self-assembly (EISA). Reprinted with permission from Lee et al., ACS Macro Letters 2017, 6, 386–392, Copyright (2017) American Chemical Society.

transport is being facilitated between the different chains of the film. Thus, PEDOT:Tos films of higher electrical conductivity (>2000 S/cm) can be produced.16–17 Furthermore, the addition of copolymers with specific orientations (plates or tubes) can induce the formation of PEDOT films with controlled structure or porosity, as shown by Lee et al. (Figure 3.4iii).18 A recent approach to the PEDOT polymerization is the chemical vapor deposited PEDOT:X. EDOT is the precursor molecule and the oxidant, like an iron salt, is the catalyst. The iron salt is sublimed at a high temperature and its vapors are exposed to the vapors of EDOT in the chamber, then, they react and form the product inside the reaction chamber (Figure 3.5i). Specifically, one of the advantages of this process is that one can tune the orientation and crystallinity of the PEDOT films, by pre-patterning the substrates with sputtering or a block copolymer. As a result, the PEDOT films of electrical conductivity as high as 2000 S/cm can be produced. The other advantage of such technique is the formation of copolymers of EDOT with other molecules, e.g. anthracene. Although the conductivity of such films is relatively poor (~20 S/cm), their transmittance (~93%) is much higher than that of PEDOT (~80%), allowing their use in optoelectronic devices.19–23 The three aforementioned methods, ICP, VPP, and CVD, involved the oxidation of the EDOT monomer from an oxidant, like iron tosylate or iron chloride, forming PEDOT:Tos or PEDOT:Cl respectively. Those techniques though, restrict the choices of the counter-ion of the PEDOT, as the oxidant acts also


Conjugated Polymers

FIGURE 3.5  (i) A schematic depicting the CVD facility for the polymerization of PEDOT. Reprinted from Im et al., Applied Physics Letters 2007, 90, 152112 with the permission of AIP Publishing. (ii) Electropolymerization of PEDOT films. The working electrode, Au (WE), the reference electrode, Ag/AgCl (RE), and the counter electrode, Pt (CE), are bathed inside the electrolyte solution (i.e. KCl in acetonitrile) that has dispersed the EDOT monomer (dot). Upon application of voltage, the EDOT is polymerized on the working electrode (right). Afterwards the working electrode is removed from the electrolyte and the gold is removed by washing the electrode in Aqua Regia (HNO3/HCl mixture). (iii) The conformation of the chains of electropolymerized PEDOT is different with respect to the different counter-ions, ClO4–, PF6– and bis(trifluoromethylsulfonyl)imide (BTFMSI). Reproduced from Culebras et al., Journal of Materials Chemistry A 2014, 2, 10109–10115 with permission of The Royal Society of Chemistry.

as the dopant. This obstacle can be overcome in a fourth approach, the electropolymerization. The electrochemical setup for the EP is a three-electrode setup with a working electrode (i.e. Au), a reference electrode (i.e. Ag/AgCl) and a counter electrode (i.e. Pt). This three-electrode configuration is placed inside a solution of EDOT with electrolyte (i.e. KCl or an ionic liquid) and a solvent (i.e. acetonitrile), as shown in Figure 3.5ii. By applying a voltage to the working electrode, the EDOT is oxidized and an insoluble PEDOT film is formed on its surface. The anion of the salt used as electrolyte becomes the counter-ion of the doped PEDOT (Table 3.1). The metal of the working electrode can be removed by a washing post-treatment with Aqua Regia. The proper choice of a counter-ion can result in PEDOT of relatively high electrical conductivity (~1000 S/cm) as a long counter-ions can induce a delocalization of the charges in the PEDOT chain (Figure 3.5iii).11, 24–25

3.2.2 Chemical Water Dispersion: PEDOT:PSS All methods presented up to now hold a main disadvantage: they lead to an insoluble PEDOT layer. Therefore it is not possible to have inks of PEDOT which can facilitate both large-scale coating or patterning through printing technologies.10 The fifth method, the dispersion polymerization, enables solution processability through the formation of a dispersion of nanoparticles of PEDOT balanced with a polyanion poly(4-styrenesulfonate), PEDOT:PSS. PEDOT:PSS was initially developed and commercialized by Bayer AG in order to provide conductive coatings for anti-electrostatic agents.26 Its easy processability from an environmentally friendly solvent, i.e. water, and good film forming properties made PEDOT:PSS technically much more useful.3 The negative charges on the sulfonate groups (SO3–) in PSS neutralize the positive charges in doped PEDOT. Meanwhile, the surface charges provided by extra sulfonic acid groups (SO3–H+) stabilize the aqueous colloidal dispersion of PEDOT:PSS complex through the Coulomb repulsion between colloidal


Charge Transport in PEDOT Thin Films TABLE 3.1  The electrical conductivities of PEDOT derivatives obtained through electropolymerization with various counter-ions Method



Counterion ClO4 BF4− PF6− NO3− SO4− Tos− N(SO2CF3)2− N(SO2C2F5)2− N(SO2C3F7)2− N(SO2C4F9)2− CF3SO3− C4F9SO3− C8F17SO3− PSS−








Conductivity at RT (S/cm)


30 200

400–650 280 120–150

55 24 49 79 41 36




30 121

125–450 130 108 106 25 86 91 8




Source: Adapted from Ref. 11

particles.27 Thus, the non-stoichiometric, excessive amount of PSS is necessary for a stable dispersion of PEDOT:PSS. The molar ratio of EDOT units to styrene sulfonate units in a highly conductive grade of commercial PEDOT:PSS dispersions (e.g. CleviosTM PH 500 and PH 1000) is 1:1.9, corresponding to a weight ratio of 1:2.5 (Table 3.2).27 Since the doping level of PEDOT is known to be approximately 1 charge per 3 EDOT units,28 the charge excess of PSS is about 6-fold. The synthesis of PEDOT:PSS is performed by oxidative polymerization of EDOT monomers in the presence of PSS in water. Since PSS does not have an oxidative effect, a distinct oxidizing agent is needed TABLE 3.2  Various grades of commercial PEDOT:PSS dispersions provided by Heraeus Deutschland GmbH Trade Name CleviosTM P CleviosTM P VP AI 4083 CleviosTM P VP CH 8000 CleviosTM PH 500 Clevios™ PH 1000

Solids Content in Water (w/w) (%)

PEDOT:PSS Ratio (w/w)

Particle Size d50 (nm)

Conductivity (S/cm)

1.2–1.4 1.3–1.7

1:2.5 1:6

80a 80a

BZFS. A known mechanism in this group is the so called “∆g mechanism” in which the constituents of the PP have slightly different g factors. In the Δg mechanism the Sz = 0 (zero spin component along the field direction) levels of, say, the spin ½ pair are split by the Zeeman interaction that leads to faster S « T0 spin mixing process at higher fields (Devir-Wolfman et al. 2014). The MFE response in this case is not limited by BZFS, but by BZFS/Δg, in the case of long decay time τd, or by ħ/(µBΔgτd), in the case of very short τd (Devir-Wolfman et al. 2014, Khachatryan, Devir-Wolfman, et al. 2016). As an example for this mechanism we mention charge transfer excitons (CTE) having τd  8 T) magneto-photo-conductance (MPC) response (Devir-Wolfman et al. 2014). Another mechanism that belongs to the second group is “thermal spin polarization” that is caused by thermal population of spin sub-levels. This mechanism is more effective at low temperatures and high magnetic fields, as long as the species decay time is not much shorter than τSL, the characteristic time it takes to reach thermal equilibrium (Wang et al. 2012, Khachatryan, Devir-Wolfman, et al. 2016). It is important to note that having spin dependent pair dissociation and/or recombination is critical for all of the above mechanisms; experimentally, spin dependent decay was verified for a number of π-conjugated polymers (Bayliss et al. 2015, Kavand et al. 2016). MFEs have been observed not only in organic devices but also in contactless OS films that are not connected by wires to any voltage or current source, both in steady state (Gautam et al. 2012) and in the pico-second (ps) time domain spectroscopy (Huynh et al. 2017). Magnetic field effect of spectrally resolved photo-induced absorption (PA) and photoluminescence (PL) [named hereafter MPA and MPL, respectively] in OS made of π-conjugated polymer films (Nguyen et al. 2008), were applied to the study of spin dependent processes. This ‘spectroscopic-sensitive’ magnetic field effect technique differs from the previously studied ‘transport-related’ MC and MEL in devices in two important aspects. (i) Since PA and PL measure directly the density of the photoexcitations (such as PP and/or TE), then MPA and MPL can be directly related to the photoexcitation spin density. Consequently, by directly comparing the MPA and MPL responses in films to those of MC and MEL in organic diodes based on the same organic active layer, the magnetic field effect in organic diodes can be related to the spin densities of the excitations formed in the device. (ii) Being a spectroscopic technique, steady state MPA can be used as a new tool to discern various long-lived photoexcitations in OS films. In addition, studies of the main spin dependent species and/or spin-mixing mechanisms that determine the MPA (and MPL) response in various forms of OS, including spin-mixing in PP species, triplet–triplet annihilation, spin-mixing among the triplet spin sublevel, and Δg mechanism of PP in polymer/fullerene blends, can be performed. Transient-MPA (t-MPA) differs from steady state MPA (ss-MPA) in that it can distinguish between fast and slow spin dependent mechanisms. For example, with ps resolved t-MPA it is possible to observe the response MFE(t,B) of photoexcitation species whose spin dependent decay time, τ, is less than, say, ~0.1 ns; species that are impossible to detect under steady state conditions. Recently, strong resonant coupling between transient photoexcited SE and TT in donor–acceptor copolymers (SE decay time is ~0.2 ns) was established via the dependence of ps-resolved t-MPA(t,B) on delay time, t, and magnetic field, B (Huynh et al. 2017).

Magnetic Field Effects in Organic Semiconductors


8.2 Review of Various Mechanisms 8.2.1 The Hyperfine Mechanism We describe here a simple approach based on the time evolution of the PP spin sublevel populations in a magnetic field which is closely related to the well-known “radical-pair” mechanism (Hayashi 2004, Morozov and Dolctorov 1991, Steiner and Ulrich 1989, Timmel et al. 1998, Verkhovlyuk et al. 2006). This model can also explain MFEs in unipolar devices, where the two spin ½ paired-polarons have the same sign charges; in this case the pair of spins form a π-dimer (i.e. bi-radical, or bipolaron (Bobbert et al. 2007, Wang, Bassler, and Vardeny 2008)). We include in the spin Hamiltonian, H0, the Zeeman, HF, and exchange terms: H 0 = H Zeeman + H HF + H ex (8.1)

In Equation (8.1) H HF is the HF interaction term in which for simplicity we include the isotropic interaction of each spin with only a single nucleus of spin I = ½ (e.g. protons),     H HF = aS1 ⋅ I1 + aS2 ⋅ I 2 , (8.2) where a is the isotropic HF interaction strength. For protons in organic molecules, the HFI constant is of the order of a(H)~0.3 µeV (or a/µB~5 mT) (Carrington and McLachlan 1967). The expression given in Equation (8.2) can be extended to the more real case of many nuclei. The next term is the electronic Zeeman interaction term:    H Zeeman = µ B ( g 1S1 + g 2S2 ) ⋅ B (8.3)

gi(~2) is the respective g-factor of each of the polarons in the PP; and µB is the Bohr magneton. Another term is the isotropic exchange interaction:   H ex = −J ex (S1 ⋅ S2 + 1 / 4) (8.4)

Here we choose for simplicity scalar g-factors and scalar exchange interaction. The energy levels and wave functions can be easily obtained by solving the Schrodinger equation in this M = (2S1 + 1) (2S2 + 1) (2I1 + 1) (2I2 + 1) = 16 (with S1 = S2=½, I1 = I2 = ½) configuration space. In the absence of the HF interaction and for B = 0, the energy levels are split into a singlet (S = 0, ES = Jex/2) and three fold degenerate triplet (S = 1, ET = −Jex/2) levels separated by an energy Jex (“zero field splitting”, ZFS). When we include the HF interaction, the ZFS energies depend now on both a and Jex. Labeling the states by |S1zS2zI1zI2z> we can still define (electronic) spin singlet, |S>, and spin triplet |T0>, |T+>, |T-> states by the following expressions (with two identical nuclei of I = ½), | S> 

1 2 2

| T0 > 

1 2

2 I 1 m1 , m2 1


(|   m1m2 > |   m1m2 >)

2 I 1 m1 , m2 1

(|   m1m2 > |   m1m2 >) (8.5)

2 I 1 1 |  m1m2 > m1 , m2 1 2 where the ± signs on the right (left) hand side of Equation (8.5) signify the |±½> (|±1>) spin states of polaron spin S = ½ (triplet spin S = 1). The M levels contain now various weights of the above spin configurations and these weights are magnetic field dependent; thus the spin configuration character of the various levels is field dependent. Imposing spin dependent dissociation rates, it is easy to understand why there is MC(B): the density of the dissociated free carriers is field dependent. Such a simple

| T > 


Conjugated Polymers

approach can explain essentially most of the experimental findings of MFE in OS devices and films, as will become clearer below. The coherent evolution of the spin pair is given by the time dependence of the spin density operator σ(t) obtained from the Liouville von Neumann equation, as

σ(t ) = exp(−iH 0t / )σ0 exp(iH 0 t / ), (8.6)

where for the Hermitian H0 we have H0 = H0†, and σ0 is determined by the initial spin configuration of the system, σ0 =


pαP α (8.7)

where α = S,T0,T± , P α(pα > 0) is the projection operator (relative weight) for the α spin configuration. The fraction ρα (t) of PP in the α spin configuration at time t is then

ρα(t ) = Tr[P ασ(t )] =



α Wnmt mn 0nm




mn= 1

where Wnm = iω nm = i(En − Em ) /  , En (n = 1,…,M) is the level energy, Omn designates the mn matrix element of the operator O; note that ρα (t) is a real quantity and both P and σ0 are Hermitian. Equation (8.8) describes the coherent evolution of spin configuration density in the absence of losses. In reality ρα (t) diminishes to zero with time. We approximate the decay process assuming that the level population decays with a rate given by γn =

∑κP α

α α nn


where κα is the decay rate for the α spin configuration. Under this assumption the exponent Wnmt in Equation (8.8) becomes Wnm = (iωnm-γnm), where γnm = γn+γm > 0. PP can contribute either to the current in the device through dissociation or to EL through fusion into SE. We denote the reaction rate at which PP that is in a spin configuration α disintegrates by Rα. Then the disintegrated density (“yield”) during a time interval dt is: dYα = Rαρα (t)dt, and in steady state

 ρ (t )dt  R  P σ ω γ  γ (8.10) The total steady state yield is therefore Y = ∑ Y and the magneto effect is defined as

Yαss  Rα






α mn 0nm

2 nm


2 nm

ss α

MY (B) = [Y (B) − Y (0)]/ Y (0) (8.11)

where Y is the measured quantity (e.g. conductance, EL, photo-conductance). For long decay times and small exchange, γnm ~10 mT. For a spin pair in complete thermal equilibrium, the fraction of pairs in each of the spin sublevels En(n = 1,..M; M = 4 for PP, M = 3 for TE) is

 bn ρth / n e


b j


, (8.16)

where bn = En/kBT and kB is the Boltzmann constant. Starting at t = 0 from any given configuration, ρn0 , the system evolves with time toward thermal equilibrium in an exponential manner, with a time constant dominated by the material spin relaxation time, τS; therefore, the thermal equilibrium fractions (Equation (8.16)) are reached after a time t >> τS. Neglecting any spin-mixing mechanism that may occur prior to thermal equilibrium, but taking into account the SP decay process, the time dependent level occupation density may be calculated using the following set of rate equations,

dρn / dt  γ nρn  ρn  

 ρ ρ  / τ i


th n


(n  1..M ). (8.17)

In Equation (8.17) the first term on the RHS describes the spin dependent decay (the level decay rate, γn, is given by Equation (8.9) above), whereas the second term in the RHS describes the evolution toward thermal equilibrium. Consequently, the spin sublevel populations become magnetic field and time dependent. Solving the set of Eqs. (8.17) for given initial conditions (e.g. only the S configuration is initially photogenerated, or equal S and T configurations for charge injection) and magnetic field B, we obtain the solutions ρn(t,B). Spin pairs eventually dissociate (contributing to the current flowing through the device) or recombine (contributing to PL and EL). We denote the spin dependent rate of these reactions (i.e. dissociation or recombination) by Rα for the α spin configuration (α = S,Tx,Ty,Tz for PP and α = Tx,Ty,Tz for TE). Consequently, the contribution of thermal spin polarization (TSP) to the measured quantity, Yth, in steady state is proportional to the total reaction yield:


Magnetic Field Effects in Organic Semiconductors

 R ρ (t , B)dt , (8.18)

Yth (B)

where Rn 

 RP α

α α nn





is the level reaction rate. The contribution of TSP to the magnetic field effect,

MYth, is then obtained from Equation (8.18): MYth(B) = [Yth(B)/Yth(0)-1]. The calculated MYth is the contribution of thermal polarization to the derived MY (here MY is one of the following; MEL, MPL, MC, or MPC) in the absence of any spin-mixing mechanism. The total measured MY is thus the weighted sum of MYth and the contribution of the dynamic spin-mixing processes (e.g. the ∆g mechanism).

8.2.5 Magnetic Field Effect in Excited-State Spectroscopies of Films; Steady State Absorbed photons with above gap energy initially generate singlet excitons (SE) in OS materials. The SE may either radiatively recombine producing photoluminescence (PL); or convert into long-lived TE via intersystem crossing. These TEs are long-lived and can give rise to phosphorescence (PH), form TT pairs, or separate into positive and negative charge polarons, some of which may form long-lived PP. All these secondary species that cascade with time have excited state optical transitions (dubbed “photoinduced absorption” or PA), which is activated by a weak probe beam. A schematic diagram of the philosophy underlying the MPA technique that we have used in our studies is presented in Figure 8.1, where X0 → X 1 is the PA transition. PA is defined as the negative fractional change in transmission, T: PA(E) ≡ (−∆T/T) = NSSβ (E), where NSS is the species steady state density, β (E) is the photoexcitation optical cross-section, and E is the probe beam photon energy. Therefore, in a magnetic field, B, PA(B) is determined by the density NSS(B); which, in turn, is given by Nss = G/Κ where G is the species generation rate and Κ is the decay rate. For B ≠ 0, the X0 level Zeeman splits according to the relevant spin multiplicity. Consequently, when Κ is spin configuration dependent, the spin content of each Zeeman split level varies, causing NSS and consequently PA to become B-dependent, leading to the B dependent magneto-response MPA(B) ≡ [PA(B)-PA(0)]/PA(0). In contrast, since it originates from singlet exciton radiative recombination, MPL(B) cannot directly originate from SE (S = 0) (which is B-independent); but rather is caused indirectly, for example via delayed generation of SE from TT annihilation. Using Equation (8.10) and assuming that the optical cross section is spin independent, Rα≡R, we have for the PA yield X1 PAX (B)













FIGURE 8.1  Schematic illustration of the magnetic field dependent pump-probe PA processes. The pump beam with above gap photon energy hVL excites the OS to the singlet exciton (SE) level (S0→S1). The SE relaxes via intersystem crossing to a triplet exciton (TE) or ionizes into separate charges forming polaron pair, PP (S1→X0). The steady state density of the X species is controlled by the spin dependent decay coefficient, κ. The incandescent probe beam monitors the photoinduced absorption, PA (X0→X 1), which is proportional to the X0 steady state density. In a magnetic field B > 0, X0 splits according to its spin multiplicity, and the population of each level becomes field dependent, resulting in a B-dependent density and PA (thus forming MPA).


Conjugated Polymers

 ρ (t )dt  f  σ

YPA (B)  R





/ γ n (B) N ss (B) (8.19)

where f is a numerical factor (of order 1). Note that γn (defined in Equation (8.9)) is B dependent since the spin configuration of each level vary with B. Consequently, MPA(B) is determined by the B dependent steady state population.

8.2.6 Time Resolved Magnetic Field Effects In this section we consider the effects of magnetic field on short-lived species. Short-lived species may not be detectible in steady state; thus, the MFE discussed above that is measured in steady state conditions is not applicable for these transient photoexcitations. An example of short-lived species that nevertheless can show MPA are SE coupled to TT pairs (multi-spin configuration) via e.g. singlet fission processes. In low band gap co-polymers both SE and TT can be optically excited, and thus may be observed in time resolved PA (namely PA(t)). Under these conditions the effect of magnetic field on such fast photoexcitations may be measured using time resolved magneto-PA (t-MPA) (Huynh et al. 2017). A key ingredient for the t-MPA formation (as is for steady state MPA) is spin dependent decay rates for the spin configurations; this leads to field dependent level population and decay rates. PA is a spectroscopic measurement that enables wavelength separation of spin specific optical transitions. Because PA(t) is proportional to the specific excited state transient population, e.g. SE or TT, and these populations are field dependent, which generates time and field dependent t-MPA(t,B). The calculation of t-MPA follows directly from the density matrix approach discussed in Section 8.2. In the following, we demonstrate it for co-polymers in which near resonance SE and TT species that can be photogenerated together. We assume that there exists resonant spin-coupling among the lowest SE and the various spin-states of the lowest TT state (i.e. TT singlet (S), triplet (T), and quintet (Q)), which are responsible for the t-MPA(B) response. To describe the t-MPA we employ an appropriate spin-Hamiltonian in a 10 × 10 Hilbert space (namely SE and 9 TT spin sub-states). The mixed SE–TT spin Hamiltonian is comprised of the TT-pair Hamiltonian, the SE term, and coupling terms, V. Denoting the TT singlet, triplet, and quintet configurations by S, T, and Q, respectively, the SE–TT coupling terms Vα = (α = S,T,Q). The total [10 × 10] SE–TT spin-Hamiltonian matrix contains the usual [9 × 9] block of the TT pair Hamiltonian (Ehrenfreund and Vardeny 2012), [1 × 1] block for the SE term, and the V coupling along the 10th row and 10th column, at the appropriate positions dictated by the coupling. The 10 × 10 Hamiltonian matrix has the following form,

. . . . . . . . .   . . . . . . . . .  . . . . . . . . .   . . . . . . . . .  . . . . TT 9 × 9matrix. . . .   . . . . . . . . .  . . . . . . . . .   . . . . . . . . .  . . . . . . . . .   . − t s t − s t s . − t

.    − t   s    t    − s  , t    s    − t    .    ESE 


Magnetic Field Effects in Organic Semiconductors

MPA (%)






0 B (mT)


FIGURE 8.2  Illustration of transient magneto-PA (t-MPA) response in the ps time domain. Using the 10 × 10 model Hamiltonian discussed in the text, the calculated t-MPA(B) response (at 0.2 ns delay) of PASE (darker) and PATT (brighter) are shown. The parameters used are the same as those in Figure 8.10.

where the full TT 9 × 9 matrix including the Zeeman term for arbitrary direction is given in (Ehrenfreund and Vardeny 2012). The letters t = VT and s = VS are the respective coupling constant parameters of SE with the triplet and singlet configurations of the TT state. ESE is the SE energy relative to the TT median energy at zero magnetic field. To obtain the SE–TT MPA(B) response, the spin Hamiltonian was solved by adding the Zeeman interaction term based on the field B in various directions with respect to the triplet axes in the TT-pair state, which is in agreement with the disordered nature of the organic film. After finding the eigen-­f unctions and eigen-energies of this Hamiltonian, we have proceeded by calculating the weight of each of the ten spin configurations in each of the eigenstates. Importantly, at t = 0+ (immediately following pulsed photoexcitation) no decay has taken place yet; thus, although the spin content of each eigenstate varies with B, the overall spin density of each spin, when summing up over the ten eigenstates, does not depend on the field. We therefore conclude that t-MPA(B) = 0 at t = 0+, as is clearly seen in the experiment. However, at t > 0 the various levels decay at rates that are spin dependent. We then calculated the densities of the SE and triplet configurations vs B based on the density matrix approach. When the decay is uniform (i.e. spin independent), the total spin density decays as a whole and again, at a fixed delay time, t-MPA(B) = 0. In contrast, when the decay rates are spin dependent, the relative level densities change with time and field leading to t-MPA(B) ≠ 0. It is seen that whereas the TT-triplet relative density increases with B, the SE relative density shows an opposite trend. An example of the full t-MPA(B) response is depicted in Figure 8.2 that reveals the similar shape but opposite response of the TT-triplet and SE spin configurations, in agreement with the experiment (Huynh et al. 2017). Another key feature of the t-MPA studies is that SE is much shorter lived than TT. Thus beyond ~0.1 ns SE is practically gone while TT stays in the system up until ~0.1 µs. Eventually, TT disintegrates into two separate TEs that remain spin entangled before losing spin coherence.

8.3 Experimental Studies 8.3.1 Magnetic Field Effects in Organic Devices at Low Fields In this section we discuss studies of MC in fields up to ~50 mT, with special attention to very small fields (B    M e . For materials of lower molecular weight, comparable ∆ H f  values are obtained whether they were produced from the melt or solution (compare Figure 10.2a and 10.2b).8  It is important to note that

Establishing the Thermal Phase Behavior


FIGURE 10.2  (A) Differential scanning calorimetry, DSC, heating (top) and cooling (bottom) thermograms obtained for melt-solidified P3HT samples of a range of molecular weights, here given as M n  (heating/cooling rate  =  10° C/min). The supercooling, ∆ T  (i.e., the difference between its T m  and T c .) for P3HT of M n   =  130  kg/mol is also given. (B) Corresponding DSC first heating thermograms measured for solution-cast P3HT films. (Adapted from Koch et al., Prog. Polym. Sci.  2013, 38, 1978)


Conjugated Polymers

molecular order for such low-molecular weight materials does not necessarily increase. Despite the fact that, in low-molecular weight materials, the macromolecules are too short to entangle, end groups can hinder good packing, limiting molecular order independent of the dilution and/or deposition method selected.8 , 12  The onset of entanglement can also be deduced from the amount of supercooling, ∆ T , given by the difference between T m  and T c  (i.e., ∆ T   =  T m  –  T c ; see Figure 10.2.). A relatively large supercooling is required to at least partly crystallize a material of high molecular weight, because of the presence of entanglements. In contrast, lower-molecular weight materials, which do not entangle (or only do so in a limited fashion), have a higher molecular mobility. This enables material transport towards crystal nuclei, resulting in a notably less pronounced ∆ T.  A strong discrepancy between high- and low-molecular weight materials not only exists in the supercooling that they require to crystallize but also in their crystal growth rate, G (Figure 10.3.).12 –  14 , 18  Highermolecular weight materials generally crystallize at considerably lower rates than their low-molecular weight counterparts. If a polymer solidifies faster than it can crystallize, that is, when the solidification rate is higher than the crystal growth rate, the material vitrifies.13 –  14 , 18  This means a phase of a very limited degree of crystallinity— a glass— forms. This happens, for instance, when a polymer is quenched from the melt or when a solvent is very rapidly evaporated from a polymer solution during spin coating. This behavior is typical for many spin-cast polymeric semiconductors and conductors, especially for those with more complex chemical structures (e.g., structures used for the creation of electron donors for organic solar cells). It is important to note in this context that the crystal growth rate, G , which features a maximum value between T m and the glass transition temperature, T g , strongly depends on the selected crystallization temperature, T c (Figure 10.3.).20 –  25 A ‘ bell-shaped’ growth rate dependence with crystallization temperature, T c , is generally observed for polymeric species (Figure 10.3.),20 –  25 which results from an interplay between nucleation and growth, expressed as the rate of formation of primary and secondary nuclei: G  ∝  exp (−  ∆ G */kT ), and material transport (i.e., diffusion), given by G  ∝ exp (Δ Φ  /kT ).23  Thereby, Δ Φ is of the form: C1  ·  T /(C 2   +   T c  –  T g ), the so-called Williams– Landel– Ferry (WLF) approximation, where C1 and C2 are universal

FIGURE 10.3  Dependency of the crystallization rate, G , on crystallization temperature, T c , used during solidification. G  is maximum at crystallization temperatures halfway between the material’ s glass transition temperature, T g , and the material’ s melting temperature, T m .12 –  14 , 18 , 76  This bell-shaped dependency results from competing processes between nucleation (dominating at lower crystallization temperatures) and diffusion processes (dominating at higher T c ,). Note also that G  is higher for materials of lower molecular mass, M w .

Establishing the Thermal Phase Behavior


constants at temperatures up to 100  K above the glass transition temperature (above T g , an Arrhenius form is used).13 , 26 , 27 Accordingly, polymer crystallization can be classified into two categories: one is predominantly valid for melt-grown systems as well as solution-grown materials using relatively high molecular weight, in which the crystal growth rate G usually is nucleation-controlled and exponentially depends on 1/(T  ·  ∆ T ); the other is valid for solution-grown polymers of high dilution or materials of relative low molecular weight, in which the growth rate G is diffusion-controlled and linearly proportional to the supercooling ∆ T .28 , 29 This scenario can be manipulated by the use of additives such as nucleation agents, which affect the nucleation process and, thus, change the competition between diffusion and nucleation at a given temperature.13 , 30 Confinement can play an opposite role often drastically reducing the number of nuclei, simply because of a reduction in volume (as discussed in more detail in Section –  35 Impact of Chain Conformation Besides affecting the entanglement density in macromolecular structures, the chain conformation of a polymeric semiconductor can play an important role in dictating phase behavior. For example, the chain conformation will affect phase transitions such as the melting temperature. This is because T m  ≈  ∆ H f /∆ S f , with ∆ H f  being the enthalpy of fusion and ∆ S f  the entropy of fusion. The more rigid the polymer backbone is, the less drastically its conformation usually changes between the liquid and the solid state; that means that ∆ S f  is low and, in turn, T m  is high.36  Rod-like conformations frequently lead also to a rich phase behavior (cf. Ref.  7), including the occurrence of liquid-crystalline phases such as thermotropic phases (liquid-crystalline, LC, phases in the solid state) or lyotropic phases (LC phases in the liquid phase), as will be also discussed in Section 10.2.4. Illustrative examples in the semiconducting polymer area are poly(thiophene) derivatives that comprise the relatively planar (‘ rigid-rod-like’ ) thieno[3,2-b ]thiophene moiety in the polymer backbone (e.g., poly[​2,5-b​is(3-​tetra​decyl​t hiop​hen-2​-yl)t​hieno​[3,2-​b]thi​ophen​e], PBTTTs) or the light-emitting poly(9,9di-n -octylfluorenyl-2,7-diyl).6 , 7  For these polymeric semiconductors, the increased chain stiffness has been shown to lead to a thermotropic behavior where transitions into liquid-crystalline phases occur as temperature is changed. Realization of such a diversity of phases with different degrees of shortand long-range order offers many other opportunities; for instance, liquid crystallinity has assisted in organizing semiconducting macromolecules, including PBTTT, into large domains via solidification from such liquid-crystalline phases. In other fields, this has led to high-strength, high-modulus polymer fibers such as those made of aramids.37 , 38  As beneficial as LC phases can be, they are sometimes difficult to detect. It also is not always straightforward to distinguish between LC-phase transitions, solid– solid crystalline transformations, and melting/crystallization processes. As a rule of a thumb, one may distinguish between melting/crystallization and LC-mesophase transitions based on the fact that materials that do not feature mesophases often display a relatively pronounced ∆ T , especially for polymeric species, while systems with LC-mesophase transitions in many cases display a small ∆ T  (see Figure 10.4.)36  Finally, we would like to emphasize here that aside from the chemical structure of the polymer, the chain conformation that a polymer can adopt also strongly depends on the environment, including temperature, solution concentration, additives (small molecular and polymeric), and geometry (e.g., degree of confinement). Accordingly, via the type of processing methodologies and processing conditions selected, one can gain some control of the polymer semiconductor’ s conformation, and thus, its assembly and structure formation, which can be used to induce the desired properties and device performance. Confinement Confinement is another universal phenomenon that can change the phase behavior of polymer semiconductors and conductors.31 –  35  The classical argument to rationalize confinement effects is a scenario where the characteristic length scales of a certain physical process (crystallization, phase separation, conformational transitions, etc.) are of a similar order of magnitude as the space available for this process to occur. For instance, in ultra-thin-film structures (such as those often used in optoelectronic


Conjugated Polymers

FIGURE 10.4  Schematic differential scanning calorimetry heating/cooling thermograms (A) for a typical semicrystalline semiconducting polymer with a melting temperature, T m , a crystallization temperature (here denoted T x ) and a glass transition temperature, T g , and (B) for a material that displays crystalline, x , liquid-crystalline (e.g., nematic, n ), and isotropic, i  (i.e., melt) phases. Adapted from Snyder et al., Royal Soc. Chem., Cambridge, 2017.

devices, including organic solar cells, field-effect transistors, and light-emitting diodes), reorganization of polymer chains in order for the material to crystallize requires a volume that is equal to or larger than the volume provided by the films. In such a scenario, physical processes such as crystallization develop differently due to the size constraint compared to systems with no size limitation. This effect is clearly more notable the larger the molecules involved are because both nucleation and crystal growth processes in macromolecules usually occur over critical length scales in the 10– 100  nm regime. Confinement also affects the overall crystallization kinetics of polymers.33 , 35  On the one hand, this originates from the fact that in a confined volume there is only a very small space/volume available for crystals to grow in. On the other hand, just as importantly, the number of nucleation centers that develop in a confined volume is orders of magnitude higher than the density of nuclei found within the bulk of a polymer. As a result, the majority of the kinetic features that affect polymer crystallization in confinement will be governed by the nucleation stage. High-Temperature Crystallization, High-Pressure Crystallization, and Annealing When processing polymers, various experimental conditions other than those discussed above can significantly influence the shape, form, and degree of perfection of the polymer crystalline entities― and, notably, their lamellar crystal thickness, l  (Figure 10.1c). This can considerably affect the phase behavior. Here, we focus on three illustrative examples: crystallization in dilute solutions and at high temperatures, melt-crystallization at high pressures, and annealing.

Establishing the Thermal Phase Behavior


Crystallization at elevated temperatures from dilute solutions, which often causes significant ‘ disentanglement’  of the constituting macromolecules (i.e., a reduction of N ), can result in considerably increased crystal thicknesses.2 , 3 , 26 , 39  This behavior is dictated by the selection of T c . For example, it has been shown for solution-grown polyethylene, which is the most prominent commodity plastic, when T c  is selected such that it is close to T m , l  increases by a factor of two or more.40  Somewhat less effective is annealing the solid polymer, post-desposition, at temperatures near their melting temperature, where the polymer chains have a relatively high molecular mobility and, thus, can partly disentangle and become more ordered.41  While an increase of l  often is achieved by annealing, the extent depends on the processing and thermal history, as well as the number of entanglements present in the system. A dramatic increase of the crystal thickness and virtual elimination of chain folding can, in contrast, be achieved by crystallizing polymers under high pressures, as demonstrated on polymers ranging from commodity materials such as high-density polyethylene3 , 42 , 43  to polymer semiconductors such as P3HT.44  In many cases, this results in fully extended-chain crystals, much like those observed for oligomers or low-molecular weight polymers discussed above.45 

10.2.2 Glass Transition Beside the crystallization and melting temperature, the glass transition temperature, T g , is highly relevant when applying polymer semiconductors in devices.36  A number of structural parameters affect the glass transition temperature, which is the temperature below which segmental relaxations are suppressed. The empirical Flory– Fox equation1 , 46 , 47  describes, for example, the increase in T g  with molecular weight, M : T g (M )  =  T g (M ∞  ) –  K/M,  where T g  (M ∞  ) is the glass transition temperature of a polymer with infinite molecular weight (infinite chain length) and K is a polymer-specific constant.46 , 47  This equation implies that, very similar to the dependence of the melting temperature on molecular weight, for shorter macromolecules (low-molecular weight materials), the T g  first increases with molecular weight before reaching a plateau value for longer-chain polymers (higher-molecular weight materials) where T g  is essentially independent of chain length. Other important parameters that affect a polymer’ s T g  are backbone stiffness and the nature of the side-chains, including their length and the degree of branching. For instance, many semiconducting and conducting polymers feature relatively rigid backbones, which often is the result of the incorporation of fused aromatic moieties into the backbone structure. This leads to high glass transition temperatures, often in excess of 100° C.47  In contrast, materials such as poly(3-alkyl thiophene)s, P3ATs, generally display relatively low glass transition temperatures. This is because their backbones consist of thiophene-repeat units, which have a comparatively low energetic barrier for rotation. P3ATs also are good examples to illustrate the effect of side chain length on T g . When increasing the length of their side-chains from a linear alkyl chain comprising four carbon atoms to one that comprises twelve carbons while keeping them linear, i.e., when moving from a butyl-substituted P3AT to a dodecyl-substituted one, the glass transition temperatures of these P3ATs are reduced from around +40° C to − 15° C.47  Besides manipulating the glass transition temperature of a given material via the chemical design of the polymer, use of additives that reduce the T g  similarly to additives that lead to the above discussed melting-point depression can be exploited. Such additives are often referred to as ‘ plasticizers’ , This is of relevance for the deposition of blends for organic photovoltaic applications where ‘ plasticizers’  sometimes are utilized to induce a desired bulk heterojunction microstructure. The extent of molecular disorder that a material adopts, which dictates how pronounced the glass transition of a material is, can in many cases be controlled by the selected solidification parameters. This in particular applies to polymer semiconductors. Because of their often complex chemical structure, they can be forced into a state that is dominated by a glassy structure and, in extreme case, can lead to full vitrification (see Figure 10.5.). This can be achieved, for instance, by rapidly quenching the semiconductor below its glass transition temperature T g — w ith a considerable effect on the final structure. This is not always easy to control. For example, when materials are processed from solution, the solidification


Conjugated Polymers

FIGURE 10.5  Schematic overview of the phase separation processes that may occur during drying of a solution that contains a solvent and two solutes: component 1   and component 2   (e.g., the donor and the acceptor in polymer solar cell blends). These processes include liquid− liquid (L– L ), liquid− solid (L– S ) and solid− solid (S– S ) phase separations as well as disorder− order transitions (D– O ), both in the liquid (left) and solid state (right). From Kouijzer et al., J. Am. Chem. Soc.  2013, 135, 12057.

rate is dictated mostly by the solvent evaporation rate and, thus, is kinetically ill-defined. In such cases, using an additive— a vitrifier— t hat acts as a glass-inducing diluent can be beneficial. The introduction of this second component hinders the crystallization of the active species, leading to an amorphous architecture that is generally kinetically metastable. As a consequence, in a fabrication step following casting and drying, the initial glassy films can be crystallized in the solid state— often on demand. Depending on the phase behavior of the semiconductor and that of the vitrifier, various crystallization routes may be opened, permitting manipulation of parameters such as the nucleation density and the crystallization rate. Vitrifiers can be optoelectronically active materials, as the examples of P3HT:fullerene and other donor:acceptor blends targeted for use in organic photovoltaic cells show.19 , 47 , 48  In these donor:acceptor blends, vitrification results because of both kinetic and thermodynamic reasons.19 , 47 , 48  Kinetic effects are predominantly determined by the T g  of the individual components (in OPV devices: the donor and the acceptor), while their mutual miscibility, strongly influenced by their chemical nature and leading to thermodynamic mixing, dictates the second phenomena. It is important to note that for vitrified systems, annealing procedures above the glass transition temperature often allow induction of molecular ordering. In the extreme case, cold-crystallization occurs.36  In blends, treated in more detail in Section 10.3, the cold crystallization can result in phase separation (Figure 10.5.). While in neat systems, a temperature above the T g  of the material at hand can usually be selected as the annealing temperature, in multicomponent systems the glass transition temperature of the intermixed phase often provides the lower temperature limit for annealing.48  The reason for this is that in many donor:acceptor blends, the solid-state structure is dominated by a three-phase morphology comprised of an acceptor-rich phase, a donor-rich phase, and an intermixed phase that in many cases is amorphous, often because it was ‘ v itrified’ .19  In the case of spin-cast binary binaries of P3HT and phenyl-C61 -butyric acid methyl ester, PC61 BM, where blending followed by rapid solvent removal usually leads to predominantly vitrified films directly after casting, annealing above the T g  of the intermixed phase leads to cold crystallization of both components and, in turn, phase separation.48 –  50  These examples demonstrate that it can be critical when working with blends to determine the glass transition temperature of not only each neat material but also the intermixed phases.36 , 48 

Establishing the Thermal Phase Behavior


Finally, we would like to highlight here that glass transition temperatures are relevant parameters for selection of post-deposition procedures such as annealing (as mentioned above); moreover, they are in many other cases important to consider as they can have a pronounced effect on the thermal stability of the final device. This is illustrated by the example of solar cells comprising a donor polymer based on 5,6-difluorobenzothiadiazole (FBT) as the acceptor unit and quarterthiophene (Th4) as the donor unit, for which an abnormally strong burn-in degradation is found caused by spinodal demixing of the donor and acceptor phase leading to dramatically reduced charge generation.51 

10.2.3 Polymorphism Polymer conductors and semiconductors can display a rich phase behavior where many different phases and crystal forms can be present— sometimes simultaneously. In many cases, several crystal motifs, i.e., polymorphs, can be induced. As for any material, polymer or molecular, organic or inorganic, this depends on the pressure and temperature that the materials is held at, and often the way the material was processed (e.g., solution concentration, casting temperature).52  The ability of many polymer semiconductors to exist in, or form, more than one crystal structure is impressive. In a single-component system, at a given temperature, T , and pressure, p , principally, only one phase can be present in complete thermodynamic equilibrium.18 , 52 , 53  However, in many organic systems the thermodynamically most stable state is very often not realized. As a consequence, crystal structures and other solid phases can occur that are not necessarily equilibrium structures. This explains the origin of the complex and rich phase behavior of polymer semiconductors, and the consequent range of microstructures and phase morphologies that can be obtained with them. An illustrative example for this diversity in structures that can be accessed with one single material can be given, again, with the ubiquitous poly(3-hexylthiophene). Using precisely defined, monodisperse, regioregular oligo(3-hexylthiophene)s (3HT)n   with n   =  4− 36, Koch et al. have shown,54  for instance, that these model compounds can feature two distinctly different solid-state structures: polymorph of Form I in which the hexyl side-chains are not interdigitated, and Form II in which they are.54  This change in structure drastically changes their phase behavior. The thermodynamic equilibrium melting temperatures of these two phases, for example, are drastically different: T m °  (Form I) ≈  298° C, while T m °  (Form II) ≈  116° C; A notable difference is also observed for the enthalpy of fusion: ∆ H f  (Form II) is about three times that of Form I; in addition, the rate of crystallization of Form II is about one order of magnitude slower than that of Form I. Accordingly, a crossover of the thermodynamically preferred Form II into the kinetically favored Form I is observed for n  =   12.54  In the regime of 10  ≤   n  ≤   21, the materials can be reversibly converted from one polymorph to another by suitable treatments. This example shows, therefore, that the molecular order and structural development that can be induced in polymer semiconductors is frequently the result of combined influences of thermodynamic and kinetic factors. This becomes even more relevant when dealing with polymer semiconductors of complex chemical structures where this interplay between thermodynamics and kinetics can be very pronounced.

10.2.4 Liquid Crystallinity Many polymer semiconductors can form liquid-crystalline phases (i.e., mesophases; see also Section Such phases are good examples of thermodynamically stable, liquid-like states with a high degree of molecular order. The presence of a LC phase in a material is interesting as several studies have suggested that they can lead to molecular architectures of improved π -stacking— or at least of increased molecular order.7 , 55 –  58  For some LC phases, e.g., those of uniaxial (nematic) LC order, the resulting properties of such mesoscopically ordered architectures can be theoretically understood using Maier− Saupe (MS) models.59 –  65   MS models have, for instance, been successfully used to describe nematic mesophases of rod-like, semiconducting polymers such as poly(alkoxy phenylene vinylene)59 as well as semi-flexible poly(alkylthiophenes).63 


Conjugated Polymers

Because of the existence of specific mesophases, liquid crystallinity can in many cases provide versatile pathways towards controlled solid-state microstructures and extend the library of architectures that can be obtained with polymer semiconductors.6 , 7 , 55 –  58  LC mesophases may also broaden the tools by which to control the structure formation of polymer semiconductors56 –  58  and, thus, the optoelectronic properties of such systems. This versatility can, however, render analysis of the phase behavior often more complex, especially as the LC phases often are ‘ hidden’  behind a broad melting process and/ or negatively affected by the proximity of the melting process to the thermal degradation temperature. Hence, significant interest exists in designing semiconducting polymers with easily accessible LC mesophases.

10.3 Multi-Component Systems The library of possible microstructures that polymer semiconductors can adopt can increase when more than one component is present in a given system― which is a common scenario in the organic electronics area. Indeed, in the polymer semiconductor field, additives and extra components are often blended or mixed with the active component to assist with processing, to introduce or enhance certain functions, or to manipulate the material’ s optoelectronic properties. Frequently used ‘ additives’  cover a wide range of systems from solvents, processing aids, insulating polymer ‘ binders’ , other active components (e.g., the acceptor in donor:acceptor solar cell blends), and beyond. These can be of small-molecular or macromolecular nature. Depending on the additive, the specific interactions between the blend components on the molecular level and, among other things, the molecular weight and molecular conformation of the various components, can affect the assembly of such a multicomponent system. Use of an additive can lead to very different phase behavior among the blend materials: from full miscibility in the liquid and the solid state (including formation of co-crystals) to full miscibility in the liquid state but partial or no miscibility in the solid state (forming, e.g., a so-called eutectic system52 , 53 ), to systems of complete immiscibility both in the liquid and the solid state. The phase behavior can have a drastic effect on overall properties and device performance;66 , 67  hence, knowledge of how to establish relevant phase diagrams that assist in understanding the phase behavior of such systems can be extremely helpful to design― from the outset― new materials systems, increase reproducibility, and target specific microstructures and phase morphologies.

10.3.1 Polymer Semiconductor:Solvent Systems Melting Depression Solvents can interact with polymers, semiconducting, conducting, or insulating, through forces such as van der Waals interactions and hydrogen bonding. This can result in a melting point depression of the solute. In semiconductor solvent systems, this usually is exploited when making a solution, e.g., for printing (see. Section The interactions between the solvent (1) and the solute (2) can thereby be quantitatively treated according to standard melting point depression concepts advanced for classical polymer systems by Flory and Huggins:68 , 69 

1 / Tm,2  1 / Tm,2 °  R / DH u ,2 Vu ,2 / Vu ,1   v1  c  v22   

where T m,2  is the melting temperature of the polymer in presence of the solvent; T m °  is the thermodynamic equilibrium melting temperature of infinitely large crystals of infinitely long chains as discussed in Section (although pragmatically often taken to be the melting temperature of the neat polymer); V u,2  is the molar volume of the polymer repeat unit; V u, 1  is the molar volume of the solvent; v 2  is the volume fraction of the polymer; v 1  is the volume fraction of the solvent (where obviously, v 1   =  1 –  v 2 ); and χ  is the Flory– Huggins interaction parameter. Of course, the equivalent equation holds for the

Establishing the Thermal Phase Behavior


depression of the melting point T m,1  of the solvent (or second component) upon addition of a semiconducting polymer. This is of relevance when discussing other systems such as donor:acceptor blends that display a eutectic behavior (see Section 10.3.1). Printing (Liquid/Solid Phase Separation) Organic semiconductors are often processed from solution, i.e., from a two-component system of the semiconductor (or another active material) and a solvent.70  One of the simplest thermodynamic phase behaviors for such a binary system is where there is complete miscibility above the dissolution temperature, T d , but partial or complete immiscibility below T d .70 , 71  Here we shall confine ourselves to binaries of a crystalline or semicrystalline material and a solvent; we will not treat systems comprising amorphous active materials. One important aspect determining the final microstructure formed by the solute (the solvent usually evaporates) is the concentration, i.e., the fraction of solute in the solvent (‘ dilution’ ). For high-molecular weight polymers, the latter parameter governs whether isolated crystals are obtained (e.g., when processing from highly dilute systems; see Section or whether molecularly less ordered, semicrystalline architectures are produced with crystalline lamellae separated by amorphous regions (e.g., via deposition from concentrated solutions).13  Crystallization phenomena need, thereby, to be carefully considered as they can drastically influence— often negatively— important processes during printing. Low room-temperature solubility frequently leads, for example, to inks of undesirably large aggregates. When dispensed into the printing cartridge, these clog the printer; hence, aggregation/crystallization of the semiconductor into larger-scale precipitates needs to be prevented. Temperature/composition cooling diagrams can assist to identify suitable protocols and printing conditions, especially with respect to selection of processing temperatures. Such (non-equilibrium) temperature/composition diagrams can readily be established based on thermal analysis. The latter can be exemplified by the binary system of PBTTT (here substituted with a C12  linear alkyl side-chain) and trichlorobenzene,70  where differential scanning calorimetry, DSC, and cooling thermograms (filled squares; Figure 10.6.) combined with visual observations at low polymer content (open squares; Figure 10.6.) were used to plot a phase diagram to identify relevant ink-jet printing parameters. The diagram reveals, for instance, that PBTTT crystallizes at concentrations in excess of 20  w t% at temperatures > 60° C. However, for more dilute systems (10  w t% polymer content and less) the onset of crystallization of PBTTT (indicated by arrows in Figure 10.6.) is found to be at temperatures 97% to 58% [35]. This indicates that steric hindrance from side chain arrangement plays a key role in determining backbone stiffness. When the alkyl side chains are forced to interact in tail-to-tail or head-to-head coupling, planarization is highly unfavorable. The role of molecular weight on the stiffness of P3HT chains can be elucidated by examining results across multiple publications. Table 11.1 provides radius of gyration values for P3HT samples dissolved in chloroform as a function of molecular weight. The radius of gyration increases as the molecular weight increases. This is expected to give rise to an increased number of chain entanglements. Besides the examination of individual polymer chains, neutron scattering has been used to examine the self-assembly of polymer aggregates in solution. This approach has been applied several times to determine the geometrical information of P3AT nanofibers including width and thickness to make comparisons with dimensions observed in thin films [31, 37–43]. Additionally, estimates of the percent

FIGURE 11.4  Dependence of persistence length extracted from SANS measurements as a function of alkyl side chain length and synthetic route (From McCulloch, B. et al. 2013. Macromolecules, 46:1899. With permission. Copyright 2018 American Chemical Society).


Conjugated Polymers TABLE 11.1  Calculated Radius of Gyration as a Function of P3HT Molecular Weight from the Literature with Similar PDI, Solvent and Concentration. Reference




Molecular Weight (kDa) PDI Solvent Concentration (mg/ml) Fitted Model

12,100 1.10 Chloroform 3 Guinier–Porod Model

35,028 1.26 Chloroform 5 Excluded Volume Model

Radius of Gyration (nm)


17,667 1.17 Chloroform 3–5 PDI-Corrected Worm-like Chain Model 6.3b

a b




Calculated from Mn and PDI 1 Calculated according to Rg2 = Lc l p 3

crystallinity can also be made by applying a mass balance approach to account for scattering from dissolved polymer chains and aggregated domains [40]. Typically, this characterization occurs in the range of 0.009 < q (A-1) < 0.07. The scattering intensity in this region can be modeled using:


I q  fv f FF D rFF

 P q  f 1  f D r 2P q (11.1) 2






Where I is the scattering intensity, q is the scattering vector, fv is the volume fraction of P3HT in solution, Dr is the scattering length density contrast, f FF is the fraction of the sample that is assembled in a specified form factor and PFF and PAmor are the form factors of the aggregate and amorphous domains respectively. The subscript FF describes the selected form factor that can take a variety of forms from cylindrical to parallelepiped depending on the system of study [38, 41, 43, 44]. The fibril domains in P3AT nanofibers have shown good agreement with a parallelepiped model of the form: PPP (q, A, B, C ) =

2 π

2π π

∫ ∫ 0



 sin(qA sin α cos β)   sin(qB sin α cos β)   sin(qC cos α)        sin (α ) dαdβ (11.2)  qA sin α cos β   qB sin α cos β   qA cos α  

here A, B and C are the fiber height, width and length, respectively, and q is the scattering vector. The amorphous form factor can be measured separately by fully dissolving the polymer before crystallization. Modulating the degree of poor and good solvents in solution is a commonly used technique to induce growth of fibrillar domains. This technique relies on changes in supersaturation of the solution to drive polymer chains to self-assemble. In a good solution, there are strong polymer–solvent interactions compared to polymer–polymer interactions causing individual chains to remain fully dissolved. As poor solvent is added, the polymer–polymer interactions become more favorable compared to polymer– solvent interactions resulting in a driving force for π-stacking and fiber formation. Figure 11.5 shows example SANS profiles that indicate that this form factor model is only valid once nanofibrillar domains have formed. This occurs at a critical fraction of poor solvent where the solvent quality induces more favorable polymer-polymer interactions relative to polymer–solvent interactions. The self-assembly of nanofibers of P3HT, P3OT and P3DDT have been investigated using the above parallelepiped form factor model. P3HT was shown to form nanofibers with widths of 22 nm and 23 nm with chloroform-hexane and 1,2-dichlorobenzene-n-dodecane, respectively. Moreover, the heights of nanofibers formed in both solvents were measured at ca. 5 nm [37, 42]. These fiber widths and heights in solution are seemingly invariant to solvent selection, molecular weight and amount of poor solvent added. Furthermore, while fiber dimensions remain constant with the addition of poor solvent, the volume fraction of fibrillar domains and fractal dimension increase. This indicates the growth of an interconnected network of nanostructures. Remarkably, the near-equilibrium fiber dimensions are



FIGURE 11.5  a) Guinier–Porod model fits at low hexane volume fractions before P3HT nanofiber formation in chloroform. Rectangular parallelepiped model fits with solubility induced nanofiber formation (From Keum, J. et al. 2013. CrystEngComm, 15:1114. With permission of the Royal Society of Chemistry), and b) Rectangular parallelepiped model fits of P3HT nanofibers formed in 1,2-dichlorobenzene with dodecane poor solvent (From Newbloom, G. et al. 2014. Soft Matter, 10:8945. With permission of the Royal Society of Chemistry).

similar to the alkyl side length changes. The widths and heights of P3OT and P3DDT nanofibers have been reported as 26.1 and 4.6 nm and 22.2 and 4.9 nm, respectively [42]. P3HT and P3OT show near identical aggregation behavior; once equilibrium is obtained, further processing enhances the network interconnectivity. In contrast, the P3DDT sample demonstrates a more complex self-assembly process with both local and global crystallization events occurring simultaneously. This could arise from the enhanced solubility of P3DDT due to the long side chains enhancing polymer–solvent interactions over polymer-polymer interactions.

11.2.2 UV–Vis Absorbance UV–Vis absorbance spectroscopy has been used extensively in the field of organic electronics to probe the electronic structure of both individual chains and aggregates [45–47]. In the solution phase, nonaggregated, amorphous conjugated polymers display a single broad absorption peak in the optical spectrum. In P3ATs, the peak position is determined by the average number of overlapping π-bonds along the chain backbone and its relation to the local solvent environment. The broad peak arises from the numerous conformational arrangements of each monomer in the chain and their ability to delocalize electrons [35]. In P3ATs, this peak generally resides in the 400–500 nm range. The use of UV–Vis absorbance spectra to quantify the planarity of individual chains in solution has been of interest to the field of conjugated polymers. UV–Vis absorbance can be used to calculate the conjugation length, the length of planar alternating double-single carbon bonds, using the following empirical correlation:

1 1  2.537 *  1.041 (11.3) n Eg

where n is the monomer conjugation number and Eg is the optical bandgap (eV) [35]. Dilute P3ATs in solution exhibit similar absorption line-shapes in a given solvent. This allows comparison of optical bandgaps through the energy (or wavelength) of the absorption peak maximum. Figure 11.6 shows the absorption spectra of four P3ATs in DCB and the position of the absorption peak maximum as the solvent is varied. Initially, the persistence length of a chain was presumed to be an upper limit for the conjugation length as planarity is a key feature linking the two length scales together. However, a larger


Conjugated Polymers

FIGURE 11.6  a) Variation of UV–Vis absorbance peak location as a function of alkyl side chain length, and b) Influence of solvent of UV–Vis absorbance peak location. (From Newbloom, G. et al. 2015. Langmuir, 31:458. With permission. Copyright 2018 American Chemical Society).

number of monomer conformations can contribute to extending the conjugation length compared to the persistence length. Specifically, it has been shown that both the cis and trans conformations allow for extended electron delocalization. Further, molecular dynamics calculations by Bredas et al. indicated that electronic conjugation can be maintained for torsion angles < 40° [47]. Newbloom et al. estimated that the conjugation length may be up to 25% higher than the measured persistence length from SANs [36]. The growth of aggregated domains in solution has also been monitored using UV–Vis absorption spectroscopy. As P3AT chains begin to self-assemble, the absorption spectra develop vibronic shoulder bands at lower energies. The 0–0 and 0–1 vibronic peaks are associated with intrachain and interchain coupling of delocalized excitons. The relative degree of aggregation can be estimated by subtracting the spectra of isolated chains from the spectra obtained from an aggregated sample. Scharsich et al. first demonstrated the utility of this approach by investigating the role of molecular weight on poor solvent induced nucleation of nanofibril domains [48]. Figure 11.7 depicts the results obtained using chloroform as a good solvent and ethyl acetate as a poor solvent. Notably, it was observed that higher molecular weight samples aggregate more readily, as shown by the higher obtainable fraction of aggregates and the ability to aggregate at a lower fraction of poor solvent. Moreover, the fraction of aggregates in solution does not exceed about 60% regardless of the molecular weight or relative degree of processing. This suggests that the majority of chain segments remain dissolved and that a thermodynamic limit exists on the size of aggregates that can form [49, 50]. While this approach has not readily been applied to study general trends in P3ATs, this approach is used extensively to develop process–structure–property relationships in P3HT. Examples of its application will be presented in Section 11.3. A notable advantage of UV–Vis absorption analysis is the ability to measure spectra of solution phase and solidified thin films. As the transition of P3ATs from solution to a thin film occurs, a red shift of the vibronic shoulder peaks occurs, indicting enhanced ordering and planarization of polymer chain backbones. Furthermore, the relative intensities of the 0–0 and 0–1 bands provide information on the exciton bandwidth and thus conjugation length of the polymers in the solid film. Work by Spano et al. developed a mathematical approach based on weakly interacting H-aggregates to quantify the degree of intrachain order [51–54]. The model is expressed as

 W  1  0.24 Ep I00   I 0 1  1  0.073 W  Ep 

2    (11.4)   



FIGURE 11.7  Variation of fraction of aggregates in solution with increasing fraction of ethyl acetate poor solvent for multiple molecular weights of P3HT dissolved in chloroform (From Scharsich, C. et al. 2012. Journal of Polymer Science, 50:442. With permission).

where W is the exciton bandwidth and Ep is the vibrational energy of the symmetric vinyl stretch (taken to be 0.18 eV). Generally, a smaller W is indicative of enhanced intrachain ordering and longer conjugation length [55, 56]. Figure 11.8 provides example thin film spectra of numerous P3AT samples, each displaying unique 0–0 and 0–1 intensities demonstrating the sensitivity of aggregate assembly on the processing conditions. Notably, a consistent trend in the exciton bandwidth as a function of alkyl side

FIGURE 11.8  a) UV–Vis absorption spectra of thin films formed by drop casting various P3ATs from chloroform (From Salammal, S. et al. 2015. European Polymer Journal, 67:199. With permission from Elsevier) and b) UV–Vis absorption spectra of thin films formed by spin-coating P3ATs from chlorobenzene (From Oosterbaan, W. et al. 2010. Advanced Functional Materials, 20:792. With permission).


Conjugated Polymers

chain length cannot be determined. The application of examining the exciton bandwidth from a process–structure–property perspective will also be presented in Section 11.3.

11.2.3 Differential Scanning Calorimetry Thermal characterization has been of interest to researchers to both understand the fundamental crystallization mechanism of conjugated polymers and develop energy efficient post deposition thermal treatments to maximize device performance. The most commonly used technique to monitor phase transitions as a function of temperature is differential scanning calorimetry (DSC). The first informative DSC experiments were conducted by Park and Levon and Liu and Chung on P3DT [57, 58]. Figure 11.9 displays typical heating and cooling curves for P3DT. In both studies, the heating curve displayed a broad peak from ca. 30 to 80 °C and another one from ca. 105 to 155 °C. The low temperature transition arises from side chain melting due to the high regioregularity in the polymers studied. The melting behavior of P3DT in the high temperature region has been shown to be highly dependent on the experimental heat rate and undercooling and can exhibit either two or three phases depending on the sample regioregularity. The first peak describes a quasi-ordered phase where thiophene rings exhibit twisting between neighboring units while the second phase describes an ordered structure in which the thiophene rings are aligned and parallel (polymorphism). The third phase, found only in the work by Liu and Chung, describes interdigitation of the long side chains orientated mainly in the trans-conformation [58]. Two phase transitions at the higher temperature region were also observed by Causin et al. for P3BT and P3DDT, while P3OT and P3HT have been shown to only display single peaks in the DSC thermograms [59]. For all samples, the same transitions are also present in the corresponding cooling curves. The design and optimization of post deposition thermal processing techniques rely on quantified knowledge of glass (Tg), melting (Tm) and crystallization (Tc) temperatures. These temperatures are also dependent on the undercooling and heating or cooling rate applied during the experiment [59, 60]. The dependence of these temperatures on the experimental conditions indicates that phase transitions in these materials are not structural transitions but rather reorganization of non-crystallized domains [57, 61]. The Tg values for P3ATs, when observed, are poorly understood and drastically vary depending on the experiment. The reported Tg for P3HT for example, varies from –113 °C to 106 °C [62, 63]. Instead, Tm and Tc considered here are applicable to improving device performance. Reported Tm and

FIGURE 11.9  Typical thermograph of P3DT at heating and cooling rates of 2.5 °C/min (From Park, K. et al. 1997. Macromolecules, 30:3175. With permission. Copyright 2018 American Chemical Society).


Poly(3-alkylthiophenes) TABLE 11.2  Melting and crystallization temperatures reported across the literature in °C References

[59] [65]















– 224–220

– 195–170

190–188 165–158

170–140 112–88



Tc temperature values are presented in Table 11.2. A wide range of studies have applied knowledge of these temperatures to induce chain reorganization after solution deposition via thermal annealing for improved device performance [16, 62–64].

11.2.4 X-Ray Scattering While neutron scattering has provided key structural information on chain and fiber dynamics in solution, X-ray scattering has been employed to probe the thin film conformations of P3AT chains in the solid state. To this end, several key structural considerations that can be resolved with X-ray scattering are of interest to researchers:

1. How do the lamellar and π-stacking distances evolve with alkyl side chain length? 2. How do the alkyl side chains organize in terms of interdigitation? 3. What is the orientation of polymer chains relative to the substrate interface? 4. How large are the crystalline domains that form?

One of the most important structure metrics that can be extracted from X-ray analysis is spacing between polymer chains in both the lamellar and π-stacking distance. Rieke et al. showed that the lamellar stacking distances of P3AT chains exhibit a monotonic increase as the side chain length increases from 12.63 Å in P3BT to 20.10 Å in P3OT to 30.48 Å in poly(3-tetradecylthiophene) (P3TDT) [17]. This monotonic increase is anticipated as disorder induced from a lack of side chain crystallization would result in the formation of regularly spaced layers. These values are in good agreement with other results obtained in the literature as shown in Figure 11.10 and quantified in Table 11.3. Notably, these lamellar packing distances can vary as a function of molecular weight. Generally, the π-spacing also increases as the side chain length increases. The existence of several P3AT polymorphs in thin films has been shown through identification of differing lamellar and π-stacking distances. Initially, two different crystal structures of P3AT crystals were proposed, referred to as Type I and II. Generally, Type I crystals are characterized by a smaller π-stacking and larger lamellar stacking distance relative to Type II. Figure 11.11 depicts the differences in these two polymorphs. This is suggestive of enhanced interdigitation of polymer side chains in Type II. In terms of processing, Type I crystals are usually formed by casting from chlorinated solvents. Type II crystals are encountered less in thin films relevant for device level applications and can arise from specific solvent preparation conditions or low molecular weight samples [20]. Notably, Type II structures can be irreversibly converted to Type I via thermal annealing. The observed lamellar and π-stacking distances for the multiple polymorphs of select P3AT samples are shown in Table 11.4. Additional studies focusing on P3HT polymorphs to further understand the structure and formation of polymorphs were conducted by selectively synthesizing oligomers and through molecular dynamics (MD) simulations [68–70]. Thermal annealing has also been deemed an important process to manipulate thin film crystallinity. Annealing of most semicrystalline polymers can result in lamellar thickening (increase in lamellar spacing distance) or lateral growth (addition of chains to the fiber via π-stacks) [16]. Abad et al. studied the change in lamellar stacking distance of P3HT and P3OT using grazing incidence X-ray diffraction (GIXRD) as the annealing temperature was varied [65]. Figure 11.12 shows the observed increase in the


Conjugated Polymers

FIGURE 11.10  X-ray diffraction patterns of P3AT nanofibers coated from anisole solutions depicting both lamellar (d100) packing distances and π-stacking distances (d020) (From Samitsu, S. et al. 2008. Macromolecules, 41:8000. With permission. Copyright 2018 American Chemical Society).

lamellar spacing of P3HT from 15.0 Å to 17.1 Å as the annealing temperature was varied from 20 °C to 125 °C. This was associated with an increase in the crystal size promoted by lateral growth. In contrast, a small decrease in the lamellar spacing was observed for P3OT with the spacing decreasing from 20.6 Å at 25 °C to 20.0 Å at 98 °C. The authors justified this trend with improved interdigitation as the annealing temperature increased. Joshi et al. also observed an increase in lamellar stacking distances in P3HT thin films as the annealing temperature increased [62]. TABLE 11.3  Lamellar stacking distances reported across the literature in Å. References [17] [66] [59] [61] [67]







12.6 12.5 12.7 13.1 12.7

16.4 16.1 16.8 – 16.5

20.1 19.8 20.2 20.1 20.5

23.9 – 23.7 – 24.1

27.2 – – 26.2 –

30.5 – – – –



FIGURE 11.11  Diffraction patterns of Type I and Type II polymorphs of P3OT and P3DDT (From Prosa, T. et al. 1996. Macromolecules, 29:3654. With permission. Copyright 2018 American Chemical Society).

Access to grazing incidence X-ray measurements (GIXRD, grazing incidence wide-angle X-ray scattering (GIWAX)) has enabled quantification of the orientation of crystalline domains relative to the substrate. Charge transport along the polymer backbones and through π-stacks have been shown to be the dominant transport pathways in crystalline P3AT domains [75]. This implies that the ‘edge-on’ orientation is preferential in organic field effect transistors where charge must move parallel to the substrate. In contrast, organic photovoltaic applications rely on the ‘face-on’ orientation for separation of holes and electrons. Figure 11.13 highlights the differences in observed diffraction patterns for edge-on and face-on orientations. Salammal et al. observed that the proportion of face-on orientated domains increased as the alkyl side chain length increased from three to eight carbons when spin-coated onto silicon surfaces from chloroform [16]. In general, the edge-on orientation in P3HT arises from spin-coating from high-boiling points solvents and functionalization of substrates with self-assembled monolayers of octadecyltrichlorosilane (OTS), hexamethyldisilane (HMDS) or 3-aminopropyl-triethoxysilance [20]. The degree of crystal orientation relative to the substrate interface can be quantified using the Herman’s orientation factor [75]. TABLE 11.4  Lamellar (d100) and π-stack (d020) spacings of differing polymorphs of P3ATs in Å P3BT











Type I Type I’ Type II










3.8 3.9 4.2

12.9 12.6 9.6

3.81 – 4.37

16.0 15.5 12.0

3.8 – 4.5

20.8 – 14.5

3.8 – 4.6

23.2 – 17.0

3.8 – 4.5

d100 27.1 – 19.8


Conjugated Polymers

FIGURE 11.12  Variation in lamellar spacing as the annealing temperature is increased for a) P3HT and b) P3OT (From Abad, J. et al. 2012. Solar Energy Materials and Solar Cells, 97:109. With permission from Elsevier).

Finally, the crystal grain size in the thin film is important to quantify as enhanced charge transport is expected to occur within ordered domains. The grain size can be calculated using the Scherrer equation with knowledge of a shape factor (typically 0.8–1) and the full width at half-maximum of a diffraction peak [75]. Abad et al. observed a larger grain size for P3HT (ca. 12 nm) compared to P3OT (ca. 9 nm) when spin-coated from toluene [65]. This contrasts with the results obtained by Samitsu et al. who produced thin films by spin-coating P3ATs fully dissolved in solvent. They observed a linear increase in the grain size as the alkyl side chain length increased, with grain sizes of 9.0, 10.5 and 11.0 nm for P3BT, P3HT and P3DT respectively [39]. Finally, Joshi et al. spin coated a similar concentration of P3HT dissolved in chloroform and reported a significantly higher grain size (ca. 23.5 nm), indicating that a wide range of processing conditions can be identified, which can greatly impact the crystal grain size [62].

11.2.5 Atomic Force Microscopy Atomic force microscopy (AFM) allows for the direct visualization of the nanofibrillar domains on length scales that are relevant to device applications. These length scales can generally vary from a few hundreds of nanometers to about 10 µm allowing for structural characterization ranging from the nano- to mesoscale. While generally used to quantitatively compare fibrillar domain sizes, packing

FIGURE 11.13  X-ray diffraction patterns of edge-on and face-on crystal structures in P3HT (From Sirringhaus et al. 1999. Nature, 401:685. With permission).



density and orientation, recent advances in film preparation and image analysis techniques have enabled the quantification of structural features from AFM images [77–79]. This has enabled the elucidation of key process–structure–property relationships in P3AT films, including the need for highly crystalline, interconnected domains that display long-range orientational order to achieve high charge carrier mobility. AFM analysis is one of the few techniques that can provide information on the planarization of nanofibers in the solid state (also TEM, e.g.). P3AT nanofibers are often depicted as having a central backbone made from planarized chains bound together through π-stacks. Amorphous segments exist on the end of each chain. Figure 11.14 provides a visualization of P3AT nanofibers. Remarkably, the fiber lengths obtained from AFM analysis display excellent agreement with those obtained from neutron scattering of nanofibers as shown in Table 11.5. A general trend of increasing fiber width as the alkyl side chain length increases is observed [71]. This can be associated with enhanced stability in the solution phase that transfers to the solid phase as the thin film deposition process occurs. Moreover, these widths represent about 40% of the length of a fully extended chain regardless of the alkyl side chain length. Despite enhanced stability associated with a longer side chain, the majority of each individual chain is not contained with the fiber backbone, but instead is either part of the amorphous matrix that surrounds these fibers or as interconnecting tie chains. A wealth of knowledge has been amassed for the impact of chain length, manipulated through molecular weight, on the fiber width using P3HT. Figure 11.15 displays observed morphologies of various molecular weight P3HT samples and the quantified relationship between molecular weight and fiber width [80, 81]. Initial studies suggested that low molecular weight P3HT can assemble into well-defined but isotropically orientated crystalline domains. In contrast, high molecular weight P3HT does not readily self-assemble into defined crystals [82]. Instead, it was suggested that π-stacks form at select points along the length of the chain to form an amorphous network [83]. As processing methods developed, this was deemed an oversimplification as fibrillar domains have been observed to form in high

FIGURE 11.14  Illustration of chain packing into semicrystalline fibers as a function of polymer molecular weight (From Brinkmann, M. et al. 2009. Macromolecules, 42:1125. With permission. Copyright 2018 American Chemical Society). TABLE 11.5  Comparison of fiber widths of P3ATs as measured by SANS and AFM

SANS Width (nm) AFM Width (nm)







23 24.1

26 25.5

– 27.7

– 27.5

22 –

[37, 42] [71]


Conjugated Polymers

FIGURE 11.15  a) Impact of molecular weight and thin film deposition process on the formation of P3HT nanofibers via AFM analysis (From Verilhac, J. et al. 2006. Synthetic Metals, 156:815. With Permission from Elsevier) and b) Saturation of P3HT nanofiber widths estimated from AFM images with increasing P3HT molecular weight (From Zhang, R. et al. 2006. Journal of American Chemical Society, 128:3480. With permission. Copyright 2018 American Chemical Society).

molecular weight P3HT. Verilhac et al. demonstrated fiber growth from high molecular weight samples by depositing thin films via a dip-coating method [77]. The ability to form nanofibers in high molecular weight samples was attributed to the slow rate of solvent evaporation which provided ample time for self-assembly processes. As expected, lower molecular weight samples can self-assemble into fibrillar structures regardless of solvent evaporation rate. Quantifying the change in fiber width as a function of the molecular weight has helped elucidate fiber formation mechanisms [81, 83]. At low molecular weights, fiber width increases as the molecular weight increases. This suggests that smaller chains can fully incorporate themselves into a fiber with few amorphous lamellar ends. Once the molecular weight reaches approximately 20 kDa, the fiber width plateaus and remains constant with increases in molecular weight. A constant fiber width with increasing molecular weight suggests a thermodynamically driven limit for the size of P3AT crystalline domains. This is consistent with the chain folded model that describes the folding of isolated polymer chains to achieve a lower energy configuration [30, 84]. The amorphous lamellar tails of individual chains have been theorized to contribute to three unique structural features: 1) an amorphous matrix around fibers, 2) folds of individual chains that reincorporate into other sections of the fiber and 3) tie chains that interconnect crystalline domains. Section 11.3 will discuss the role of these amorphous chain tails as they relate to improved charge transport.



FIGURE 11.16  Workflow of AFM fiber vectorization procedure (From Persson, N. et al. 2017. Chemistry of Materials, 29(1):3. With permission. Copyright 2018 American Chemical Society).

Recent advances in computational image analysis enabled quantification of a wide range of structural parameters from AFM images. Persson et al. have developed an open source fiber analysis software package that vectorizes all fibers in an AFM image to enable calculations of nanofiber persistence length, fiber packing density, orientational order parameters and size distributions [78, 79]. The outline of the process flow to perform this vectorization is shown in Figure 11.16. The usefulness of this software package was demonstrated for a variety of P3HT processing methods that will be presented in Section 11.3. It was observed that smaller nanofibers tend to aggregate at the P3HT–substrate interface and pack more densely compared to the P3HT–air interface. As the polymer–air interface is generally more accessible for AFM studies, these results highlight one difficulty in developing robust structure–property methods for P3AT materials. The distribution of microstructures through the entirety of the film can vastly differ.

11.2.6 Charge Carrier Mobility The goal of detailed structural characterization is the development of robust structure-property relationships to enable organic electronic devices. To achieve this goal, the organic field effect transistor serves as an accessible and simple platform to extract electronic properties in the form of the charge carrier mobility. While mobility results are not an intrinsic property and are dependent upon transistor geometry and dimensions, gate and drain testing voltages and testing environment, they provide a useful basis to compare process–structure–property results [85]. Since the early development of P3AT organic field effect transistors (OFETs) in the late 1980s, the impact of a wide range of processing conditions on charge carrier mobility has been investigated. As the role of side chain engineering is primarily designed to manipulate polymer solubility, early investigations studied the role of alkyl length with otherwise constant processing conditions. Consensus was established that long alkyl side chains are detrimental to charge transport [86, 87]. Work by Bao et al., Kaneto et al. and Babel et al. observed decreases in the mobility as the side chain length surpassed eight carbons [88–90]. This was associated with the formation of an insulating layer by the alkyl chains that disrupted hopping pathways during charge transport. This trend of poor mobility results for long alkyl


Conjugated Polymers

FIGURE 11.17  a) Charge carrier mobility values for P3ATs dissolved in chloroform deposited via spin-coating (From Babel, A. et al. 2005. Synthetic Metals, 148:169. With permission from Elsevier), and b) drop-casting (From Suave, G. et al. 2010. Journal of Materials Chemistry, 20:3195. With permission of the Royal Society of Chemistry).

side chain lengths was shown to be a tunable parameter in work by Sauve et al. [91]. In their study, they showed that deposition of P3OT and P3DDT on OTS-8 treated surfaces produced films with higher mobility than P3HT deposited on non-treated silicon surfaces. By treating the surface with alkyl silanes to enhance surface hydrophobicity, polymer–substrate interactions were minimized. This improved the self-assembly of polymers with longer side chains at the interface, providing improved pathways for transport. However, Suave et al. concluded that short side chain polymers may be preferable for future studies due to their reduced sensitivity to surface treatment. Figure 11.17 highlights observed trends in mobility as a function of alkyl side chain length. Inconsistencies in trends demonstrate the heavy dependence of electrical properties on processing conditions. Interestingly, poly(3-hexylthiophene) has developed into the model poly(3-aklythiophene) polymer for process–structure–property relationships, despite P3BT exhibiting higher charge carrier mobility values in numerous studies. One hypothesis for enhanced charge transport for P3BT involves the packing behavior at the dielectric–semiconductor interface [92]. As charge transport in OFETs occurs in the first few nanometers from the interface, a smaller lamellar packing distance enables enhanced packing of edge-on chains. This provides additional pathways for charges to avoid traps. A further justification for improved charge carrier mobility in P3BT compared to P3HT is attributed to a reduction in the hole injection barrier allowing more holes to flow through the active layer [93]. Ultimately, disagreement between studies involving P3BT and P3HT indicates that a wide range of morphological features impact mobility that may have been preferentially optimized in a given study. Exceptions to previously established trends for P3HT-based devices have been a common occurrence over the past decade. This issue has received enhanced attention recently with the rise of big data for materials discovery. In 2016, Persson et al. collected process–property information on over 200 P3HT transistors from over 19 publications into an open source platform to explore global trends [94]. This approach could generally validate established trends from the literature, yet also highlighted the complexity of these systems. For example, one of the earliest process–property relationships observed was an increase in charge carrier mobility with polymer molecular weight. Figure 11.18 demonstrates the generality of this trend across all studies. However, mobility values can vary by multiple orders of magnitude for a given molecular weight. To investigate the impact of multiple influential processing variables, Persson et al. extracted five devices from the database that displayed the most similar processing conditions, as shown in Figure 11.17. These five devices were fabricated with the following similar conditions: 1) neat chloroform as the solvent, 2) spin-coated onto the substrate, 3) on a bottom-gate, bottom contact (BGBC) geometry with gold electrodes, 4) from molecular weight > 20 kDa, 5) no annealing, 6)



FIGURE 11.18  a) Hole mobility vs Mn (in kilodaltons) observed across over 200 devices from the literature, b) Reported mobility values of devices with progressively tighter processing constraints and c) Summary of top five devices conforming to similar processing conditions with vastly different charge carrier mobility values (From Persson, N. et al. 2016. Current Opinion in Solid State and Materials Science, 20:338. With permission from Elsevier).

no solution pretreatment besides dissolution and 7) transistor channels > 10 μm. Despite all these similarities in processing conditions, the device mobilities span over two orders of magnitude. This study highlights the complexity in drawing robust process–property relationships due to the massive design space available in the manufacturing of organic field effect transistors.

11.2.7 Concluding Remarks The wide range of accessible structures and corresponding electronic properties of P3AT-based devices demonstrates the vast tunability and versatility of organic electronics. Minor changes to the length of the alkyl side change can have significant impact on the packing behavior of individual chains. In more recent studies, researchers are designing synthetic techniques to create polymer chains with multiple alkylthiophene monomer units to leverage solubility enhancements of longer side chains with improved packing associated with the shorter side chains. While many structural motifs have already been discovered, it is highly feasible that a wide range of other structures with unique properties can be realized. The vast amount of accumulated knowledge on the family of poly(3-alkylthiophene) polymers provides an unparalleled opportunity to elucidate informative process–structure–property relationships relevant to the behavior of conjugated polymers in general. Structural features and properties described


Conjugated Polymers

above display a strong dependence on polymer molecular weight, polydispersity, regioregularity, solvent choice and processing history. Advances in ‘Big Data’ and materials informatics techniques are only just beginning to leverage existing knowledge to identify promising regions for future experimentation. Developing a complete understanding of P3AT polymers can aid in the rapid advancement of novel conjugated polymer systems. While this chapter focuses solely on organic transistors as the final application, knowledge of the fundamental solution and thin film behavior of P3AT chains can inform processing decisions for a wide range of other applications. As an example, the wide-spread commercial availability of P3AT polymers enables process optimization of phase separation in photovoltaic devices as new high performance n-type polymers develop. Use of a consistent p-type polymer family may enable materials informatics approaches to be applied to photovoltaic processing as well. As the organic electronics field continues to mature, the fundamental understandings obtained from single-component P3AT devices will be invaluable to the development of new technologies.

11.3 Advances in Solution Processing Methods As discussed in the previous section, the role of processing of poly(3-alyklthiophenes) has a profound impact on the structure and resulting properties of deposited thin films. Processing strategies to control the morphology can be separated into three main categories: 1) solution-state preprocessing, 2) controlled deposition methods and 3) solid state post processing. Solution state preprocessing involves disrupting the local solvent environment to induce nucleation and growth of semicrystalline π–π stacked nanofibers [95, 96]. Controlled deposition methods involve manipulation of shear-aligning techniques and solvent evaporation rates to obtain ordered microstructures [63, 97]. Solid-state post processing involves thermal annealing or mechanical manipulation techniques [73]. All of these have been intensely studied in the literature for P3HT, with transferability projected to other semicrystalline conjugated polymers. Solution state preprocessing methods have received significant attention due to their industrial applicability. In this section, recent advances in manipulating and controlling the microstructure and thus charge mobility via solution preprocessing will be discussed. A wide range of processing methods have been proposed to promote self-assembly of nanofibers. These techniques include solvent solubility tuning [26, 98–101], ultrasonication (Son) [96, 102–104], UV irradiation [105–108], solution aging [109, 110] and microfluidics [95], each with their own unique proposed mechanism of self-assembly and accessible nanofibril structures. Stark contrasts in the fibril morphology can be easily visualized using AFM images analyzed via advanced image processing techniques. Persson et al. compiled 100 AFM images from nine unique solution preprocessing techniques produced in the same laboratory setting [78]. These images were analyzed using open source software depicted above in Figure 11.16 to quantify the fiber length density (ρFL), degree of orientational order (Sfull), and mean fiber length. Fiber length density has been defined as the summed length of all fibers per unit area. The orientation order parameter for a given image was calculated according to:

Sfull  2 cos2 qn  1 (11.5)

where qn is the angle between an individual fiber pixel and the average orientation of the population. This order parameter varies between 0, indicating an isotropic arrangement of fibers, and 1, indicating perfect alignment. The degree of orientational alignment can be visualized by presenting processed AFM images as orientation maps, where the color of a fiber is indicative of its orientation. Examples of orientation maps that highlight structural features captured by Sfull and ρFL are shown in Figure 11.19. Notably, similar processing methods appear to cluster around unique values of the structural descriptors. Son+Poor Solvent thin films display a medium packing density and low alignment with minimal tunability as a function of processing. In comparison, Son+Age thin films exhibit both a medium



FIGURE 11.19  Illustration of fiber alignment (Sfull), fiber length density (ρFL), and fraction of nematic domains (a) that arise from various solution processing techniques and film deposition methods. Orientation maps provide visual representation of influence of extracted perimeter on thin film structure (From Persson, N. et al. 2017. Applied Materials and Interfaces, 9:36090. With permission. Copyright 2018 American Chemical Society).

degree of alignment and packing density. Microfluidic+UV processed fibers generally display a high fiber length density, with the degree of processing controlling the fiber alignment. Analysis conducted in this manner can be used to support hypothesized self-assembly mechanisms. Persson et al. analyzed the alignment and packing density results of the three mentioned techniques and summarized differing bundle formations elucidated in the original publications [78]. In the case of poor co-solvent addition with sonication, the presence of the unfavorable solvent environment and cavitation from sonication will induce nucleation of P3HT nanofibers. However, the presence of the poor cosolvent results in few tie chains as the fringe chains coil up to decrease their interaction with the solvent. Upon thin film deposition, each fiber will be exposed to a unique shear force leading to an isotropic film. In the case of sonication and aging, the cavitation effect will once again provide the nucleation step. The aging time allows fibers to diffuse through the solution and for tie chains to extend and form interconnections. These interconnections keep the network together during deposition leading to enhanced local alignment. An extreme case of tie chain interconnectivity is hypothesized to arise from microfluidic processing via the formation of shish-kebabs. In terms of charge carrier mobility, μmicrofluidic > μSon+Age > μSon+Poor Solvent, indicating the importance of tie chains for enhanced global charge transport.


Conjugated Polymers

Expanding this analysis for a more inclusive array of solution preprocessing techniques shows a general correlation between sfull and charge carrier mobility. This has been attributed to tie chains that lead to a network that enables orientational alignment during deposition and a percolative pathway for enhanced charge transport. Figure 11.20 depicts hypothesized structural motifs that fringe chains at the end of fibers can exhibit. While AFM analysis presents a generalizable structure–property relationship that is applicable across a wide range of solution preprocessing techniques, it provides little insight on the role of processing to obtain these structures. It is more desirable to seek complete process–structure–property relationships to be able to apply knowledge on an industrial scale, or to transfer knowledge to another polymer system. However, developing global process–structure–property relationships even under highly controlled processing conditions is a challenging task. Work by Wang et al. on the microfluidic processing of P3HT and Chu et al. on UV-induced control of fibrillar domains exemplify the use of the structural descriptors from Section 11.2 to develop process–structure–property relationships [95, 110]. The work by Wang et al. was motivated by the development of a solution processing technique that could be linked directly to a roll-to-roll deposition environment in a continuous flow process [95]. Solutions of P3HT dissolved in chloroform were flowed through a microfluidic tube that contained a cooling-enhanced nucleation step, and UV-induced growth step. Residence times for each step were kept constant, resulting in the use of the flowrate to modulate shish-kebab bundle formation. Figure 11.21 indicates that a flow rate of 0.25 m/s yielded optimal results. This was justified via optimal structural results as indicated by high solution aggregate fraction with aligned nanofibrils and minimized π-stacking distance and thin film exciton bandwidth. These descriptors suggest the development of long-range order on the nano- through mesoscale that provides a percolative network for enhanced charge transport. As the flow rate was both increased and decreased to 0.6 m/s and 0.10 m/s respectively, similar suboptimal values were obtained for the π-stacking distance, Herman’s orientation factor and exciton bandwidth. Only differences in the solution aggregate fraction and AFM images were observed. It was suggested that the

FIGURE 11.20  a) Decay of orientation order as a function of image length and associated schematic of interfiber interactions (From Persson, N. et al. 2017. Chemistry of Materials, 29:3. With permission). Copyright 2018 American Chemical Society), b) General relationship between orientation alignment and charge carrier mobility as a function of processing methods (From Persson, N. et al. 2017 Accounts of Chemical Research, 50:932. With permission. Copyright 2018 American Chemical Society), and c) Schematic of charge transfer mechanism associated with aligned, interconnected fibrillar domains (From Persson, N. et al. 2017. Applied Materials and Interfaces, 9:36090. With permission. Copyright 2018 American Chemical Society).



FIGURE 11.21  Key electrical and structural characterization techniques utilized to elucidate the mechanism of self-assembly during microfluidic processing of P3HT (From Wang, G. et al. 2015. Chemistry of Materials, 9:8220. With permission. Copyright 2018 American Chemical Society).

higher flow rate produces too many nuclei leading to insufficient growth of interconnected domains and vice-versa for a lower flow rate. This effectively demonstrates the utility of detailed process and structural descriptors to develop mechanistic insights associated with new processing conditions. Work by Chu et al. strived to monitor the time-dependent self-assembly of nanofibers via a nucleation and growth mechanism induced by low dose UV-irradiation and solution aging [110]. The impact of fiber orientation relative to OFET source and drain electrodes was also investigated. P3HT was dissolved in chloroform, exposed to UV light for 8 minutes and left to age in solution. Structural measurements of both the solution and thin film characteristics were incrementally measured across a 24-hour period. Figure 11.22 portrays the observed mobility values and structural descriptors as a function of solution aging time. An increase in the solution aggregate fraction is observed across the 24 hours, confirming the time-dependent self-assembly mechanism. Furthermore, a decrease in the thin film exciton bandwidth and an increase of orientational order is observed, indicating improved chain planarization and network formation as a function of time. Mechanistically, a general trend of enhanced charge transport when fiber backbones are aligned parallel to the electrodes was attributed to tie chains providing efficient transport pathways between crystalline domains. The excellent agreement between structural descriptors in both the solution and thin film phases suggests that in-situ monitoring could be adapted in industrial applications to optimize electrical properties during device fabrication.


Conjugated Polymers

FIGURE 11.22  Key electrical and structural descriptors utilized to monitor the time dependent self-assembly of P3HT nanofibers after UV-induced nucleation (From Chu, P. et al. 2016. Chemistry of Materials, 28:9099. With permission. Copyright 2018 American Chemical Society).

11.4 Deposition Methods Morphology of conjugated polymers plays a critical role for the intrinsic charge transport characteristics and performance of polymer-based devices [21, 111–114]. It has been shown that the conjugated polymer/P3HT crystalline structure significantly influences efficient charge hopping between transport sites and in turn influences device performance [100, 104, 106, 107, 110]. In the past decade, rapid progress has been achieved in conjugated polymer morphology control through solution-based processing methods [96, 115–121]. How to precisely control the morphology of the active layer for achieving high performance electronic devices is a key challenge during solution deposition. As discussed in the previous section, polymer solution pre-treatments will facilitate conjugated polymer aggregation to create crystalline structures. In this section, film deposition techniques that used conjugated polymer solutions, including spin-coating, drop-casting, inkjet printing, dip-coating, solution shearing and blade coating will be introduced and described.

11.4.1 Spin-Coating Spin-coating is a very common technique that uses a solution to form semiconductor films [122]. The film thickness can be easily adjusted by controlling the spin speed, solvent boiling point, solution concentration and viscosity. Recently, this method has been used to consider the simultaneous, one-step



deposition of a mixture of a semiconductor (P3HT) and a soft insulating polymer. The resultant phase separation and self-organized network structure led to improvements in morphology and device performance [123]. Figure 11.23 illustrates the observed phase separation.

11.4.2 Drop-Casting Drop-casting is a facile quasi-equilibrium process which includes the casting of semiconductor and solvent evaporation to generate a thin film [24]. Control of solvent evaporation is the key factor for the film morphology. Improved device performance has been achieved by controlling the solvent evaporation including the use of mixed solvents, saturated solvent environments, surface treatments and sealed chambers [124, 125].

11.4.3 Inkjet Printing Inkjet printing involves ejection of a jet of ink from a chamber and subsequent deposition of the droplet onto a desired substrate. The droplet undergoes a similar drying process to the drop-casting method. Inkjet printing facilitates a type of computer-controlled printing system that forms a desired pattern by propelling the ink onto a substrate. Final film morphology is related to the ink viscosity and substrate type. High-performing semiconducting films fabricated by this method have been achieved by precisely controlling and delicately balancing the various processing parameters [126, 127]. Figure 11.24 demonstrates the inkjet printing technology.

11.4.4 Dip-Coating Dip-coating includes the vertical withdrawal of a substrate dipped into a semiconductor solution [128]. Figure 11.25 highlights the main features of a dip-coating apparatus. Film morphology and thickness can be adjusted by controlling the withdrawal velocity, solution temperature, solvent evaporation rate and substrate pattern. Xue et al. studied the formation of different P3HT morphologies that could be induced to form by dip-coating [129]. Figure 11.26 presents the AFM height images of these morphologies.

FIGURE 11.23  AFM phase images of P3HT/PMMA films with different P3HT/PMMA ratios obtained on bare silicon substrates. The P3HT/PMMA ratios are a) 1:99, b) 2:98, c) 3:97, d) 5:95, e) 8:92 and f) 100:0. In all images, the scale bar represents 200 nm. (From Qui, L. et al. 2008. Advanced Materials, 20:1141. With permission).


Conjugated Polymers

FIGURE 11.24  Inkjet printing of organic single-crystal thin films. a) Schematic of the process. Antisolvent ink (A) is first inkjet-printed (step 1), and then solution ink (B) is overprinted sequentially to form intermixed droplets confined to a predefined area (step 2). Semiconducting thin films grow at liquid–air interfaces of the droplet (step 3), before the solvent fully evaporates (step 4). b) Micrographs of a 20 × 7 array of inkjet-printed C8-BTBT single-crystal thin-films. c) Crossed Nicols polarized micrographs of the film. d), Expanded micrograph of the film, showing stripes caused by molecular-layer steps. e) Atomic-force microscopy image and the height profile (below) showing the step-and-terrace structure on the film surfaces. (From Minemawari, H. 2011. Nature, 475:364. With permission).

11.4.5 Solution Shearing Solution shearing is a highly versatile coating technique where a movable top shearing blade holds a semiconductor droplet above a temperature-controlled substrate [130]. The film thickness can be controlled by the gap size of the blade to the surface. Other coating parameters like substrate surface energy, surface tension of the fluid, viscosity and surface temperature will also influence the film formation,

FIGURE 11.25  Schematic of experimental set-up for dip coating an OTS patterned silicon wafer from the P3HT/ toluene solution. (From Xue, L. et al. 2010. Nanotechnology, 21:145303. With permission).



FIGURE 11.26  AFM height images with an area of 20 μm × 20 μm and the corresponding small areas of 5 μm × 5 μm of P3HT films prepared from 1 mg ml−1 87k-P3HT/toluene solution aged for ((a), (b)) zero days, ((c)–(e)) one day, ((f)–(h)) ten days by dip-coating process, respectively. The red dashed lines in ((a), (c) and (f)) indicate the location of the borderline between the SiOx and OTS areas. (From Xue, L. et al. 2010. Nanotechnology, 21:145303. With permission).

morphology and thickness. This method has been of increasing interest for the fabrication of semiconductor layers and electrodes; and many studies have demonstrated that the solution shearing method can be used to tune the semiconductor packing and crystalline texture, leading to high performance devices. Chang et al. demonstrated that this method can be used to precisely control the P3HT fiber orientation to produce very high performance devices [97]. Figure 11.27 illustrates the high degree of alignment that can be achieved via a solution shearing deposition method. Becerril et al. investigated how the shear speed and substrate temperature influence the semiconductor film morphology [130]. For example, in a P3HT:PCBM blend system, they found that P3HT aggregation and conjugation length are highly dependent on the shearing speed when the solvent evaporation rate is kept constant [131]. Figure 11.28 presents differences in the surface morphologies of these P3HT:PCBM blend systems deposited via solution shearing. The rapid advancement of thin film deposition methods and morphology control strategies during recent years has brought us closer to the bright future promised by organic electronics applications. Choosing appropriate processing techniques and deposition parameters is as important as the choice of semiconductor solution processing methods.


Conjugated Polymers

FIGURE 11.27  AFM phase images (2 μm × 2 μm) of P3HT thin films deposited by blade-coating UV−P3HT solutions with (a) 0 h, (b) 3 h, (c) 8 h, (d) 15 h and (e) 24 h solution-aging time. Image analysis (5 μm × 5 μm) of P3HT thin films deposited by blade-coating UV−P3HT solutions with (f) 0 h, (g)3 h, (h) 8 h, (i) 15 h, and (h) 24 h solution-aging time. (k) Polar plots and (l) orientational order parameter (S2D) of blade-coated films as a function of solution-aging time. (From Chu, P. et al. 2016. Chemistry of Materials, 28:9099. With permission. Copyright 2018 American Chemical Society).

11.5 Semiconductor Crystalline Structure in Flexible and Stretchable Devices As discussed above, semiconductor crystalline structure plays a crucial role for charge transport in organic transistors. The crystalline structure formation is a complex phenomenon that depends on molecular parameters (i.e., molecular weight, concentration, solvent), thermodynamic parameters (i.e., annealing temperature) and processing parameters (i.e., film deposition method). Organic semiconductors have received much attention as an important component in flexible/stretchable electronics [132–134]. The electrical performance of some polymer semiconductors is now comparable to or exceeds that of amorphous Si-based transistors. Such attractive materials can be deposited through simple solution processes, which is advantageous for scale-up. If the electrical performance of these semiconductors under strain can be maintained, these materials can be used to enable flexible/stretchable electronics. Three classical approaches to maintain the electrical performance of semiconducting films under strain have been engineered during the last decades. The first approach is based on geometric designs. Conventional brittle semiconductors are geometrically patterned into curved shapes or deposition of



FIGURE 11.28  AFM images of P3HT:PCBM thin film fabricated by: (a) spin-coating, (b) blade coating, (c) blade coating at 60 °C, (d) blade and spin-coating, (e) blade and spin-coating from chlorobenzene solution and (f) blade and spin-coating from dichlorobenzene solution. The P3HT:PCBM films in (a)–(d) are prepared from toluene solution. (From Chang, Y. et al. 2009. Organic Electronics, 10:741. With permission from Elsevier).

semiconductor films onto a prestrained substrate leads to the formation of periodic wavy shapes upon release of the strain [133, 135, 136]. In this approach, the semiconducting films can be stretched when stress is applied on the underlying substrate and that stress is consumed in straightening the designed shapes. This strategy demonstrates the possibility of imparting stretchability to brittle semiconducting films while maintaining their electrical properties under a certain strain. However, it is still a challenge to retain electrical performance under large strain or deformation using this approach owing to the limit of the dimensions of the semiconductors and their adhesion to the elastic substrate. The second approach is to directly synthesize new semiconductors through the incorporation of modified side chains or dynamic non-covalent molecular units into the polymer chains [137–140]. The flexible molecular segments in the structure will endow intrinsic stretchability to the newly designed semiconducting polymer that can accommodate strain and mechanical stimuli through molecular deformation. This approach is promising and has potential for E-skin or other wearable applications. Nevertheless, the elaborate synthesis process associated with the intrinsically elastic semiconductor polymers and limited stretchability limits their applications. The third route is to blend semiconducting polymers with dissimilar insulating elastomeric polymers to achieve stretchable semiconducting nanocomposites via continuous semiconducting networks [142, 143]. The resulting network structure


Conjugated Polymers

FIGURE 11.29  (A) Aggregated nanofibrillar structures in the ‘solution state’. (B) The film separates into two metastable layers and finally an IPN structure is formed. (C) SEM and AFM images of the IPN structure. Note: To illustrate the fiber aggregation and IPN structure clearly, the cartoons are not drawn to scale. (From Zhang, G. et al. 2017. Chemistry of Materials, 29:7645. With permission. Copyright 2018 American Chemical Society).

in the film will enhance the polymer chain dynamics and is necessary to achieve good charge carrier transport. More recently, P3HT crystalline structures were well controlled in a blend system through phase separation. The resulting films even presented enhanced stretchability compared with the pure soft insulating materials. The electrical properties of P3HT films are not only maintained but greatly enhanced. More importantly, the charge-transport properties of the resulting films under 100% strain almost retain their original level. The phase separated semiconductor crystalline structure proved advantageous because the solubility and surface energy differences between P3HT and polydimethylsiloxane (PDMS) led to increased interactions between the conjugated polymer chains, leading to a highly networked fibrillar structure that extended into the cross-linked PDMS host. In turn, the networked semiconducting film helps increase the fluidity of PDMS chains in the biphasic region, thereby affording enhanced stretchability. The resultant network structure not only improved the effectiveness of charge transport pathways and elasticity but also provided for environmental stability of the devices. Figure 11.29 depicts an example of this hypothesized network structure. On the basis of this versatile semiconducting film, a new practical lift-off approach was explored to integrate the stretchable components for a large area transistor array through solution processing. The obtained robust transistor arrays exhibiting charge carrier mobilities above 1.0 cm2/ V s with excellent durability, even under 100% strain as shown in Figure 11.30. This achievement will have great impact on stretchable optoelectronic devices for practical applications and represents a promising direction for industrial-scale production of stretchable displays and wearable electronic devices. Only one of the more recent results with regard to semiconductor films for stretchable transistors has been discussed in this review. Here, the semiconducting polymer crystalline network in the stretchable active layer was guaranteed by inducing vertical phase-separation with the elastic counterpart, while the elastic, insulating polymer acted as a gate-dielectric or passivation layer. Through appropriate selection of materials and processing conditions for the semiconducting crystalline structure for the active



FIGURE 11.30  Photograph showing the high visual transparency of the transistor arrays without (A) and with (B) strain. (C) Transfer curves obtained from P3HT/PDMS (0.49 wt% P3HT: left), DPP-DTT/PDMS (0.83 wt% DPPDTT: middle), and DPPDPyBT/PDMS (0.62 wt%, DPPDPyBT: right), in its original condition, under 100 % strain (parallel, perpendicular to the charge transport direction), and release condition. (D) Changes in mobility of the stretchable devices in their original condition, under 100% strain and after release. (E) Changes in the mobility after multiple stretching-releasing cycles at 100% strain parallel to the charge transport direction. (From Zhang, G. et al. 2017. Chemistry of Materials, 29:7645. With permission. Copyright 2018 American Chemical Society).

layer, blending semiconductors with soft materials is a promising and effective method for the fabrication of flexible/stretchable electronics.

11.6 Conclusions This chapter provides a broad summary of the research endeavors to elucidate key process–structure– property relationships in the poly(3-alkylthiophene)s. This class of conjugated polymers has received a remarkable degree of attention, leading to the development of a vast array of processing conditions and structural motifs, affording a wide range of electrical properties. Along with improvements in device performance, significant advances in characterization techniques and methodologies to probe these


Conjugated Polymers

structures have been realized. However, inconsistencies between reported results and a lack of generalizable and robust process–structure–property relationships suggest that there is still much to be learned from a fundamental perspective. As the organic electronics field continues to mature from a laboratory curiosity to commercial manufacturing, poly(3-alkylthiophene)s are expected to play a major role. Their ease of synthesis and widespread availability makes them prime candidates to explore the scalability and feasibility of transferring from spin-coating and dip-coating single devices to solution sheared, high-volume, roll-to-roll manufacturing. At the same time, the poly(3-alkylthiophene)s can greatly impact the development of new technologies, including transparent and stretchable devices. With a push for light-weight, flexible, and cost-effective electronics, it is very likely that P3ATs will continue to be model conjugated polymers to push the boundaries of accessible technologies.


1. Rughooputh, S.D.D.V.; Hotta, S.; Heeger, A.J., et al. 1987. “Chromism of soluble polythienylenes”. Journal of Polymer Science: Part B: Polymer Physics, 25, 1071–1078. 2. Assadi, A.; Svensson, C.; Willander, M., et al. 1988. “Field‐effect mobility of poly(3‐hexylthiophene)”. Applied Physics Letters, 53, 195–197. 3. Holliday, S.; Donaghey, J.; McCulloch, I. 2013. “Advances in charge carrier mobilities of semiconducting polymers used in organic transistors”. Chemistry of Materials, 26, 647–663. 4. Earmme, T.; Hwang, Y.J.; Murari, N.M., et al. 2013. “All-polymer solar cells with 3.3% efficiency based on naphthalene diimide-selenophene copolymer acceptor”. Journal of the American Chemical Society, 135, 14960–14963. 5. Sekine, C.; Tsubata, Y.; Yamada, T., et al. 2014. “Recent progress of high performance polymer oled and opv materials for organic printed electronics”. Science and Technology of Advanced Materials, 15, 034203. 6. Rotzoll, R.; Mohapatra, S.; Olariu, V., et al. 2006. “Radio frequency rectifiers based on organic thin-film transistors”. Applied Physics Letters, 88, 123502. 7. Pu, K.; Shuhendler, A.J.; Jokerst, J.V., et al. 2014. “Semiconducting polymer nanoparticles as photoacoustic molecular imaging probes in living mice”. Nature Nanotechnology, 9, 233–239. 8. Knopfmacher, O.; Hammock, M.L.; Appleton, A.L., et al. 2014. “Highly stable organic polymer field-effect transistor sensor for selective detection in the marine environment”. Nature Communications, 5, 2954. 9. Sirringhaus, H. 2014. “25th anniversary article: Organic field-effect transistors: The path beyond amorphous silicon”. Advanced Materials, 26, 1319–1335. 10. Motaung, D.; Malgas, G.; Arendse, C. 2011. “Insights into the stability and thermal degradation of p3ht:C60 blended films for solar cell applications”. Journal of Materials Science, 46, 4942–4952. 11. Manceau, M.; Rivaton, A.; Gardette, J., et al. 2009. “The mechanism of photo- and thermooxidation of poly(3-hexylthiophene) (p3ht) reconsidered”. Polymer Degradation and Stability, 94, 898–907. 12. Sato, M.; Tanaka, S.; Kaeriyama, K. 1986. “Soluble conducting polythiophenes”. Journal of the Chemical Society, Chemical Communications, 873–874. 13. Yoshino, K.; Love, P.; Onoda, M., et al. 1988. “Dependence of absorption spectra and solubility of poly(30alkylthiophene) on molecular structure of solvent”. Japanese Journal of Applied Physics, 27, L2388-L2392. 14. Das, S.; Chatterjee, D.; Ghosh, R., et al. 2015. “Water soluble polythiophenes: Preparation and applications”. RSC Advances, 5, 20160–20177. 15. Zhu, J.; Han, Y.; Kumar, R., et al. 2015. “Controlling molecular ordering in solution-state conjugated polymers”. Nanoscale, 7, 15134–15141.



16. Salammal, S.; Mikayelyan, E.; Grigorian, S., et al. 2012. “Impact of thermal annealing on the semicrystalline nanomorphology of spin-coated thin films of regioregular poly(3-alkylthiophene)s as observed by high-resolution transmission electron microscopy and grazing incidence x-ray diffraction”. Macromolecules, 45, 5575–5585. 17. Chen, T.; Yu, X.; Rieke, R. 1995. “Regiocontrolled synthesis of poly(3-alkylthophenes) mediated by rieke zinc: Their characterization and solid-state properties”. Journal of American Chemical Society, 117, 233–244. 18. Amou, S.; Hama, O.; Shitato, K., et al. 1998. “Head-to-tail regioregularity of poly(3-hexylthiophene) in oxidative coupling polymerization with fecl3”. Journal of Polymer Science Part A: Polymer Chemistry, 37, 1943–1948. 19. Baggioli, A.; Famulari, A. 2014. “On the inter-ring torsion potential of regioregular p3ht: A first principles reexamination with explicit side chains”. Physical Chemistry Chemical Physics, 16, 3983–3994. 20. Brinkmann, M. 2011. “Structure and morphology control in thin films of regioregular poly(3hexylthiophene)”. Journal of Polymer Science Part B: Polymer Physics, 49, 1218–1233. 21. Chang, M.; Lim, G.; Park, B., et al. 2017. “Control of molecular ordering, alignment, and charge transport in solution-processed conjugated polymer thin films”. Polymers, 9, 212. 22. Agbolaghi, S.; Zenoozi, S. 2017. “A comprehensive review on poly(3-alkylthiophene)-based crystalline structures, protocols and electronic applications”. Organic Electronics, 51, 362–403. 23. Wang, H.; Xu, Y.; Yu, X., et al. 2013. “Structure and morphology control in thin films of conjugated polymers for an improved charge transport”. Polymers, 5, 1272–1324. 24. Kim, D.H.; Han, J.T.; Park, Y.D., et al. 2006. “Single-crystal polythiophene microwires grown by self-assembly”. Advanced Materials, 18, 719–723. 25. Kim, D.; Park, Y.; Jang, Y., et al. 2005. “Solvent vapor-induced nanowire formation in poly(3-hexylthiophene) thin films”. Macromolecular Rapid Communications, 26, 834–839. 26. Oh, J.; Shin, M.; Lee, T., et al. 2012. “Self-seeded growth of poly(3-hexylthiophene) (p3ht) nanofibrils by a cycle of cooling and heating in solutions”. Macromolecules, 45, 7504–7513. 27. Brinkmann, M.; Chandezon, F.; Pansu, R., et al. 2009. “Epitaxial growth of highly oriented fibers of semiconducting polymers with a shish-kebab-like superstructure”. Advanced Functional Materials, 19, 2759–2766. 28. Malik, S.; Jana, T.; Nandi, A. 2001. “Thermoreversible gelation of regioregular poly(3-hexylthiophene) in xylene”. Macromolecules, 34, 275–281. 29. Yamamoto, T.; Komarudin, D.; Arai, M., et al. 1998. “Extensize studies on π-stacking of poly(3alkylthiophene-2m4-diyl)s and poly(4-alkylthiazole-2,5-diyl)s by optical spectroscopy, nmr analysis, light scattering analysis, and x-ray crystallography “ Journal of the American Chemical Society, 120, 2047–2057. 30. Hu, D.; Yu, J.; Wong, K., et al. 2000. “Collapse of stiff conjugated polymers with chemical defects into ordered, cylindrical conformations”. Nature, 405, 1030–1033. 31. Chen, C.; Chan, S.; Li, J., et al. 2010. “Formation and thermally-induced disruption of nanowhiskers in poly(3-hexylthiophene)/xylene gel studied by small-angle x-ray scattering”. Macromolecules, 43, 7305–7311. 32. Wu, C.; Wang, X. 1998. “Gobule-to-coil transition of a single homopolymer chain in solution”. Physical Review Letters, 80, 4092–4094. 33. Grosberg, A.; Kuznetsov, D. 1992. “Quantitative theory of the globule-to-coil transistion. 1. Link density distribution in a globule and its radius of gyration”. Macromolecules, 25, 1970–1978. 34. Aime, J.; Schott, B. 1989. “Structural study of conducting polymers in solution”. Synthetic Metals, 28, 11. 35. McCulloch, B.; Ho, V.; Hoarfrost, M., et al. 2013. “Polymer chain shape of poly(3-alkylthiophenes) in solution using small-angle neutron scattering”. Macromolecules, 46, 1899–1907.


Conjugated Polymers

36. Newbloom, G.M.; Hoffmann, S.M.; West, A.F., et al. 2015. “Solvatochromism and conformational changes in fully dissolved poly(3-alkylthiophene)s”. Langmuir, 31, 458–468. 37. Keum, J.; Xiao, K.; Ivanov, I., et al. 2013. “Solvent quality-induced nucleation and growth of parallelepiped nanorods in dilute poly(3-hexylthiophene) (p3ht) solution and the impact on the crystalline morphology of solution-cast thin film”. CrystEngComm, 15, 1114–1124. 38. Yang, H.; LeFevre, S.; Ryu, C., et al. 2007. “Solubility-driven thin film structures of regioregular poly(3-hexyl thiophene) using volatile solvents”. Applied Physics Letters, 90, 172116–172118. 39. Samitsu, S.; Shimomura, T.; Ito, K. 2008. “Nanofiber preparation by whisker method using solvent-soluble conducting polymers”. Thin Solid Films, 516, 2478–2486. 40. Newbloom, G.; Kim, F.; Jenekhe, S., et al. 2011. “Mesoscale morphology and charge transport in colloidal networks of poly(3-hexylthiophene)”. Macromolecules, 44, 3801–3809. 41. Liu, J.; Arif, M.; Zou, J., et al. 2009. “Controlling poly(3-hexylthiophene) crystal dimension: Nanowhiskers and nanoribbons”. Macromolecules, 42, 9390–9393. 42. Newbloom, G.M.; de la Iglesia, P.; Pozzo, L.D. 2014. “Controlled gelation of poly(3-alkylthiophene) s in bulk and in thin-films using low volatility solvent/poor-solvent mixtures”. Soft Matter, 10, 8945–8954. 43. Mittelbach, P.; Porod, G. 1961. Acta Physica Austriaca, 14, 185. 44. Rahimi, K.; Botiz, I.; Agumba, J., et al. 2014. “Light absorption of poly(3-hexylthiophene) single crystals”. RSC Advances, 4, 11121–11123. 45. Roehling, J.; Arslan, I.; Moulé, A. 2012. “Controlling microstructure in poly(3-hexylthiophene) nanofibers”. Journal of Materials Chemistry, 22, 2498–2506. 46. Niles, E.; Roehling, J.; Yamagata, H., et al. 2012. “J-aggregate behavior in poly-3-hexylthiophene nanofibers”. The Journal of Physical Chemistry Letters, 3, 259–263. 47. Brédas, J.L.; Street, G.B.; Thémans, B., et al. 1985. “Organic polymers based on aromatic rings (polyparaphenylene, polypyrrole, polythiophene): Evolution of the electronic properties as a function of the torsion angle between adjacent rings”. The Journal of Chemical Physics, 83, 1323–1329. 48. Scharsich, C.; Lohwasser, R.; Sommer, M., et al. 2012. “Control of aggregate formation in poly(3hexylthiophene) by solvent, molecular weight, and synthetic method”. Journal of Polymer Science Part B: Polymer Physics, 50, 442–453. 49. Panzer, F.; Sommer, M.; Bässler, H., et al. 2015. “Spectroscopic signature of two distinct h-aggregate species in poly(3-hexylthiophene)”. Macromolecules, 48, 1543–1553. 50. Panzer, F.; Bassler, H.; Lohwasser, R., et al. 2014. “The impact of polydispersity and molecular weight on the order-disorder transition in poly(3-hexylthiophene)”. Journal of Physical Chemistry Letters, 5, 2742–2747. 51. Yamagata, H.; Spano, F.C. 2012. “Interplay between intrachain and interchain interactions in semiconducting polymer assemblies: The hj-aggregate model”. The Journal of Chemical Physics, 136, 184901. 52. Spano, F.C.; Silva, C. 2014. “H- and j-aggregate behavior in polymeric semiconductors”. Annual Review of Physical Chemistry, 65, 477–500. 53. Spano, F.C. 2006. “Absorption in regio-regular poly(3-hexyl)thiophene thin films: Fermi resonances, interband coupling and disorder”. Chemical Physics, 325, 22–35. 54. Spano, F.C. 2005. “Modeling disorder in polymer aggregates: The optical spectroscopy of regioregular poly(3-hexylthiophene) thin films”. The Journal of Chemical Physics, 122, 234701. 55. Clark, J.; Silva, C.; Friend, R.H., et al. 2007. “Role of intermolecular coupling in the photophysics of disordered organic semiconductors: Aggregate emission in regioregular polythiophene”. Physical Review Letters, 98, 206406. 56. Clark, J.; Chang, J.; Spano, F., et al. 2009. “Determining exciton bandwidth and film microstructure in polythiophene films using linear absorption spectroscopy”. Applied Physics Letters, 94, 163306.



57. Park, K.C.; Levon, K. 1997. “Order-disorder transition in the electroactive polymer poly(3-dodecylthiophene)”. Macromolecules, 30, 3175–3183. 58. Liu, S.L.; Chung, T.S. 2000. “Crystallization and melting behavior of regioregular poly(3-dodecylthiophene)”. Polymer, 41, 2781–2793. 59. Causin, V.; Marega, C.; Marigo, A. 2005. “Crystallization and melting behavior of poly(3-butylthiophene), poly(3-octylthiophene), and poly(3-dodecylthiophene)”. Macromolecules, 38, 409–415. 60. Peng, Y.; He, Z.; Li, H., et al. 2016. “Understanding the phase behavior from multiple-step isothermally crystallized poly(3-hexylthiophene)s”. Polymer, 98, 61–69. 61. Samitsu, S.; Shimomura, T.; Heike, S., et al. 2008. “Effective production of poly(3-alkylthiophene) nanofibers by means of whisker method using anisole solvent: Structural, optical, and electrical properties”. Macromolecules, 41, 8000–8010. 62. Joshi, S.; Grigorian, S.; Pietsch, U., et al. 2008. “Thickness dependence of the crystalline structure and hole mobility in thin films of low molecular weight poly(3-hexylthiophene)”. Macromolecules, 41, 6800–6808. 63. Cho, S.; Lee, K.; Yuen, J., et al. 2006. “Thermal annealing-induced enhancement of the field-effect mobility of regioregular poly(3-hexylthiophene) films”. Journal of Applied Physics, 100, 114503. 64. Zen, A.; Pfaum, J.; Hirschmann, S., et al. 2004. “Effect of molecular weight and annealing of poly(3hexylthiophene)s on the performance or organic field-effect transistors”. Advanced Functional Materials, 14, 757–764. 65. Abad, J.; Espinosa, N.; Ferrer, P., et al. 2012. “Molecular structure of poly(3-alkyl-thiophenes) investigated by calorimetry and grazing incidence x-ray scattering”. Solar Energy Materials and Solar Cells, 97, 109–118. 66. Gustafsson, G.; Inganas, O.; Osterholm, H., et al. 1991. “X-ray diffraction and infra-red spectroscopy studies of orientated poly(3-alkylthiophenes)”. Polymer, 32, 1574–1450. 67. Yuan, Y.; Shu, J.; Kolman, K., et al. 2016. “Multiple chain packing and phase composition in regioregular poly(3-butylthiophene) films”. Macromolecules, 49, 9493–9506. 68. Koch, F.P.; Smith, P.; Heeney, M. 2013. ““Fibonacci's route” to regioregular oligo(3-hexylthiophene) s”. Journal of the American Chemical Society, 135, 13695–13698. 69. Koch, F.P.; Heeney, M.; Smith, P. 2013. “Thermal and structural characteristics of oligo(3-hexylthiophene)s (3ht)n, n = 4-36”. Journal of the American Chemical Society, 135, 13699–13709. 70. Zhugayevych, A.; Mazaleva, O.; Naumov, A., et al. 2018. “Lowest-energy crystalline polymorphs of p3ht”. The Journal of Physical Chemistry C, 122, 9141–9151. 71. Yuan, Y.; Zhang, J.; Sun, J., et al. 2011. “Polymorphism and structural transition around 54 °C in regioregular poly(3-hexylthiophene) with high crystallinity as revealed by infrared spectroscopy”. Macromolecules, 44, 9341–9350. 72. Prosa, T.; Winokur, M.; McCullough, R. 1996. “Evidence of a novel side chain structure in regioregular poly(3-alkylthiophenes)”. Macromolecules, 29, 3654–3656. 73. McCullough, R.; Tristram-Nagle, S.; Williams, S., et al. 1993. “Self-orienting head-to-tail poly(3alkylthiophenes): New insights on structure-property relationships in conducting polymers”. Journal of American Chemical Society, 115, 4910–4911. 74. Meille, S.; Romita, V.; Caronna, T., et al. 1997. “Influence of molecular weight and regioregularity on the polymorphic behavior of poly(3-decylthiophenes)”. Macromolecules, 30, 7898–7905. 75. Rivnay, J.; Mannsfeld, S.C.; Miller, C.E., et al. 2012. “Quantitative determination of organic semiconductor microstructure from the molecular to device scale”. Chemical Reviews, 112, 5488–5519. 76. Jordens, S.; Isa, L.; Usov, I., et al. 2013. “Non-equilibrium nature of two-dimensional isotropic and nematic coexistence in amyloid fibrils at liquid interfaces”. Nature Communications, 4, 1917. 77. Verilhac, J.; LeBlevennec, G.; Djurado, D., et al. 2006. “Effect of macromolecular parameters and processing conditions on supramolecular organisation, morphology and electrical transport properties in thin layers of regioregular poly(3-hexylthiophene)”. Synthetic Metals, 156, 815–823.


Conjugated Polymers

78. Persson, N.E.; Rafshoon, J.; Naghshpour, K., et al. 2017. “High-throughput image analysis of fibrillar materials: A case study on polymer nanofiber packing, alignment, and defects in organic field effect transistors”. ACS Applied Materials & Interfaces, 9, 36090–36102. 79. Persson, N.; McBride, M.; Grover, M., et al. 2016. “Automated analysis of orientational order in images of fibrillar materials”. Chemistry of Materials, 29, 3–14. 80. Zhang, R.; Li, B.; Lovu, M., et al. 2006. “Nanostructure dependence of field-effect mobility in regioregular poly(3-hexylthiophene) thin film effect transistors”. Journal of American Chemical Society, 128, 3480–3481. 81. Kline, R.; McGehee, M.; Kadnikova, E., et al. 2003. “Controlling the field-effect mobility of regioregular polythiophene by changing the molecular weight”. Advanced Materials, 15, 1519–1522. 82. Kline, R.; McGehee, M.; Kadnikova, E., et al. 2005. “Dependence of regioregular poly(3-hexylthiophene) film morphology and field-effect mobility on molecular weight”. Macromolecules, 38, 3312–3319. 83. Chang, J.; Clark, J.; Zhao, N., et al. 2006. “Molecular-weight depdence of interchain polaron delocalization and exciton bandwidth in high-mobility conjugated polymers”. Physical Review B, 74, 115318–115329. 84. Han, Y.; Guo, Y.; Chang, Y., et al. 2014. “Chain folding in poly(3-hexylthiophene) crystals”. Macromolecules, 47, 3708–3712. 85. Horowitz, G. 1998. “Organic field effect transistors.” Advanced Materials, 10, 365–377. 86. Savagatrup, S.; Printz, A.; Wu, H., et al. 2015. “Viability of stretchable poly(3-heptylthiophene) (p3hpt) for organic solar cells and field-effect transistors”. Synthetic Metals, 203, 208–214. 87. Oosterbaan, W.; Bolsée, J.; Gadisa, A., et al. 2010. “Alkyl-chain-length-independent hole mobility via morphological control with poly(3-alkylthiophene) nanofibers”. Advanced Functional Materials, 20, 792–802. 88. Kaneto, K.; Lim, W.; Takashima, W., et al. 2000. “Alkyl chain length dependence of field-effect mobilities in regioregular poly(3-alkylthiophene) films”. Japanese Journal of Applied Physics, 39, L372–374. 89. Bao, Z.; Feng, Y.; Dodabalapur, A., et al. 1998. “High-performance plastic transistor fabricated by printing techniques”. Chemistry of Materials, 9, 1299–1302. 90. Babel, A.; Jenekhe, S. 2005. “Alkyl chain length dependence of the field-effect carrier mobility in regioregular poly(3-alkylthiophene)s”. Synthetic Metals, 148, 169–173. 91. Sauvé, G.; Javier, A.; Zhang, R., et al. 2010. “Well-defined, high molecular weight poly(3-alkylthiophene)s in thin-film transistors: Side chain invariance in field-effect mobility”. Journal of Materials Chemistry, 20, 3195. 92. Park, Y.; Kim, D.; Jang, Y., et al. 2006. “Effect of side chain length on molecular ordering and fieldeffect mobility in poly(3-alkylthiophene) transistors”. Organic Electronics, 7, 514–520. 93. Lee, H.; Cho, J.; Cho, K., et al. 2013. “Alkyl side chain length modulates the electronic structure and electrical characteristics of poly(3-alkylthiophene) thin films”. The Journal of Physical Chemistry C, 117, 11764–11769. 94. Persson, N.; McBride, M.; Grover, M., et al. 2016. “Silicon valley meets the ivory tower: Searchable data repositories for experimental nanomaterials research”. Current Opinion in Solid State and Materials Science, 20, 338–343. 95. Wang, G.; Persson, N.; Chu, P.H., et al. 2015. “Microfluidic crystal engineering of π-conjugated polymers”. ACS Nano, 9, 8220–8230. 96. Aiyar, A.R.; Hong, J.I.; Izumi, J., et al. 2013. “Ultrasound-induced ordering in poly(3-hexylthiophene): Role of molecular and process parameters on morphology and charge transport”. ACS Applied Materials & Interfaces, 5, 2368–2377. 97. Chang, M.; Choi, D.; Egap, E. 2016. “Macroscopic alignment of one-dimensional conjugated polymer nanocrystallites for high-mobility organic field-effect transistors”. ACS Applied Materials & Interfaces, 8, 13484–13491.



98. Park, Y.; Lee, H.; Choi, Y., et al. 2009. “Solubility-induced ordered polythiophene precursors for high-performance organic thin-film transistors”. Advanced Functional Materials, 19, 1200–1206. 99. Lee, Y.; Oh, J.Y.; Son, S.Y., et al. 2015. “Effects of regioregularity and molecular weight on the growth of polythiophene nanofibrils and mixes of short and long nanofibrils to enhance the hole transport”. ACS Applied Materials & Interfaces, 7, 27694–27702. 100. Kleinhenz, N.; Rosu, C.; Chatterjee, S., et al. 2015. “Liquid crystalline poly(3-hexylthiophene) solutions revisited: Role of time-dependent self-assembly”. Chemistry of Materials, 27, 2687–2694. 101. Chang, M.; Choi, D.; Fu, B.; Reichmanis, E. 2013. “Solvent based hydrogen bonding: Impact on poly(3-hexylthiophene) nanoscale morphology and charge transport characteristics”. ACS Nano, 7, 5402–5413. 102. Zhao, K.; Xue, L.; Liu, J., et al. 2010. “A new method to improve poly(3-hexyl thiophene) (p3ht) crystalline behavior: Decreasing chains entanglement to promote order-disorder transformation in solution”. Langmuir, 26, 471–477. 103. Zhao, K.; Khan, H.; Li, R., et al. 2013. “Entanglement of conjugated polymer chains influences molecular self-assembly and carrier transport”. Advanced Functional Materials, 23, 6024–6035. 104. Choi, D.; Chang, M.; Reichmanis, E. 2015. “Controlled assembly of poly(3-hexylthiophene): Managing the disorder to order transition on the nano- through meso-scales”. Advanced Functional Materials, 25, 920–927. 105. Xue, X.; Chandler, G.; Zhang, X., et al. 2015. “Oriented liquid crystalline polymer semiconductor films with large ordered domains”. ACS Applied Materials & Interfaces, 7, 26726–26734. 106. Chang, M.; Lee, J.; Kleinhenz, N., et al. 2014. “Photoinduced anisotropic supramolecular assembly and enhanced charge transport of poly(3-hexylthiophene) thin films”. Advanced Functional Materials, 24, 4457–4465. 107. Chang, M.; Lee, J.; Chu, P.H., et al. 2014. “Anisotropic assembly of conjugated polymer nanocrystallites for enhanced charge transport”. ACS Applied Materials & Interfaces, 6, 21541–21549. 108. Chang, M.; Choi, D.; Wang, G., et al. 2015. “Photoinduced anisotropic assembly of conjugated polymers in insulating polymer blends”. ACS Applied Materials & Interfaces, 7, 14095–14103. 109. Kleinhenz, N.; Persson, N.; Xue, Z., et al. 2016. “Ordering of poly(3-hexylthiophene) in solutions and films: Effects of fiber length and grain boundaries on anisotropy and mobility”. Chemistry of Materials, 28, 3905–3913. 110. Chu, P.H.; Kleinhenz, N.; Persson, N., et al. 2016. “Toward precision control of nanofiber orientation in conjugated polymer thin films: Impact on charge transport”. Chemistry of Materials, 28, 9099–9109. 111. Virkar, A.A.; Mannsfeld, S.; Bao, Z., et al. 2010. “Organic semiconductor growth and morphology considerations for organic thin-film transistors”. Advanced Materials, 22, 3857–3875. 112. Liu, S.; Wang, W.M.; Briseno, A.L., et al. 2009. “Controlled deposition of crystalline organic semiconductors for field-effect-transistor applications”. Advanced Materials, 21, 1217–1232. 113. Lim, J.A.; Lee, H.S.; Lee, W.H., et al. 2009. “Control of the morphology and structural development of solution-processed functionalized acenes for high-performance organic transistors”. Advanced Functional Materials, 19, 1515–1525. 114. Diao, Y.; Shaw, L.; Bao, Z., et al. 2014. “Morphology control strategies for solution-processed organic semiconductor thin films”. Energy & Environmental Science, 7, 2145–2159. 115. Wang, S.; Kiersnowski, A.; Pisula, W., et al. 2012. “Microstructure evolution and device performance in solution-processed polymeric field-effect transistors: The key role of the first monolayer”. Journal of the American Chemical Society, 134, 4015–4018. 116. Park, M.S.; Aiyar, A.; Park, J.O., et al. 2011. “Solvent evaporation induced liquid crystalline phase in poly(3-hexylthiophene)”. Journal of the American Chemical Society, 133, 7244–7247. 117. Park, J.; Lee, S.; Lee, H.H. 2006. “High-mobility polymer thin-film transistors fabricated by solvent-assisted drop-casting”. Organic Electronics, 7, 256–260.


Conjugated Polymers

118. Park, B.; Aiyar, A.; Park, M.S., et al. 2011. “Conducting channel formation in poly(3-hexylthiophene) field effect transistors: Bulk to interface”. The Journal of Physical Chemistry C, 115, 11719–11726. 119. Diao, Y.; Tee, B.C.; Giri, G., et al. 2013. “Solution coating of large-area organic semiconductor thin films with aligned single-crystalline domains”. Nature Materials, 12, 665–671. 120. Chang, J.-F.; Sun, B.; Breiby, D.W., et al. 2004. “Enhanced mobility of poly(3-hexylthiophene) transistors by spin-coating from high-boiling-point solvents”. Chemistry of Materials, 16, 4772–4776. 121. Aiyar, A.R.; Hong, J.-I.; Reichmanis, E. 2012. “Regioregularity and intrachain ordering: Impact on the nanostructure and charge transport in two-dimensional assemblies of poly(3-hexylthiophene)”. Chemistry of Materials, 24, 2845–2853. 122. Chua, L.L.; Ho, P.K.H.; Sirringhaus, H., et al. 2004. “Observation of field‐effect transistor behavior at self‐organized interfaces”. Advanced Materials, 16, 1609. 123. Qiu, L.; Lim, J.A.; Wang, X., et al. 2008. “Versatile use of vertical-phase-separation-induced bilayer structures in organic thin-film transistors”. Advanced Materials, 20, 1141–1145. 124. Wang, S.; Kappl, M.; Liebewirth, I., et al. 2012. “Organic field-effect transistors based on highly ordered single polymer fibers”. Advanced Materials, 24, 417–420. 125. Goto, O.; Tomiya, S.; Murakami, Y., et al. 2012. “Organic single-crystal arrays from solution-phase growth using micropattern with nucleation control region”. Advanced Materials, 24, 1117–1122. 126. Minemawari, H.; Yamada, T.; Matsui, H., et al. 2011. “Inkjet printing of single-crystal films”. Nature, 475, 364–367. 127. Arias, A.C.; Ready, S.E.; Lujan, R., et al. 2004. “All jet-printed polymer thin-film transistor activematrix backplanes”. Applied Physics Letters, 85, 3304–3306. 128. Rogowski, R.Z.; Dzwilewski, A.; Kemerink, M., et al. 2011. “Solution processing of semiconducting organic molecules for tailored charge transport properties”. The Journal of Physical Chemistry C, 115, 11758–11762. 129. Xue, L.; Gao, X.; Zhao, K., et al. 2010. “The formation of different structures of poly(3-hexylthiophene) film on a patterned substrate by dip coating from aged solution”. Nanotechnology, 21, 145303. 130. Becerril, H.A.; Roberts, M.E.; Liu, Z., et al. 2008. “High‐performance organic thin‐film transistors through solution‐sheared deposition of small‐molecule organic semiconductors”. Advanced Materials, 20, 2588–2594. 131. Reinspach, J.A.; Diao, Y.; Giri, G., et al. 2016. “Tuning the morphology of solution-sheared p3ht:Pcbm films”. ACS Applied Materials & Interfaces, 8, 1742–1751. 132. Someya, T.; Sekitani, T.; Iba, S., et al. 2004. “A large-area, flexible pressure sensor matrix with organic field-effect transistors for artificial skin applications”. Proceedings of the National Academy of Sciences of the United States of America, 101, 9966–9970. 133. Kim, D.H.; Lu, N.; Ghaffari, R., et al. 2011. “Materials for multifunctional balloon catheters with capabilities in cardiac electrophysiological mapping and ablation therapy”. Nature Materials, 10, 316–323. 134. Kaltenbrunner, M.; Sekitani, T.; Reeder, J., et al. 2013. “An ultra-lightweight design for imperceptible plastic electronics”. Nature, 499, 458–463. 135. Sun, Y.; Choi, W.M.; Jiang, H., et al. 2006. “Controlled buckling of semiconductor nanoribbons for stretchable electronics”. Nature Nanotechnology, 1, 201–207. 136. Shyu, T.C.; Damasceno, P.F.; Dodd, P.M., et al. 2015. “A kirigami approach to engineering elasticity in nanocomposites through patterned defects”. Nature Materials, 14, 785–789. 137. Sekitani, T.; Nakajima, H.; Maeda, H., et al. 2009. “Stretchable active-matrix organic light-emitting diode display using printable elastic conductors”. Nature Materials, 8, 494–499. 138. Park, M.; Im, J.; Shin, M., et al. 2012. “Highly stretchable electric circuits from a composite material of silver nanoparticles and elastomeric fibres”. Nature Nanotechnology, 7, 803–809.



139. Oh, J.Y.; Rondeau-Gagne, S.; Chiu, Y.C., et al. 2016. “Intrinsically stretchable and healable semiconducting polymer for organic transistors”. Nature, 539, 411–415. 140. Liang, J.; Li, L.; Niu, X., et al. 2013. “Elastomeric polymer lissght-emitting devices and displays”. Nature Photonics, 7, 817–824. 141. Kim, K.S.; Zhao, Y.; Jang, H., et al. 2009. “Large-scale pattern growth of graphene films for stretchable transparent electrodes”. Nature, 457, 706–710. 142. Zhang, G.; McBride, M.; Persson, N., et al. 2017. “Versatile interpenetrating polymer network approach to robust stretchable electronic devices”. Chemistry of Materials, 29, 7645–7652. 143. Xu, J.; Wang, S.; Wang, G.-J.N., et al. 2017. “Highly stretchable polymer semiconductor films through the nanoconfinement effect”. Science, 355, 59–64.

12 Microstructural Characterization of Conjugated Organic Semiconductors by X-Ray Scattering 12.1 Introduction.......................................................................................391 12.2 Fundamentals of X-Ray Scattering.................................................394 Wide-Angle X-Ray Scattering (WAXS)  •  Small Angle X-Ray Scattering (SAXS)

12.3 Applications in Conjugated Semiconductors (Selected Examples)............................................................................................410

Maged Abdelsamie and Michael F. Toney

Crystal Structure and Molecular Packing of Small-Molecules for Organic Thin-Film Transistor (OTFT)  •  Estimation of Volume Fraction of Phases in Bulk Heterojunction (BHJ) Photovoltaics  •  Probing the Surface and the Bulk of Small-Molecule Thin Films  •  Microstructural Evolution for P3HT:PCBM During Spin-Coating from One Solvent  •  In Situ GISAXS for Probing Phase-Separation Evolution using Multiphase Modeling Based on TSI  •  Co-Solvent Processing for Reducing Domains Over-Coarsening by Influencing the Liquid-Liquid Phase Separation

12.4 Summary and Outlook..................................................................... 418 Acknowledgment........................................................................................... 419 References....................................................................................................... 419

12.1 Introduction Conjugated organic semiconductors have received tremendous attention in the past decades due to their availability, diversity, flexibility, the potential for large-area processing into flexible or rigid substrates, and tunability through chemical synthesis of a virtually infinite number of structures [1–4]. Thus, they offer the promise of significant benefits for a wide range of optoelectronic applications such as organic thin film transistors (OFETs), biosensors, logic integrated circuits, organic photovoltaics, organic light emitting diodes (OLEDs), organic thermoelectrics, photodetectors, and photoresistors [5, 6]. For optoelectronic applications, organic semiconductors are used mainly in the form of thin films that are deposited using a variety of coating techniques including physical-vapor-deposition and solution-processing methods such as spin-coating, slot-die-coating, blade-coating, spray-coating, and printing [6, 7]. The processing routes and conditions, in addition to the choice of organic material, have significant 391


Conjugated Polymers

implications on the microstructure of thin films, and hence their optoelectronic, thermal, and mechanical properties impacting their functionality and durability in devices. Therefore, establishing processing​–micr​ostru​cture​–prop​erty–​perfo​rmanc​e relationships has been a focus for numerous studies [8, 9]. In-depth characterization of the microstructure is crucial for establishing such relationships, while understanding and precise control of the microstructure are necessary for customization of the thin film properties towards the appropriate functionality and optimal performance in optoelectronic devices. Organic thin films are utilized in optoelectronic applications to form the active layer or the buffer/ contact layers (i.e. interlayer between electrodes and active layer), and less often, the electrodes [5]. The organic molecules are made of building blocks of different lengths and structures including small-molecules, medium-sized-molecules, homopolymers, and copolymers [5]. The anisotropic chemical structure of the molecular building blocks and the weak intermolecular forces between molecules add complexity to the packing structure and promote the presence of a relatively high density of defects as compared to inorganics, making the microstructural characterization challenging. The complexity of the microstructure is further amplified when blending two or more compounds in composite systems, such as bulk heterojunction (BHJ) for photovoltaics [10]. Detailed characterization of the microstructure requires quantitative and qualitative determination of a variety of microstructural parameters at different length scales from the molecular scale to the nanoscale and macroscale morphology. X-ray scattering techniques have been highly successful to uncover microstructural features at many length scales and are used with other thin film characterization techniques for a comprehensive description of the microstructure [8, 9]. At the molecular scale, X-ray diffraction is used for the determination of the crystal packing structure and the molecular orientation with respect to the device architecture. In particular, wide-angle X-ray scattering (WAXS) is widely used to uncover the packing structure at the molecular scale, ranging from Angstrom (Å) to a few nanometers (nm) [11]. Optoelectronics properties in conjugated semiconductors are defined by the intermolecular and intramolecular π–π interactions and their overlap along the π-conjugated system which are highly affected by the intermolecular π–π stacking distance and the slip angle between adjacent molecules [12–15]. Such π–π interactions, and hence optoelectronics properties, are anisotropic due to the asymmetry of the packing structure, highlighting the need for the determination of molecular orientations. For instance, fast charge transport is facilitated along specific stacking directions, making it important to define molecular orientation with respect to the electrodes, where charges are collected [16, 17]. Moreover, organic thin film durability and stability are highly impacted by intermolecular forces between the building blocks which are defined by the packing structure [18–21]. For instance, as shown in Figure 12.1a–f, two single-crystal polymorphs of the same small-molecule (Cl2-NDI), obtained at different processing conditions, exhibit different crystal packing structure and hence distinct functional properties when used in thin film transistors [21]. Although both phases exhibited excellent transistor performance, the α-phase achieved higher electron mobility (~ 2.5 times) as compared to β-phase, but β-phase exhibited better thermal stability and maintained superior electron mobility at high temperature. Such distinct properties are attributed solely to the crystal packing structure as both polymorphs were grown as single crystals. Unlike the previous example of single crystals, microstructural characterization of organic thin films becomes more complex for polycrystalline films where there is contribution from crystal defects and grain boundaries, and even more complex when an amorphous phase is present, such as in semicrystalline materials, see Figure 12.1g. For semicrystalline organic semiconductors, the optoelectronic properties are affected by not only the crystal packing structure but also the number of repeat units (e.g. monomers) involved in π–π interactions, defining the conjugation length, and thus, affected by size and shape of the ordered domains [22–24]. In addition, ordered crystalline domains offer a fast pathway for charge transport, making the determination of their size and content of great value [25]. As represented schematically in Figure 12.1g, in semicrystalline thin films, the microstructure is composed of amorphous and crystalline domains. In these films, in addition to the crystal packing structure and defects, other microstructural features become important, such as the size, shape, and distribution of crystallites. WAXS is used to reveal valuable information about the degree of crystallinity and the crystalline correlation length; a parameter

Microstructural Characterization of Conjugated Organic Semiconductors


FIGURE 12.1  (a–b) Molecular structure of the small-molecule Cl 2-NDI in α-phase (a) and β-phase (b) [21]. (c–d) Crystal packing structures of α-phase (c) with herringbone packing structure and β-phase (d) with 2D brick-wall packing in the (ab)-plane. The π–π stacking for α-phase has d-spacing of 3.27 Å, corresponding to slipping angle of 62°, while π–π stacking d-spacing for β-phase is 3.29 Å and 3.32 Å, corresponding to slipping angles of 41° and 32°, respectively (−CH2C3F7 is omitted for clarity). (e–f) Crystal packing structure of α-phase (e) and β-phase (f) in the (ac)-plane with tilt angles of ~ 67° and ~ 90° on the substrate, respectively. Adapted with permission [21]. Copyright 2015, Nature Publishing Group. (g–h) Schematics represent the plane-view microstructure of neat semicrystalline polymer film (g) consisting of crystalline and amorphous domains, and polymer: fullerene blend film (h) consisting of pure polymer (crystalline and amorphous), pure fullerene, and a mixed phase.

related to the average size of ordered regions and contains information about the disorder and defects within crystallites [8, 26]. Through careful analysis of the shape (i.e. broadening) of the scattering profile, the contribution from the crystalline domain sizes and defects can be decoupled [27, 28]. Moreover, scattering intensity from WAXS gives information about the degree of crystallinity, which determines the fractional amount of crystalline phase in the film. In neat films, small-angle X-ray scattering (SAXS) makes use of the contrast between the amorphous and crystalline domains to give information about their average domain size [8, 26]. For composite systems, such as in BHJ blends of two organics, as represented in Figure 12.1h for a polymer: fullerene blend film, the microstructure is composed of pure polymer domains (crystalline and amorphous), pure fullerene domains, and a mixed phase of the polymer and fullerene. In such films, microstructural characterization involves the evaluation of additional features such as the average domain size of pure phases and the degree of mixing (i.e. volume percentage of the mixed phase). In this case, WAXS can be used to decouple the contribution of each component/ phase into the scattering profile and to measure the volume fraction of each component/phase, while SAXS can be used to calculate the average domain size relying on the electron density contrast between different phases/components [29, 30]. Typically, organic semiconductors are weakly ordered requiring a long exposure and/or high flux in order to efficiently probe the microstructure. Both lab-based and synchrotron-based X-rays have been used for organic thin film characterization, however, lab-based X-rays are limited to strongly scattering films and thick powder samples due to low flux and intensity of the X-ray beam [31]. On the other hand, synchrotron-based X-rays provide high flux, intense, and collimated X-ray beams allowing characterization of weakly ordered samples such as most organic thin films; which is the focus in this chapter


Conjugated Polymers

[31]. One should pay attention to the possible structural changes due to long exposure to the X-ray beam (i.e. beam damage) that may lead to misleading results. Although the high flux from synchrotron X-rays allows minimal exposure, it can accelerate the beam damage. Investigating the microstructure as a function of exposure time and X-ray dose helps to define the maximum suitable exposure before beam damage can occur, below which data should be collected. This practice is used to ensure that the collected data is not affected by beam damage. In this chapter, we provide an overview of the contribution of X-ray scattering to microstructure characterization of organic thin films used in optoelectronic applications. Firstly, we begin with explaining the fundamentals of X-ray scattering, covering grazing incidence wide-angle x-ray scattering (GIWAXS) for probing the molecular order and orientation at the small-length-scale and grazing incidence small-angle X-ray scattering (GISAXS) for probing the phase separation at the large-scale. We then provide selected examples to highlight the applications of GIWAXS and GISAXS in probing the microstructure for both static thin films and in situ during thin film formation.

12.2 Fundamentals of X-ray Scattering In this section, we provide a brief review of the fundamentals of X-ray scattering techniques used in the characterization of organic semiconductors with a focus on thin film characterization. Firstly, we cover the fundamentals of wide-angle X-ray scattering (WAXS) and its application to probe the molecular order and orientation at the small-length-scale. Then, we explain the fundamentals of small-angle x-ray scattering (SAXS) in probing the phase separation at the large-scale.

12.2.1 Wide-Angle X-ray Scattering (WAXS) WAXS is typically sensitive to the ordered (crystalline) regions and reveals information of the microstructural features at the small-length scale (molecular scale) such as the intermolecular and intermolecular repeating building blocks in ordered domains. X-ray scattering from these molecular-scale features occurs at wide scattering angles, thus named WAXS. It is often used in the grazing incidence geometry, in order to increase the scattering signal, to characterize the microstructure of organic thin films through resolving the crystal structure and texture orientation and estimating the relative degree of crystallinity and the crystalline correlation length [8, 30]. In the following subsections, we cover the principles of the characterization of these microstructural parameters by WAXS. Basics of WAXS WAXS is an X-ray scattering technique focused on analyzing the Bragg diffraction at wide angles corresponding to microstructural features with length scales from Å to a few nm using high-energy (or hard) X-rays. In WAXS, an incident X-ray beam impinges onto the sample, where a portion of the beam is diffracted by the repetitive (periodic) crystallographic planes. The diffraction angle (2θ) is related to the spacing of the periodic planes (d) and wavelength (λ) of the incident beam through Bragg’s law of diffraction (with its general formula: nl  2d sin q ; where [n = 1, 2, 3,…] is the order of reflection) [32]. One powerful tool for predicting crystallographic scattering patterns uses construction of reciprocal space and Ewald’s sphere [32]. The Ewald sphere is a geometric representation of the relationship between the incident beam vector, the diffracted vectors, the reciprocal lattice vector, and the angles of diffraction, with the recognition that the diffracted vectors in Ewald sphere are associated with an elastic scattering. As shown in Figure 12.2A, parallel crystallographic planes in real space can be represented as points in reciprocal space where each point corresponds to a set of lattice planes labeled with Miller indices (hkl).  The reciprocal space vector (q ) direction is normal to the lattice planes and its magnitude is defined by the inter-planar spacing of the lattice planes as q = 2π / d , see Figure 12.2A. The Ewald sphere can be constructed by a sphere with a radius related to the wavelength of the incident X-ray beam by 2p / l


Microstructural Characterization of Conjugated Organic Semiconductors


FIGURE 12.2  (A) Schematic representation of the diffraction condition; the reciprocal-space is represented as dots correspond to the Fourier transform of the crystallographic planes in real space (represented as parallel lines); the Ewald sphere, represented as circle  (sphere   in 3D) with a radius = 2π / λ , constitutes all wave-vectors of the incident and diffracted X-ray beams ( ki and ko , respectively); diffraction occurs when the reciprocal lattice points intersect  the Ewald sphere where the wave-vector in reciprocal space ( q , known    as scattering vector) for each diffraction spot is determined by the incident and diffracted wave-vectors by ( q  ko  ki ) [32]. (B) Schematic representation of grazing incidence geometry of WAXS experiment where shallow incident angles (αi 10%) polymer paired with multiple fullerene acceptors108,112, 11.7%-efficiency devices processed with hydrocarbon solvents83, and over 11%-efficiency devices based on a regioregular polymer85. Inspired by Collins’s P-SoXS work in organic transistors and initial work on all-polymer OSCs, Tumbleston et al.59 discovered that molecular orientation relative to donor/acceptor (D/A) interfaces can be an important parameter that controls the performance of polymer:fullerene OSCs. Shown in Figure 13.12a is the schematic of a pair of PNDT-DTBT-based polymers and their processing conditions. The fluorinated polymer shows a face-on orientation while the nonfluorinated polymer shows a slightly edgeon orientation relative to the D/A interface. They found that both device short-circuit current density (Jsc) and fill factor (FF) are strongly correlated with the degree of molecule orientation (DMO) (see Figure 13.12b), which can be determined from the energy-dependent P-SoXS data. By comparing three sets of donor/acceptor copolymers, this study concluded that preferential face-on polymer orientation relative to the discrete D/A interface is critical to achieving high-efficiency OSC devices. This work thus established a strong relationship between the DMO and device parameters (Jsc, FF). Scattering anisotropy has been observed in many polymer:polymer solar cell systems54,79,113. In a high-performance all-polymer system54, the addition of a trace amount of high boiling-point solvent additive can simultaneously optimize both average domain purity and orientational ordering, which collectively lead to a much higher device FF of 66.8% in comparison to that of the reference additive-free device. Such orientational ordering, along with molecular-scale mixing driven by a fundamental miscibility of the materials are very important to the processes of charge separation and recombination in devices124–126. Soft X-ray scattering is a unique tool

FIGURE 13.12  (a) Chemical structures of two PNDT-DTBT polymers where the X atoms are either fully hydrogen or fully fluorine); (b) Device FF and J sc   as a function of the relative degree of molecular orientation (DMO). Tumbleston et al. Nature Photonics  2014, 8, 385– 391.

Soft X-Ray Scattering Characterization of Polymer Semiconductors


to assess such orientation and the role (if any) relative molecular orientation at an interface plays in the performance of devices can now be assessed although full interpretation of the information encoded in the anisotropic scattering remains challenging until better analysis tools have been developed. Aided by R-SoXS results, structure–morphology–performance relationships can be established7. These characterizations allow the community, for example, to demonstrate strong correlations between a number of morphology and performance parameters. A notable example is the frequently observed monotonic and even linear ISI1/2–FF relations. In a model polymer:fullerene system (Figure 13.13), Yan and co-workers demonstrated that addition of a hydrocarbon solvent additive PN to the host solvent TMB can significantly decrease the domain spacing and increase the average domain purity, together leading to a record-efficiency polymer:fullerene OSC with a certified efficiency of 11.5%. The device FF showed a strong relation with ISI1/2, which scales with the standard deviation of the composition and is impacted by the average domain purity. In 2015, Mukherjee et al.123 for the first time established a mostly linear relation between FF and ISI1/2 in organic solar cells based on small molecule:PCBM blends. In early 2017, Ye et al.63 observed a strictly linear relation between ISI1/2 and FF in five polymer:fullerene blends, where the polymers used are structurally similar. This relation has been verified in many subsequent studies67,85,90,92,127–129 and inspires new understanding of structure–function relations. Based on the linear ISI1/2–FF relation assumption, Ye et al. further developed a quantitative model7 that can relate the effective Flory–Huggins interaction parameter χ to ISI1/2 and device FF (see Figure 13.14) in a model system of Poly[​N-9'-​hepta​decan​yl-2,​7-car​bazol​e-alt​-5,5-​(4',7​'-di-​2-thi​enyl-​2',1'​,3'-b​enzot​ hiadi​azole​)] (PCDTBT):PCBM by using a combination of probe tools including STXM, R-SoXS at the Advanced Light Source (ALS), and secondary ion mass spectroscopy (SIMS). This relation is also delineated across numerous high and low performing materials systems, including fullerene and nonfullerene acceptors. Ye et al. showed that molecular interactions, as encoded in this measurable χ parameter, control the achievable OSC morphology and monotonically relate to ISI. Only unfavorable interactions (high χ) lead to the formation of a favorable OSC morphology with high purity of the mixed domains and high performance. Crucially, the experimental observations are quite consistent with the computational results from atomic molecular dynamics simulations but are only valid if the acceptor concentration remains above the percolation threshold9. In some cases, highly complex and multilength-scale morphology is observed55,64,65,68,79,84,114,130,131 and the morphology analysis is not that straightforward in these OSCs. In a model polymer:fullerene system where four different molecular weight PBDTTPDs are used, Mukherjee et al.65 revealed that the root-mean-square composition variance of the smallest domain, i.e. its standard deviation, is the most

FIGURE 13.13  (a) Chemical structures of TMB, PN and PffBT4T-C9 C13  . (b) R-SoXS profiles of PffBT4T-C9 C13 :PC71 BM-based films processed from TMB (normalized ISI of 0.74) or TMB+PN (normalized ISI of 1.00). Zhao et al. Nature Energy 2016, 1, 15027.


Conjugated Polymers

FIGURE 13.14  (a) Device FF of PCDTBT:PC71 BM films cast at room temperature from the fast-drying solvent chloroform, subsequently annealed for a constant time (10 minutes) as a function of annealing temperature as indicated. (b) Device FF of PCDTBT:PC71 BM films as a function of χ aa . Squares and spheres represent Batch 1 and Batch 2, respectively. The solid curved lines are directly derived from a [ FF  k f0  f1 f2  f0   const ] model that is based on the binodal ϕ 1  and ϕ 2  encoded by χ aa . (c) Lorentz-corrected R-SoXS profiles of PCDTBT:PC71 BM films with χ  = 1.19 (annealed at 90 °C ) and χ  = 0.54 (annealed at 200 °C ), i.e. in the two-phase and one-phase region, respectively. Only a featureless profile with a sloping background and low scattering intensity is observed once devices are annealed at 200 °C . (d) Plot of relative ISI acquired from R-SoXS against χ aa  for PCDTBT (Batch 1):PC71 BM films. ISI of the film processed at 20° C is set to 1. The solid curved line is directly derived from a quantitative model. Ye et al., Nature Materials  2018, 17, 253– 260.

c. ritical factor that determines the device Jsc and FF of polymer:fullerene blends, while the overall ISI1/2 has no clear relation with these performance metrics. This can be qualitatively understood as the largest domains observed had a length scale comparable to or larger than the film thickness. Recent studies by Ye et al. indicate that this observation might be general and that the characteristics of smallest domains critically determine device performance of both spin-coated and blade-coated nonfullerene OSCs51,53. As illustrated in Figure 13.15, Ye et al. reported a method to precisely control the multi-length scale morphology by gradually reducing the concentration of a green solvent additive used in blade-coated films. It is found that a lower amount of the halogen-free solvent additive diphenyl ether (DPE) results in a higher Jsc and FF. As a result, the additive-free eco-friendly all-polymer OSCs show a significantly higher performance in comparison to the OSCs with additives. Using R-SoXS and electrical methods, they revealed that the device Jsc and FF can be correlated to the increased mean-square composition variations at the smallest length-scales, interpreted to be an increased volume fraction (φs) of the smallest

Soft X-Ray Scattering Characterization of Polymer Semiconductors


FIGURE 13.15  (a) Schematic of blade-coating all-polymer OSC in air and chemical structures of the polymer semiconductors used; (b) Normalized and Lorentz corrected R-SoXS profiles of blade-coated OSCs with varying volume amount of the additive DPE; (c) Relations of device J sc   and FF with the volume fraction of the smallest domains. Adapted with permission from Ye et al., Adv. Funct. Mater. , 2017, 27, 1702016.

domains with fixed purity. These studies together suggest that high volume fraction or composition variation of smallest domains is a critical requirement for achieving high-efficiency OSCs with multilength scale morphology. This new finding might lead to the prediction of the multi-length scale morphology and its relation to the device performance metrics. As the field of OSCs continues to grow as new fullerene-free materials are created to improve device efficiency and stability132, R-SoXS again shows great promise in characterizing the complex morphology of nonfullerene systems. Jsc and FF of nonfullerene OSCs can be often explained by the morphological parameters extracted from R-SoXS in a wider range of nonfullerene OSC devices7,90,92,101,103,128,133,134. These include many high-efficiency all-polymer solar cells54,84,95,105,109,135 and the most recent record-efficiency (~11%) fullerene-free OSCs by scalable blade-coating67.

13.3.3 Soft X-Ray Scattering of Multi-Component Semiconducting Polymer Blends R-SoXS inherently brings chemical sensitivity and, as such, can be used to investigate multi-component systems, where the pair-wise contrast of the different components can be tuned by the selection of the incident X-ray energy. Compared with traditional OSCs based on binary donor:acceptor blends, OSCs integrating multiple donor or acceptor materials in one photoactive layer have emerged as a promising strategy to improve device performance136-138. The third component plays versatile roles in ternary-blend OSCs, including complementary light harvesting, facilitating exciton dissociation, enhancing charge transport, and optimizing the film morphology. In 2016, Bo’s group139 employed a fullerene derivative (PC71BM) and a nonfullerene small molecule (3,9-​bis(2​-meth​ylene​-(3-(​1,1-d​icyan​ometh​ylene​)-ind​a none​))-5,​5,11,​11-te​t raki​s(4-h​exylp​henyl​)-dit​hieno​ [2,3-​d:2’,​3’-d’​]-s-i​ndace​no[1,​2-b:5​,6-b’​]dith​iophe​ne) (ITIC) as dual acceptors to fabricate a new type of ternary OSC device for the first time. A remarkably higher PCE of 10.4% was obtained in the optimal ternary devices, which is ~40% higher than the ITIC-based and PCBM-based control binary devices.


Conjugated Polymers

Ternary blend OSCs with different ratios of the ITIC:PC71BM were analyzed with R-SoXS. The large shift of R-SoXS peaks suggests that the domain spacing of the PPBDTBT:ITIC devices are significantly decreased after adding PC71BM. This also tracks the changes of device Jsc. Additionally, they found that the device FF varies positively with the increasing weight ratio of PC71BM in acceptor materials. The influence of the third component on the film morphology in ternary blends was studied with R-SoXS. With increasing content of PC71BM, the relative mean square composition variations of these blends are 0.09, 0.26, 0.50, 0.68, 0.91, and 1, respectively. The higher purity of the mixed domains can reduce bimolecular recombination, which can explain the changes of FF in the OSCs. The mean-square composition variation FF data sets in this study again point to the generality of the mostly linear relations observed in previous studies and the importance of domain purity. The current R-SoXS technique is focused on the in-plane morphology probed by the transmission geometry for the sake of easy setup and straightforward data analysis. Grazing incidence R-SoXS (GI-RSoXS)35,140 is a new type of soft X-ray scattering that can reveal both vertical and lateral morphology of layered polymer films. In the case of multi-component blends, GI-RSoXS is even more powerful. The strong dependence of optical constants of organic/polymeric materials in the energy range of 260–300 eV allows for construction of the 3D morphology. In 2016, Müller-Buschbaum et al.141 applied GI-RSoXS to reveal the 3D morphology evolution in the benchmark PCPDTBT:PCBM blend films with an added photosensitizer of perylene diimide (PDI) in the presence and absence of a solvent additive. They were able to reveal the morphological pictures of this multi-component system on the presence of octanedithiol (ODT) and PDI. The GI-RSoXS results suggest that additive-free blend films show only vertical phase separation while films processed with solvent additive also show lateral phase separation. They showed that the addition of the nonfullerene sensitizer PDI tends to facilitate lateral phase separation between PCPDTBT and PCBM by expelling PCBM from the amorphous polymer matrix, as demonstrated in Figure 13.16. Together, this work shows that the sizes and compositions of different phases and phase separation mechanisms of such a complex system could be clearly characterized with this GI-RSoXS tool. To date, R-SoXS has also contributed significantly to the fundamental understanding of alloy formation and phase separation of other types of ternary OSCs including polymer:small molecule:fullerene ble nds131,142–145, polymer:polymer:fullerene blends146–151, polymer:polymer:polymer blends152 and fullerenefree blends153–155. In addition, R-SoXS has been successfully applied to more complex systems, including quaternary blend OSCs156. Overall, this fundamental understanding from R-SoXS may stimulate new breakthroughs on the complex morphology of OSCs and related electronics.

13.4 Conclusions and Outlook In this chapter, we reviewed the basic principles and morphological parameters that dictate the morphology and applications of soft X-ray scattering in the field of polymer semiconductors. Specifically,

FIGURE 13.16  Energy-dependent GI-RSoXS patterns and graphical representations of the morphology of the OSC based on a multi-component blend PCPDTBT:PCBM:PDI. Adapted with permission from Schaffer et al., Polymer , 2016, 105, 357– 367.

Soft X-Ray Scattering Characterization of Polymer Semiconductors


resonant soft X-ray scattering combines soft X-ray spectroscopy with small angle X-ray scattering and thus offers statistical information for organic materials over a large length scale range from ~10 nm to ~1 μm. Due to the unique chemical and orientational sensitivity near the carbon absorption edge, R-SoXS can quantify the domain characteristics (spacing and average purity) and orientational ordering relative to the interface can be probed with P-SoXS. Particularly, measurements of composition variations which are sensitive to domain purity are not easily obtainable with other techniques. In addition, GI-RSoXS is capable of obtaining both vertical and lateral information of polymer blends. In particular, we have discussed the most recent progress in organic electronics with the aid of soft X-ray scattering. It is clear that these quantified parameters from soft X-ray scattering measurements allow the community to establish many critical performance–morphology, processing–morphology, and miscibility–function relations in the emerging ternary-blend and fullerene-free OSC devices, in addition to the conventional polymer:fullerene OSC devices. These relations would eventually enable the prediction of best processing parameters and material combinations for the manufacture of highperformance OSCs. Despite the great achievements from soft X-ray scattering characterizations of polymer semiconductors, there are still some limitations for conventional R-SoXS characterizations. First, real-time information can currently not be obtained from conventional R-SoXS measurements, and there are very limited studies on the development of novel soft X-ray scattering tools under environmental control (solution, temperature, electric field, etc.). The in-situ/in-operando measurements41,157–161 on the basis of soft X-ray scattering will allow profound progress in the knowledge of morphology evolution of polymer semiconductors during film formation, post-deposition process, and aging/stability assessment of organic devices. Second, the scattering vector q probed with the typical R-SoXS configuration is generally below 1 nm-1. In order to construct a more comprehensive morphological picture, higher q range (1-3 nm-1) is needed for some polymer:nonfullerene molecule systems where the domain size is lower than 10 nm. In this regard, developing super high-q scattering geometries or utilizing other resonant energies, such as the oxygen162, nitrogen, or fluorine K-edge range is anticipated8. Third, conventional R-SoXS measurements only provide in-plane (lateral) information while out-of-plane and phase information are missing. X-ray ptychography163–168, actively constructed at the Advanced Light Source, might be highly beneficial to the phase retrieval of R-SoXS data and visualization of 3D morphology. Data in reciprocal space could be inverted to real space maps or even tomograms of the device morphology. Beyond the successful and rapid application in polymer semiconductors, soft X-ray scattering techniques can also probe the morphology of other soft matter such as block copolymers33,61,169–171, membranes172, ferroelectric semiconductors173, liquid crystals174–177, metal–organic frameworks, and perovskites178. The continued development of soft X-ray scattering can improve the precise measurement and correlation of morphology to device processes, allowing for the identification of processing routes to control and optimize the hierarchical structure within these devices. We expect that the methods described here can inspire the application of this tool to other emerging important material classes.

Acknowledgments The authors would like to thank the support for this work from the ONR grants N000141512322 and N000141712204, and NSF INFEWS grant CBET 1639429. Samuel J. Stuard also acknowledges support by the SEAS program under NSF grant DGE-1633587. Beamline at the Advanced Light Source, Lawrence Berkeley National Laboratory is supported by the Director of the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. We gratefully acknowledge all the researchers who contributed to the advances, our thinking, and applications of R-SoXS. Specifically, we thank Eliot Gann and Brian Collins for fruitful discussions and critical comments over the years and particularly during the revisions of this chapter.


Conjugated Polymers


1. Beaujuge, P. M.; Fréchet, J. M. J. Molecular design and ordering effects in π-functional materials for transistor and solar cell applications. Journal of the American Chemical Society 2011, 133, 20009–20029. 2. Henson, Z. B.; Mullen, K.; Bazan, G. C. Design strategies for organic semiconductors beyond the molecular formula. Nature Chemistry 2012, 4, 699–704. 3. Jinno, H.; Fukuda, K.; Xu, X. M.; Park, S.; Suzuki, Y.; Koizumi, M.; Yokota, T.; Osaka, I.; Takimiya, K.; Someya, T. Stretchable and waterproof elastomer-coated organic photovoltaics for washable electronic textile applications. Nature Energy 2017, 2, 780–785. 4. Heeger, A. J. Semiconducting polymers: the third generation. Chemical Society Reviews 2010, 39, 2354–2371. 5. Treat, N. D.; Westacott, P.; Stingelin, N. The power of materials science tools for gaining insights into organic semiconductors. Annual Review of Materials Research 2015, 45, 459–490. 6. Dou, L.; You, J.; Hong, Z.; Xu, Z.; Li, G.; Street, R. A.; Yang, Y. 25th anniversary article: a decade of organic/polymeric photovoltaic research. Advanced Materials 2013, 25, 6642–6671. 7. Ye, L.; Hu, H.; Ghasemi, M.; Wang, T.; Collins, B. A.; Kim, J.-H.; Jiang, K.; Carpenter, J.; Li, H.; Li, Z.; McAfee, T.; Zhao, J.; Chen, X.; Lai, J. Y. L.; Ma, T.; Bredas, J.-L.; Yan, H.; Ade, H. Quantitative relations between interaction parameter, miscibility and function in organic solar cells. Nature Materials 2018, 17, 253–260. 8. Jiao, X.; Ye, L.; Ade, H. Quantitative morphology-performance correlations in organic solar cells: insights from soft X-ray scattering. Advanced Energy Materials 2017, 7, 1700084. 9. Ye, L.; Collins, B. A.; Jiao, X.; Zhao, J.; Yan, H.; Ade, H. Miscibility-function relations in organic solar cells: significance of optimal miscibility in relation to percolation. Advanced Energy Materials 2018, 8, 1703058. 10. Huang, Y.; Kramer, E. J.; Heeger, A. J.; Bazan, G. C. Bulk heterojunction solar cells: morphology and performance relationships. Chemical Reviews 2014, 114, 7006–7043. 11. Lu, L.; Zheng, T.; Wu, Q.; Schneider, A. M.; Zhao, D.; Yu, L. Recent advances in bulk heterojunction polymer solar cells. Chemical Reviews 2015, 115, 12666–12731. 12. Brabec, C. J.; Heeney, M.; McCulloch, I.; Nelson, J. Influence of blend microstructure on bulk heterojunction organic photovoltaic performance. Chemical Society Reviews 2011, 40, 1185–1199. 13. Menke, S. M.; Ran, N. A.; Bazan, G. C.; Friend, R. H. Understanding energy loss in organic solar cells: toward a new efficiency regime. Joule 2017, 2, 1–11. 14. Kuei, B.; Gomez, E. D. Chain conformations and phase behavior of conjugated polymers. Soft Matter 2016, 13, 49–67. 15. Jackson, N. E.; Savoie, B. M.; Marks, T. J.; Chen, L. X.; Ratner, M. A. The next breakthrough for organic photovoltaics? The Journal of Physical Chemistry Letters 2015, 6, 77–84. 16. Treat, N. D.; Chabinyc, M. L. Phase separation in bulk heterojunctions of semiconducting polymers and fullerenes for photovoltaics. Annual Review of Physical Chemistry 2014, 65, 59–81. 17. Patel, S. N.; Glaudell, A. M.; Peterson, K. A.; Thomas, E. M.; O'Hara, K. A.; Lim, E.; Chabinyc, M. L. Morphology controls the thermoelectric power factor of a doped semiconducting polymer. Science Advances 2017, 3, e1700434. 18. Zhang, Q.; Sun, Y.; Xu, W.; Zhu, D. Organic thermoelectric materials: emerging green energy materials converting heat to electricity directly and efficiently. Advanced Materials 2014, 26, 6829–6851. 19. Russ, B.; Glaudell, A.; Urban, J. J.; Chabinyc, M. L.; Segalman, R. A. Organic thermoelectric materials for energy harvesting and temperature control. Nature Reviews Materials 2016, 1, 16050. 20. Di, C.-A.; Xu, W.; Zhu, D. Organic thermoelectrics for green energy. National Science Review 2016, 3, 269–271. 21. Jansen-van Vuuren, R. D.; Armin, A.; Pandey, A. K.; Burn, P. L.; Meredith, P. Organic photodiodes: the future of full color detection and image sensing. Advanced Materials 2016, 28, 4766–4802.

Soft X-Ray Scattering Characterization of Polymer Semiconductors


22. Gong, X.; Tong, M.; Xia, Y.; Cai, W.; Moon, J. S.; Cao, Y.; Yu, G.; Shieh, C. L.; Nilsson, B.; Heeger, A. J. High-detectivity polymer photodetectors with spectral response from 300 nm to 1450 nm. Science 2009, 325, 1665–1667. 23. Rim, Y. S.; Bae, S.-H.; Chen, H.; De Marco, N.; Yang, Y. Recent progress in materials and devices toward printable and flexible sensors. Advanced Materials 2016, 28, 4415–4440. 24. Abbaszadeh, D.; Kunz, A.; Wetzelaer, G. A. H.; Michels, J. J.; Cra c̆ iun, N. I.; Koynov, K.; Lieberwirth, I.; Blom, P. W. M. Elimination of charge carrier trapping in diluted semiconductors. Nature Materials 2016, 15, 628–633. 25. Venkateshvaran, D.; Nikolka, M.; Sadhanala, A.; Lemaur, V.; Zelazny, M.; Kepa, M.; Hurhangee, M.; Kronemeijer, A. J.; Pecunia, V.; Nasrallah, I.; Romanov, I.; Broch, K.; McCulloch, I.; Emin, D.; Olivier, Y.; Cornil, J.; Beljonne, D.; Sirringhaus, H. Approaching disorder-free transport in highmobility conjugated polymers. Nature 2014, 515, 384–388. 26. Nikolka, M.; Nasrallah, I.; Rose, B.; Ravva, M. K.; Broch, K.; Sadhanala, A.; Harkin, D.; Charmet, J.; Hurhangee, M.; Brown, A.; Illig, S.; Too, P.; Jongman, J.; McCulloch, I.; Bredas, J.-L.; Sirringhaus, H. High operational and environmental stability of high-mobility conjugated polymer field-effect transistors through the use of molecular additives. Nature Materials 2017, 16, 356–362. 27. Sirringhaus, H. 25th anniversary article: organic field-effect transistors: the path beyond amorphous silicon. Advanced Materials 2014, 26, 1319–1335. 28. Xu, J.; Wang, S.; Wang, G.-J. N.; Zhu, C.; Luo, S.; Jin, L.; Gu, X.; Chen, S.; Feig, V. R.; To, J. W. F.; Rondeau-Gagné, S.; Park, J.; Schroeder, B. C.; Lu, C.; Oh, J. Y.; Wang, Y.; Kim, Y.-H.; Yan, H.; Sinclair, R.; Zhou, D.; Xue, G.; Murmann, B.; Linder, C.; Cai, W.; Tok, J. B. H.; Chung, J. W.; Bao, Z. Highly stretchable polymer semiconductor films through the nanoconfinement effect. Science 2017, 355, 59–64. 29. Oh, J. Y.; Rondeau-Gagné, S.; Chiu, Y.-C.; Chortos, A.; Lissel, F.; Wang, G.-J. N.; Schroeder, B. C.; Kurosawa, T.; Lopez, J.; Katsumata, T.; Xu, J.; Zhu, C.; Gu, X.; Bae, W.-G.; Kim, Y.; Jin, L.; Chung, J. W.; Tok, J. B. H.; Bao, Z. Intrinsically stretchable and healable semiconducting polymer for organic transistors. Nature 2016, 539, 411–415. 30. Noriega, R.; Rivnay, J.; Vandewal, K.; Koch, F. P. V.; Stingelin, N.; Smith, P.; Toney, M. F.; Salleo, A. A general relationship between disorder, aggregation and charge transport in conjugated polymers. Nature Materials 2013, 12, 1038–1044. 31. Carpenter, J. H.; Hunt, A.; Ade, H. Characterizing morphology in organic systems with resonant soft X-ray scattering. Journal of Electron Spectroscopy and Related Phenomena 2015, 200, 2–14. 32. McNeill, C. R.; Ade, H. Soft X-ray characterisation of organic semiconductor films. Journal of Materials Chemistry C 2013, 1, 187–201. 33. Liu, F.; Brady, M. A.; Wang, C. Resonant soft X-ray scattering for polymer materials. European Polymer Journal 2016, 81, 555–568. 34. Mukherjee, S.; Herzing, A. A.; Zhao, D. L.; Wu, Q. H.; Yu, L. P.; Ade, H.; DeLongchamp, D. M.; Richter, L. J. Morphological characterization of fullerene and fullerene-free organic photovoltaics by combined real and reciprocal space techniques. Journal of Materials Research 2017, 32, 1921–1934. 35. Gann, E.; Watson, A.; Tumbleston, J. R.; Cochran, J.; Yan, H.; Wang, C.; Seok, J.; Chabinyc, M.; Ade, H. Topographic measurement of buried thin-film interfaces using a grazing resonant soft X-ray scattering technique. Physical Review B 2014, 90, 245421. 36. Su, G. M.; Cordova, I. A.; Brady, M. A.; Prendergast, D.; Wang, C. Combining theory and experiment for X-ray absorption spectroscopy and resonant X-ray scattering characterization of polymers. Polymer 2016, 99, 782–796. 37. Wang, C.; Hexemer, A.; Nasiatka, J.; Chan, E. R.; Young, A. T.; Padmore, H. A.; Schlotter, W. F.; Lüning, J.; Swaraj, S.; Watts, B.; Gann, E.; Yan, H.; Ade, H. Resonant soft X-ray scattering of polymers with a 2D detector: initial results and system developments at the advanced light source. IOP Conference Series: Materials Science and Engineering 2010, 14, 012016.


Conjugated Polymers

38. Collins, B. A.; Bokel, F. A.; DeLongchamp, D. M. Organic photovoltaic morphology. In C. Brabec, U. Scherf, V. Dyakonov, eds., Organic Photovoltaics. Wiley-VCH Verlag GmbH & Co. KGaA 2014; pp 377–420. 39. Chen, W.; Nikiforov, M. P.; Darling, S. B. Morphology characterization in organic and hybrid solar cells. Energy & Environmental Science 2012, 5, 8045–8074. 40. Liu, F.; Gu, Y.; Jung, J. W.; Jo, W. H.; Russell, T. P. On the morphology of polymer-based photovoltaics. Journal of Polymer Science Part B: Polymer Physics 2012, 50, 1018–1044. 41. Richter, L. J.; DeLongchamp, D. M.; Amassian, A. Morphology development in solution-processed functional organic blend films: an in situ viewpoint. Chemical Reviews 2017, 117, 6332–6366. 42. Rivnay, J.; Mannsfeld, S. C. B.; Miller, C. E.; Salleo, A.; Toney, M. F. Quantitative determination of organic semiconductor microstructure from the molecular to device scale. Chemical Reviews 2012, 112, 5488–5519. 43. Liu, F.; Gu, Y.; Shen, X.; Ferdous, S.; Wang, H.-W.; Russell, T. P. Characterization of the morphology of solution-processed bulk heterojunction organic photovoltaics. Progress in Polymer Science 2013, 38, 1990–2052. 44. DeLongchamp, D. M.; Kline, R. J.; Fischer, D. A.; Richter, L. J.; Toney, M. F. Molecular characterization of organic electronic films. Advanced Materials 2011, 23, 319–337. 45. Attwood, D.; Sakdinawat, A. X-Rays and Extreme Ultraviolet Radiation: Principles and Applications. Cambridge University Press 2016. 46. Attwood, D. Soft X-Rays and Extreme Ultraviolet Radiation: Principles and Applications. Cambridge University Press 2000. 47. Gann, E.; Young, A. T.; Collins, B. A.; Yan, H.; Nasiatka, J.; Padmore, H. A.; Ade, H.; Hexemer, A.; Wang, C. Soft X-ray scattering facility at the advanced light source with real-time data processing and analysis. Review of Scientific Instruments 2012, 83, 045110. 48. Stribeck, N. X-Ray Scattering of Soft Matter. Springer 2007. 49. Fang, S. J.; Haplepete, S.; Chen, W.; Helms, C. R.; Edwards, H. Analyzing atomic force microscopy images using spectral methods. Journal of Applied Physics 1997, 82, 5891–5898. 50. Gann, E. Using resonant soft X-rays to reveal internal organic thin film morphology. PhD Thesis 2013, North Carolina State University, Raleigh NC. 51. Ye, L.; Zhao, W. C.; Li, S. S.; Mukherjee, S.; Carpenter, J. H.; Awartani, O.; Jiao, X. C.; Hou, J. H.; Ade, H. High-efficiency nonfullerene organic solar cells: critical factors that affect complex multilength scale morphology and device performance. Advanced Energy Materials 2017, 7, 1602000. 52. Li, W.; Ye, L.; Li, S.; Yao, H.; Ade, H.; Hou, J. A high efficiency organic solar cell enabled by strong intramolecular electron push-pull effect of non-fullerene acceptor. Advanced Materials 2018, 30, 1707170. 53. Ye, L.; Xiong, Y.; Li, S.; Ghasemi, M.; Balar, N.; Turner, J.; Gadisa, A.; Hou, J.; O'Connor, B. T.; Ade, H. Precise manipulation of multi-length scale morphology and its influence on eco-friendly printed all-polymer solar cells. Advanced Functional Materials 2017, 27, 1702016. 54. Ye, L.; Jiao, X. C.; Zhou, M.; Zhang, S. Q.; Yao, H. F.; Zhao, W. C.; Xia, A. D.; Ade, H.; Hou, J. H. Manipulating aggregation and molecular orientation in all-polymer photovoltaic cells. Advanced Materials 2015, 27, 6046–6054. 55. Ye, L.; Xiong, Y.; Yao, H.; Gadisa, A.; Zhang, H.; Li, S.; Ghasemi, M.; Balar, N.; Hunt, A.; O’Connor, B. T.; Hou, J.; Ade, H. High performance organic solar cells processed by blade coating in air from a benign food additive solution. Chemistry of Materials 2016, 28, 7451–7458. 56. Alqahtani, O.; Babics, M.; Gorenflot, J.; Savikhin, V.; Ferron, T.; Balawi, A.; Paulke, A.; Kan, Z.; Pope, M.; Clulow, A. J.; Wolf, J.; Burn, P. L.; Ian, G.; Dieter, N.; Toney, M. F.; Laquai, F.; Beaujuge, P. M.; Collins, B. A. Mixed domains enhance charge generation and extraction in bulk-heterojunction solar cells with small-molecule donors. Advanced Energy Materials 2018, 8, 1702941 . 57. Crist, B.; Morosoff, N. Small-angle X-ray scattering of semicrystalline polymers. II. Analysis of experimental scattering curves. Journal of Polymer Science: Polymer Physics Edition 1973, 11, 1023–1045.

Soft X-Ray Scattering Characterization of Polymer Semiconductors


58. Khambatta, F. B. Small-angle X-ray and small-angle light scattering studies of the morphology of polymer blends. PhD Thesis 1976, University of Massachusetts, Amherst MA. 59. Tumbleston, J. R.; Collins, B. A.; Yang, L.; Stuart, A. C.; Gann, E.; Ma, W.; You, W.; Ade, H. The influence of molecular orientation on organic bulk heterojunction solar cells. Nature Photonics 2014, 8, 385–391. 60. Roe, R.-J. Methods of X-Ray and Neutron Scattering in Polymer Science. Oxford University Press 2000. 61. Ferron, T.; Pope, M.; Collins, B. A. Spectral analysis for resonant soft X-ray scattering enables measurement of interfacial width in 3D organic nanostructures. Physical Review Letters 2017, 119, 167801. 62. Limpert, E.; Stahel, W. A.; Abbt, M. Log-normal distributions across the sciences: keys and clues: on the charms of statistics, and how mechanical models resembling gambling machines offer a link to a handy way to characterize log-normal distributions, which can provide deeper insight into variability and probability—normal or log-normal: that is the question. BioScience 2001, 51, 341–352. 63. Ye, L.; Jiao, X. C.; Zhang, S. Q.; Yao, H. F.; Qin, Y. P.; Ade, H.; Hou, J. H. Control of mesoscale morphology and photovoltaic performance in diketopyrrolopyrrole-based small band gap terpolymers. Advanced Energy Materials 2017, 7, 1601138. 64. Mukherjee, S.; Proctor, C. M.; Bazan, G. C.; Nguyen, T. Q.; Ade, H. Significance of average domain purity and mixed domains on the photovoltaic performance of high-efficiency solution-processed small-molecule BHJ solar cells. Advanced Energy Materials 2015, 5, 1500877. 65. Mukherjee, S.; Jiao, X. C.; Ade, H. Charge creation and recombination in multi-length scale polymer:fullerene BHJ solar cell morphologies. Advanced Energy Materials 2016, 6, 1600699. 66. Collins, B. A.; Li, Z.; Tumbleston, J. R.; Gann, E.; McNeill, C. R.; Ade, H. Absolute measurement of domain composition and nanoscale size distribution explains performance in PTB7:PC71BM solar cells. Advanced Energy Materials 2013, 3, 65–74. 67. Ye, L.; Xiong, Y.; Zhang, Q.; Li, S.; Wang, C.; Jiang, Z.; Hou, J.; You, W.; Ade, H. Surpassing 10% efficiency benchmark for nonfullerene organic solar cells by scalable coating in air from single nonhalogenated solvent. Advanced Materials 2018, 30, 1705485. 68. Ma, W.; Tumbleston, J. R.; Ye, L.; Wang, C.; Hou, J.; Ade, H. Quantification of nano- and mesoscale phase separation and relation to donor and acceptor quantum efficiency, Jsc, and FF in polymer:fullerene solar cells. Advanced Materials 2014, 26, 4234–4241. 69. Ma, W.; Ye, L.; Zhang, S.; Hou, J.; Ade, H. Competition between morphological attributes in the thermal annealing and additive processing of polymer solar cells. Journal of Materials Chemistry C 2013, 1, 5023–5030. 70. Cser, F. About the Lorentz correction used in the interpretation of small angle X-ray scattering data of semicrystalline polymers. Journal of Applied Polymer Science 2001, 80, 2300–2308. 71. Chen, W.; Xu, T.; He, F.; Wang, W.; Wang, C.; Strzalka, J.; Liu, Y.; Wen, J.; Miller, D. J.; Chen, J.; Hong, K.; Yu, L.; Darling, S. B. Hierarchical nanomorphologies promote exciton dissociation in polymer/fullerene bulk heterojunction solar cells. Nano Letters 2011, 11, 3707–3713. 72. Harald, A.; Cheng, W.; Hongping, Y. The case for soft X-rays: improved compositional contrast for structure and morphology determination with real and reciprocal space methods. IOP Conference Series: Materials Science and Engineering 2010, 14, 012020. 73. Ade, H.; Hitchcock, A. P. NEXAFS microscopy and resonant scattering: composition and orientation probed in real and reciprocal space. Polymer 2008, 49, 643–675. 74. Stohr, J. NEXAFS spectroscopy, Springer Series in Surface Sciences Vol. 25. Springer 1992. 75. Watts, B.; Swaraj, S.; Nordlund, D.; Lüning, J.; Ade, H. Calibrated NEXAFS spectra of common conjugated polymers. The Journal of Chemical Physics 2011, 134, 024702. 76. Nahid, M. M.; Gann, E.; Thomsen, L.; McNeill, C. R. NEXAFS spectroscopy of conjugated polymers. European Polymer Journal 2016, 81, 532–554.


Conjugated Polymers

77. Collins, B. A.; Cochran, J. E.; Yan, H.; Gann, E.; Hub, C.; Fink, R.; Wang, C.; Schuettfort, T.; McNeill, C. R.; Chabinyc, M. L.; Ade, H. Polarized X-ray scattering reveals non-crystalline orientational ordering in organic films. Nature Materials 2012, 11, 536–543. 78. Gann, E.; Collins, B. A.; Tang, M. L.; Tumbleston, J. R.; Mukherjee, S.; Ade, H. Origins of polarization-dependent anisotropic X-ray scattering from organic thin films. Journal of Synchrotron Radiation 2016, 23, 219–227. 79. Ye, L.; Jiao, X.; Zhao, W.; Zhang, S.; Yao, H.; Li, S.; Ade, H.; Hou, J. Manipulation of domain purity and orientational ordering in high performance all-polymer solar cells. Chemistry of Materials 2016, 28, 6178–6185. 80. Bokel, F. A.; Engmann, S.; Herzing, A. A.; Collins, B. A.; Ro, H. W.; DeLongchamp, D. M.; Richter, L. J.; Schaible, E.; Hexemer, A. In situ X-ray scattering studies of the influence of an additive on the formation of a low-bandgap bulk heterojunction. Chemistry of Materials 2017, 29, 2283–2293. 81. Ye, L.; Zhang, S. Q.; Ma, W.; Fan, B. H.; Guo, X.; Huang, Y.; Ade, H.; Hou, J. H. From binary to ternary solvent: morphology fine-tuning of D/A blends in PDPP3T-based polymer solar cells. Advanced Materials 2012, 24, 6335–6341. 82. Love, J. A.; Collins, S. D.; Nagao, I.; Mukherjee, S.; Ade, H.; Bazan, G. C.; Nguyen T.-Q. Interplay of solvent additive concentration and active layer thickness on the performance of small molecule solar cells. Advanced Materials 2014, 26, 7308–7316. 83. Zhao, J.; Li, Y.; Yang, G.; Jiang, K.; Lin, H.; Ade, H.; Ma, W.; Yan, H. Efficient organic solar cells processed from hydrocarbon solvents. Nature Energy 2016, 1, 15027. 84. Ye, L.; Jiao, X. C.; Zhang, H.; Li, S. S.; Yao, H. F.; Ade, H.; Hou, J. H. 2D-conjugated benzodithiophene-based polymer acceptor: design, synthesis, nanomorphology, and photovoltaic performance. Macromolecules 2015, 48, 7156–7163. 85. Zhong, H. L.; Ye, L.; Chen, J. Y.; Jo, S. B.; Chueh, C. C.; Carpenter, J. H.; Ade, H.; Jen, A. K. Y. A regioregular conjugated polymer for high performance thick-film organic solar cells without processing additive. Journal of Materials Chemistry A 2017, 5, 10517–10525. 86. Liu, F.; Wang, C.; Baral, J. K.; Zhang, L.; Watkins, J. J.; Briseno, A. L.; Russell, T. P. Relating chemical structure to device performance via morphology control in diketopyrrolopyrrole-based low band gap polymers. Journal of the American Chemical Society 2013, 135, 19248–19259. 87. Cai, W.; Liu, P.; Jin, Y.; Xue, Q.; Liu, F.; Russell, T. P.; Huang, , F.; Yip, H.-L.; Cao, Y. Morphology evolution in high-performance polymer solar cells processed from nonhalogenated solvent. Advanced Science 2015, 2, 1500095. 88. Liu, F.; Zhao, W.; Tumbleston, J. R.; Wang, C.; Gu, Y.; Wang, D.; Briseno, A. L.; Ade, H.; Russell, T. P. Understanding the morphology of PTB7:PCBM blends in organic photovoltaics. Advanced Energy Materials 2014, 4, 1301377. 89. Guldal, N. S.; Berlinghof, M.; Kassar, T.; Du, X.; Jiao, X.; Meyer, M.; Ameri, T.; Osvet, A.; Li, N.; Destri, G. L.; Fink, R. H.; Ade, H.; Unruh, T.; Brabec, C. J. Controlling additive behavior to reveal an alternative morphology formation mechanism in polymer:fullerene bulk-heterojunctions. Journal of Materials Chemistry A 2016, 4, 16136–16147. 90. Zhao, W.; Ye, L.; Li, S.; Liu, X.; Zhang, S.; Zhang, Y.; Ghasemi, M.; He, C.; Ade, H.; Hou, J. Environmentally-friendly solvent processed fullerene-free organic solar cells enabled by screening halogen-free solvent additives. Science China Materials 2017, 60, 697–706. 91. Song, X.; Gasparini, N.; Ye, L.; Yao, H.; Hou, J.; Ade, H.; Baran, D. Controlling blend morphology for ultrahigh current density in nonfullerene acceptor-based organic solar cells. ACS Energy Letters 2018, 669–676. 92. Liu, X. Y.; Ye, L.; Zhao, W. C.; Zhang, S. Q.; Li, S. S.; Su, G. M.; Wang, C.; Ade, H.; Hou, J. H. Morphology control enables thickness-insensitive efficient nonfullerene polymer solar cells. Materials Chemistry Frontiers 2017, 1, 2057–2064.

Soft X-Ray Scattering Characterization of Polymer Semiconductors


93. Hu, H.; Chow, P. C. Y.; Zhang, G.; Ma, T.; Liu, J.; Yang, G.; Yan, H. Design of donor polymers with strong temperature-dependent aggregation property for efficient organic photovoltaics. Accounts of Chemical Research 2017, 50, 2549–2528. 94. Min, J.; Jiao, X. C.; Ata, I.; Osvet, A.; Ameri, T.; Bauerle, P.; Ade, H.; Brabec, C. J. Time-dependent morphology evolution of solution-processed small molecule solar cells during solvent vapor annealing. Advanced Energy Materials 2016, 6, 1502579. 95. Roland, S.; Schubert, M.; Collins, B. A.; Kurpiers, J.; Chen, Z.; Facchetti, A.; Ade, H.; Neher, D. Fullerene-free polymer solar cells with highly reduced bimolecular recombination and field-independent charge carrier generation. The Journal of Physical Chemistry Letters 2014, 5, 2815–2822. 96. Min, J.; Jiao, X.; Sgobba, V.; Kan, B.; Heumüller, T.; Rechberger, S.; Spiecker, E.; Guldi, D. M.; Wan, X.; Chen, Y.; Ade, H.; Brabec, C. J. High efficiency and stability small molecule solar cells developed by bulk microstructure fine-tuning. Nano Energy 2016, 28, 241–249. 97. Tumbleston, J. R.; Stuart, A. C.; Gann, E.; You, W.; Ade, H. Fluorinated polymer yields high organic solar cell performance for a wide range of morphologies. Advanced Functional Materials 2013, 23, 3463–3470. 98. Albrecht, S.; Tumbleston, J. R.; Janietz, S.; Dumsch, I.; Allard, S.; Scherf, U.; Ade, H.; Neher, D. Quantifying charge extraction in organic solar cells: the case of fluorinated PCPDTBT. Journal of Physical Chemistry Letters 2014, 5, 1131–1138. 99. Stuart, A. C.; Tumbleston, J. R.; Zhou, H.; Li, W.; Liu, S.; Ade, H.; You, W. Fluorine substituents reduce charge recombination and drive structure and morphology development in polymer solar cells. Journal of the American Chemical Society 2013, 135, 1806–1815. 100. He, X.; Mukherjee, S.; Watkins, S.; Chen, M.; Qin, T.; Thomsen, L.; Ade, H.; McNeill, C. R. Influence of fluorination and molecular weight on the morphology and performance of PTB7:PC71BM solar cells. The Journal of Physical Chemistry C 2014, 118, 9918–9929. 101. Zhao, F.; Dai, S.; Wu, Y.; Zhang, Q.; Wang, J.; Jiang, L.; Ling, Q.; Wei, Z.; Ma, W.; You, W.; Wang, C.; Zhan, X. Single-junction binary-blend nonfullerene polymer solar cells with 12.1% efficiency. Advanced Materials 2017, 29, 1700144. 102. Zhang, Q.; Kelly, M. A.; Bauer, N.; You, W. The curious case of fluorination of conjugated polymers for solar cells. Accounts of Chemical Research 2017, 50, 2401–2409. 103. Bauer, N.; Zhang, Q.; Zhao, J.; Ye, L.; Kim, J.-H.; Constantinou, I.; Yan, L.; So, F.; Ade, H.; Yan, H.; You, W. Comparing non-fullerene acceptors with fullerene in polymer solar cells: a case study with FTAZ and PyCNTAZ. Journal of Materials Chemistry A 2017, 5, 4886–4893. 104. Li, S.; Ye, L.; Zhao, W.; Zhang, S.; Mukherjee, S.; Ade, H.; Hou, J. Energy-level modulation of small-molecule electron acceptors to achieve over 12% efficiency in polymer solar cells. Advanced Materials 2016, 28, 9423–9429. 105. Zhou, Y.; Kurosawa, T.; Ma, W.; Guo, Y.; Fang, L.; Vandewal, K.; Diao, Y.; Wang, C.; Yan, Q.; Reinspach, J.; Mei, J.; Appleton, A. L.; Koleilat, G. I.; Gao, Y.; Mannsfeld, S. C. B.; Salleo, A.; Ade, H.; Zhao, D.; Bao, Z. High performance all-polymer solar cell via polymer side-chain engineering. Advanced Materials 2014, 26, 3767–3772. 106. Li, W. T.; Abrecht, S.; Yang, L. Q.; Roland, S.; Tumbleston, J. R.; McAfee, T.; Yan, L.; Kelly, M. A.; Ade, H.; Neher, D.; You, W. Mobility-controlled performance of thick solar cells based on fluorinated copolymers. Journal of the American Chemical Society 2014, 136, 15566–15576. 107. Li, W.; Yang, L.; Tumbleston, J. R.; Yan, L.; Ade, H.; You, W. Controlling molecular weight of a high efficiency donor-acceptor conjugated polymer and understanding its significant impact on photovoltaic properties. Advanced Materials 2014, 26, 4456–4462. 108. Ma, W.; Yang, G. F.; Jiang, K.; Carpenter, J. H.; Wu, Y.; Meng, X. Y.; McAfee, T.; Zhao, J. B.; Zhu, C. H.; Wang, C.; Ade, H.; Yan, H. Influence of processing parameters and molecular weight on the morphology and properties of high-performance PffBT4T-2OD:PC71BM organic solar cells. Advanced Energy Materials 2015, 5, 1501400.


Conjugated Polymers

109. Kang, H.; Uddin, M. A.; Lee, C.; Kim, K.-H.; Nguyen, T. L.; Lee, W.; Li, Y.; Wang, C.; Woo, H. Y.; Kim, B. J. Determining the role of polymer molecular weight for high-performance all-polymer solar cells: its effect on polymer aggregation and phase separation. Journal of the American Chemical Society 2015, 137, 2359–2365. 110. Kim, J. H.; Gadisa, A.; Schaefer, C.; Yao, H. F.; Gautam, B. R.; Balar, N.; Ghasemi, M.; Constantinou, I.; So, F.; O'Connor, B. T.; Gundogdu, K.; Hou, J. H.; Ade, H. Strong polymer molecular weightdependent material interactions: impact on the formation of the polymer/fullerene bulk heterojunction morphology. Journal of Materials Chemistry A 2017, 5, 13176–13188. 111. Liu, F.; Chen, D.; Wang, C.; Luo, K.; Gu, W.; Briseno, A. L.; Hsu, J. W. P.; Russell, T. P. Molecular weight dependence of the morphology in P3HT:PCBM solar cells. ACS Applied Materials & Interfaces 2014, 6, 19876–19887. 112. Liu, Y.; Zhao, J.; Li, Z.; Mu, C.; Ma, W.; Hu, H.; Jiang, K.; Lin, H.; Ade, H.; Yan, H. Aggregation and morphology control enables multiple cases of high-efficiency polymer solar cells. Nature Communications 2014, 5, 5293. 113. Diao, Y.; Zhou, Y.; Kurosawa, T.; Shaw, L.; Wang, C.; Park, S.; Guo, Y.; Reinspach, J. A.; Gu, K.; Gu, X.; Tee, B. C. K.; Pang, C.; Yan, H.; Zhao, D.; Toney, M. F.; Mannsfeld, S. C. B.; Bao, Z. Flowenhanced solution printing of all-polymer solar cells. Nature Communications 2015, 6, 7955. 114. Ro, H. W.; Downing, J. M.; Engmann, S.; Herzing, A. A.; DeLongchamp, D. M.; Richter, L. J.; Mukherjee, S.; Ade, H.; Abdelsamie, M.; Jagadamma, L. K.; Amassian, A.; Liu, Y.; Yan, H. Morphology changes upon scaling a high-efficiency, solution-processed solar cell. Energy & Environmental Science 2016, 9, 2835–2846. 115. Vakhshouri, K.; Smith, B. H.; Chan, E. P.; Wang, C.; Salleo, A.; Wang, C.; Hexemer, A.; Gomez, E. D. Signatures of intracrystallite and intercrystallite limitations of charge transport in polythiophenes. Macromolecules 2016, 49, 7359–7369. 116. Swaraj, S.; Wang, C.; Yan, H.; Watts, B.; Lüning, J.; McNeill, C. R.; Ade, H. Nanomorphology of bulk heterojunction photovoltaic thin films probed with resonant soft X-ray scattering. Nano Letters 2010, 10, 2863–2869. 117. Kesava, S. V.; Fei, Z.; Rimshaw, A. D.; Wang, C.; Hexemer, A.; Asbury, J. B.; Heeney, M.; Gomez, E. D. Domain compositions and fullerene aggregation govern charge photogeneration in polymer/ fullerene solar cells. Advanced Energy Materials 2014, 4, 1400116. 118. Yan, H.; Collins, B. A.; Gann, E.; Wang, C.; Ade, H.; McNeill, C. R. Correlating the efficiency and nanomorphology of polymer blend solar cells utilizing resonant soft X-ray scattering. ACS Nano 2012, 6, 677–688. 119. Huang, W.; Gann, E.; Chandrasekaran, N.; Thomsen, L.; Prasad, S. K. K.; Hodgkiss, J. M.; Kabra, D.; Cheng, Y. B.; McNeill, C. R. Isolating and quantifying the impact of domain purity on the performance of bulk heterojunction solar cells. Energy & Environmental Science 2017, 10, 1843–1853. 120. He, Z.; Xiao, B.; Liu, F.; Wu, H.; Yang, Y.; Xiao, S.; Wang, C.; Russell, T. P.; Cao, Y. Single-junction polymer solar cells with high efficiency and photovoltage. Nature Photonics 2015, 9, 174–179. 121. Zhang, Q.; Kan, B.; Liu, F.; Long, G.; Wan, X.; Chen, X.; Zuo, Y.; Ni, W.; Zhang, H.; Li, M.; Hu, Z.; Huang, F.; Cao, Y.; Liang, Z.; Zhang, M.; Russell, T. P.; Chen, Y. Small-molecule solar cells with efficiency over 9%. Nat Photonics 2014, 9, 35. 122. Liu, J.; Chen, S. S.; Qian, D. P.; Gautam, B.; Yang, G. F.; Zhao, J. B.; Bergqvist, J.; Zhang, F. L.; Ma, W.; Ade, H.; Inganas, O.; Gundogdu, K.; Gao, F.; Yan, H. Fast charge separation in a non-fullerene organic solar cell with a small driving force. Nature Energy 2016, 1, 16089. 123. Mukherjee, S.; Proctor, C. M.; Tumbleston, J. R.; Bazan, G. C.; Nguyen, T. Q.; Ade, H. Importance of domain purity and molecular packing in efficient solution-processed small-molecule solar cells. Advanced Materials 2015, 27, 1105–1111. 124. Gautam, B. R.; Younts, R.; Li, W.; Yan, L.; Danilov, E.; Klump, E.; Constantinou, I.; So, F.; You, W.; Ade, H.; Gundogdu, K. Charge photogeneration in organic photovoltaics: role of hot versus cold charge-transfer excitons. Advanced Energy Materials 2016, 6, 1301032.

Soft X-Ray Scattering Characterization of Polymer Semiconductors


125. Ma, W.; Tumbleston, J. R.; Wang, M.; Gann, E.; Huang, F.; Ade, H. Domain purity, miscibility, and molecular orientation at donor/acceptor interfaces in high performance organic solar cells: paths to further improvement. Advanced Energy Materials 2013, 3, 864–872. 126. Ma, W.; Reinspach, J.; Zhou, Y.; Diao, Y.; McAfee, T.; Mannsfeld, S. C. B.; Bao, Z. A.; Ade, H. Tuning local molecular orientation-composition correlations in binary organic thin films by solution shearing. Advanced Functional Materials 2015, 25, 3131–3137. 127. Yao, H.; Li, Y.; Hu, H.; Chow, P. C. Y.; Chen, S.; Zhao, J.; Li, Z.; Carpenter, J. H.; Lai, J. Y. L.; Yang, G.; Liu, Y.; Lin, H.; Ade, H.; Yan, H. A facile method to fine-tune polymer aggregation properties and blend morphology of polymer solar cells using donor polymers with randomly distributed alkyl chains. Advanced Energy Materials 2017, 7, 1701895. 128. Li, S.; Ye, L.; Zhao, W.; Liu, X.; Zhu, J.; Ade, H.; Hou, J. Design of a new small-molecule electron acceptor enables efficient polymer solar cells with high fill factor. Advanced Materials 2017, 29, 1704051. 129. Ran, N. A.; Love, J. A.; Heiber, M. C.; Jiao, X.; Hughes, M. P.; Karki, A.; Wang, M.; Brus, V. V.; Wang, H.; Neher, D.; Ade, H.; Bazan, G. C.; Nguyen T.-Q. Charge generation and recombination in an organic solar cell with low energetic offsets. Advanced Energy Materials 2017, 7, 1701073. 130. Gao, K.; Miao, J.; Xiao, L.; Deng, W.; Kan, Y.; Liang, T.; Wang, C.; Huang, F.; Peng, J.; Cao, Y.; Liu, F.; Russell, T. P.; Wu, H.; Peng, X. Multi-length-scale morphologies driven by mixed additives in porphyrin-based organic photovoltaics. Advanced Materials 2016, 28, 4727–4733. 131. Fang, J.; Wang, Z.; Zhang, J.; Zhang, Y.; Deng, D.; Wang, Z.; Lu, K.; Ma, W.; Wei, Z. Understanding the impact of hierarchical nanostructure in ternary organic solar cells. Advanced Science 2015, 2, 1500250. 132. Hou, J.; Inganäs, O.; Friend, R. H.; Gao, F. Organic solar cells based on non-fullerene acceptors. Nature Materials 2018, 17, 119–128. 133. Yao, H. F.; Ye, L.; Hou, J. X.; Jang, B.; Han, G. C.; Cui, Y.; Su, G. M.; Wang, C.; Gao, B. W.; Yu, R. N.; Zhang, H.; Yi, Y. P.; Woo, H. Y.; Ade, H.; Hou, J. H. Achieving highly efficient nonfullerene organic solar cells with improved intermolecular interaction and open-circuit voltage. Advanced Materials 2017, 29, 1700254. 134. Li, S. S.; Ye, L.; Zhao, W. C.; Zhang, S. Q.; Mukherjee, S.; Ade, H.; Hou, J. H. Energy-level modulation of small-molecule electron acceptors to achieve over 12% efficiency in polymer solar cells. Advanced Materials 2016, 28, 9423–9429. 135. Guo, C.; Lin, Y.-H.; Witman, M. D.; Smith, K. A.; Wang, C.; Hexemer, A.; Strzalka, J.; Gomez, E. D.; Verduzco, R. Conjugated block copolymer photovoltaics with near 3% efficiency through microphase separation. Nano Letters 2013, 13, 2957–2963. 136. Lu, L.; Kelly, M. A.; You, W.; Yu, L. Status and prospects for ternary organic photovoltaics. Nature Photonics 2015, 9, 491–500. 137. Yang, L.; Yan, L.; You, W. Organic solar cells beyond one pair of donor–acceptor: ternary blends and more. The Journal of Physical Chemistry Letters 2013, 4, 1802–1810. 138. Zhang, S. Q.; Ye, L.; Hou, J. H. Breaking the 10% efficiency barrier in organic photovoltaics: morphology and device optimization of well-known PBDTTT polymers. Advanced Energy Materials 2016, 6, 1502529. 139. Lu, H.; Zhang, J.; Chen, J.; Liu, Q.; Gong, X.; Feng, S.; Xu, X.; Ma, W.; Bo, Z. Ternary-blend polymer solar cells combining fullerene and nonfullerene acceptors to synergistically boost the photovoltaic performance. Advanced Materials 2016, 28, 9559–9566. 140. Ruderer, M. A.; Wang, C.; Schaible, E.; Hexemer, A.; Xu, T.; Müller-Buschbaum, P. Morphology and optical properties of P3HT:MEH-CN-PPV blend films. Macromolecules 2013, 46, 4491–4501. 141. Schaffer, C. J.; Wang, C.; Hexemer, A.; Müller-Buschbaum, P. Grazing incidence resonant soft X-ray scattering for analysis of multi-component polymer-fullerene blend thin films. Polymer 2016, 105, 357–367.


Conjugated Polymers

142. Zhang, J.; Zhang, Y.; Fang, J.; Lu, K.; Wang, Z.; Ma, W.; Wei, Z. Conjugated polymer–small molecule alloy leads to high efficient ternary organic solar cells. Journal of the American Chemical Society 2015, 137, 8176–8183. 143. Zhang, G.; Zhang, K.; Yin, Q.; Jiang, X.-F.; Wang, Z.; Xin, J.; Ma, W.; Yan, H.; Huang, F.; Cao, Y. High-performance ternary organic solar cell enabled by a thick active layer containing a liquid crystalline small molecule donor. Journal of the American Chemical Society 2017, 139, 2387–2395. 144. Li, H.; Lu, K.; Wei, Z. Polymer/small molecule/fullerene based ternary solar cells. Advanced Energy Materials 2017, 7, 1602540. 145. Wang, Z.; Zhang, Y.; Zhang, J.; Wei, Z.; Ma, W. Optimized “alloy-parallel” morphology of ternary organic solar cells. Advanced Energy Materials 2016, 6, 1502456. 146. Ghasemi, M.; Ye, L.; Zhang, Q.; Yan, L.; Kim, J. H.; Awartani, O.; You, W.; Gadisa, A.; Ade, H. Panchromatic sequentially cast ternary polymer solar cells. Advanced Materials 2017, 29, 1604603. 147. Lu, L.; Xu, T.; Chen, W.; Landry, E. S.; Yu, L. Ternary blend polymer solar cells with enhanced power conversion efficiency. Nature Photonics 2014, 8, 716. 148. Lu, L.; Chen, W.; Xu, T.; Yu, L. High-performance ternary blend polymer solar cells involving both energy transfer and hole relay processes. Nature Communications 2015, 6, 7327. 149. Zhang, Q.; Kelly, M. A.; Hunt, A.; Ade, H.; You, W. Comparative photovoltaic study of physical blending of two donor–acceptor polymers with the chemical blending of the respective moieties. Macromolecules 2016, 49, 2533–2540. 150. Du, X.; Jiao, X.; Rechberger, S.; Perea, J. D.; Meyer, M.; Kazerouni, N.; Spiecker, E.; Ade, H.; Brabec, C. J.; Fink, R. H.; Ameri, T. Crystallization of sensitizers controls morphology and performance in Si-/C-PCPDTBT-sensitized P3HT:ICBA ternary blends. Macromolecules 2017, 50, 2415–2423. 151. Gasparini, N.; Jiao, X. C.; Heumueller, T.; Baran, D.; Matt, G. J.; Fladischer, S.; Spiecker, E.; Ade, H.; Brabec, C. J.; Ameri, T. Designing ternary blend bulk heterojunction solar cells with reduced carrier recombination and a fill factor of 77%. Nature Energy 2016, 1, 16118. 152. Li, Z.; Xu, X.; Zhang, W.; Meng, X.; Genene, Z.; Ma, W.; Mammo, W.; Yartsev, A.; Andersson, M. R.; Janssen, R. A. J.; Wang, E. 9.0% power conversion efficiency from ternary all-polymer solar cells. Energy & Environmental Science 2017, 10, 2212–2221. 153. Jiang, K.; Zhang, G.; Yang, G.; Zhang, J.; Li, Z.; Ma, T.; Hu, H.; Ma, W.; Ade, H.; Yan, H. Multiple cases of efficient nonfullerene ternary organic solar cells enabled by an effective morphology control method. Advanced Energy Materials 2018, 8, 1701370. 154. Fu, H.; Wang, Z.; Sun, Y. Advances in non-fullerene acceptor based ternary organic solar cells. Solar RRL 2017, 1, 1700158. 155. Zhong, L.; Gao, L.; Bin, H.; Hu, Q.; Zhang, Z.-G.; Liu, F.; Russell, T. P.; Zhang, Z.; Li, Y. High efficiency ternary nonfullerene polymer solar cells with two polymer donors and an organic semiconductor acceptor. Advanced Energy Materials 2017, 7, 1602215. 156. Yang, Y.; Chen, W.; Dou, L.; Chang, W.-H.; Duan, H.-S.; Bob, B.; Li, G.; Yang, Y. High-performance multiple-donor bulk heterojunction solar cells. Nature Photonics 2015, 9, 190–198. 157. Güldal, N. S.; Kassar, T.; Berlinghof, M.; Unruh, T.; Brabec, C. J. In situ characterization methods for evaluating microstructure formation and drying kinetics of solution-processed organic bulkheterojunction films. Journal of Materials Research 2017, 32, 1855–1879. 158. Schaffer, C. J.; Palumbiny, C. M.; Niedermeier, M. A.; Jendrzejewski, C.; Santoro, G.; Roth, S. V.; Müller-Buschbaum, P. A direct evidence of morphological degradation on a nanometer scale in polymer solar cells. Advanced Materials 2013, 25, 6760–6764. 159. Liu, F.; Ferdous, S.; Schaible, E.; Hexemer, A.; Church, M.; Ding, X.; Wang, C.; Russell, T. P. Fast printing and in situ morphology observation of organic photovoltaics using slot-die coating. Advanced Materials 2015, 27, 886–891. 160. Chou, K. W.; Yan, B.; Li, R.; Li, E. Q.; Zhao, K.; Anjum, D. H.; Alvarez, S.; Gassaway, R.; Biocca, A.; Thoroddsen, S. T.; Hexemer, A.; Amassian, A. Spin-cast bulk heterojunction solar cells: a dynamical investigation. Advanced Materials 2013, 25, 1923–1929.

Soft X-Ray Scattering Characterization of Polymer Semiconductors


161. Schaffer, C. J.; Palumbiny, C. M.; Niedermeier, M. A.; Burger, C.; Santoro, G.; Roth, S. V.; MüllerBuschbaum, P. Morphological degradation in low bandgap polymer solar cells – an in operando study. Advanced Energy Materials 2016, 6, 1600712. 162. Guo, C.; Kozub, D. R.; Vajjala Kesava, S.; Wang, C.; Hexemer, A.; Gomez, E. D. Signatures of multiphase formation in the active layer of organic solar cells from resonant soft X-ray scattering. ACS Macro Letters 2013, 2, 185–189. 163. Pfeiffer, F. X-ray ptychography. Nature Photonics 2018, 12, 9–17. 164. Shapiro, D. A.; Yu, Y.-S.; Tyliszczak, T.; Cabana, J.; Celestre, R.; Chao, W.; Kaznatcheev, K.; Kilcoyne, A. L. D.; Maia, F.; Marchesini, S.; Meng, Y. S.; Warwick, T.; Yang, L. L.; Padmore, H. A. Chemical composition mapping with nanometre resolution by soft X-ray microscopy. Nature Photonics 2014, 8, 765. 165. David, A. S.; Rich, C.; Peter, D.; Maryam, F.; John, J.; Kilcoyne, A. L. D.; Stefano, M.; Howard, P.; Singanallur, V. V.; Tony, W.; Young-Sang, Y. Ptychographic imaging of nano-materials at the advanced light source with the nanosurveyor instrument. Journal of Physics: Conference Series 2017, 849, 012028. 166. Venkatakrishnan, S. V.; Farmand, M.; Yu, Y. S.; Majidi, H.; Benthem, K. V.; Marchesini, S.; Shapiro, D. A.; Hexemer, A. Robust X-ray phase ptycho-tomography. IEEE Signal Processing Letters 2016, 23, 944–948. 167. Shapiro, D.; Roy, S.; Celestre, R.; Chao, W.; Doering, D.; Howells, M.; Kevan, S.; Kilcoyne, D.; Kirz, J.; Marchesini, S.; Seu, K. A.; Schirotzek, A.; Spence, J.; Tyliszczak, T.; Warwick, T.; Voronov, D.; Padmore, H. A. Development of coherent scattering and diffractive imaging and the COSMIC facility at the advanced light source. Journal of Physics: Conference Series 2013, 425, 192011. 168. Wise, A. M.; Weker, J. N.; Kalirai, S.; Farmand, M.; Shapiro, D. A.; Meirer, F.; Weckhuysen, B. M. Nanoscale chemical imaging of an individual catalyst particle with soft X-ray ptychography. ACS Catalysis 2016, 6, 2178–2181. 169. Wang, C.; Lee, D. H.; Hexemer, A.; Kim, M. I.; Zhao, W.; Hasegawa, H.; Ade, H.; Russell, T. P. Defining the nanostructured morphology of triblock copolymers using resonant soft X-ray scattering. Nano Letters 2011, 11, 3906–3911. 170. Nakatani, Y.; Harada, T.; Takano, A.; Yamada, M.; Watanabe, T. Evaluation of block copolymer structure using soft X-ray scattering. Journal of Photopolymer Science and Technology 2017, 30, 77–82. 171. Brady, M. A.; Ku, S.-Y.; Perez, L. A.; Cochran, J. E.; Schmidt, K.; Weiss, T. M.; Toney, M. F.; Ade, H.; Hexemer, A.; Wang, C.; Hawker, C. J.; Kramer, E. J.; Chabinyc, M. L. Role of solution structure in self-assembly of conjugated block copolymer thin films. Macromolecules 2016, 49, 8187–8197. 172. Culp, T. E.; Ye, D.; Paul, M.; Roy, A.; Behr, M. J.; Jons, S.; Rosenberg, S.; Wang, C.; Gomez, E. W.; Kumar, M.; Gomez, E. D. Probing the internal microstructure of polyamide thin-film composite membranes using resonant soft X-ray scattering. ACS Macro Letters 2018, 927–932. 173. Su, G. M.; Lim, E.; Kramer, E. J.; Chabinyc, M. L. Phase separated morphology of ferroelectric– semiconductor polymer blends probed by synchrotron X-ray methods. Macromolecules 2015, 48, 5861–5867. 174. Zhu, C.; Wang, C.; Young, A.; Liu, F.; Gunkel, I.; Chen, D.; Walba, D.; Maclennan, J.; Clark, N.; Hexemer, A. Probing and controlling liquid crystal helical nanofilaments. Nano Letters 2015, 15, 3420–3424. 175. Salamonczyk, M.; Vaupotic, N.; Pociecha, D.; Wang, C.; Zhu, C.; Gorecka, E. Structure of nanoscale-pitch helical phases: blue phase and twist-bend nematic phase resolved by resonant soft X-ray scattering. Soft Matter 2017, 13, 6694–6699. 176. Zhu, C.; Tuchband, M. R.; Young, A.; Shuai, M.; Scarbrough, A.; Walba, D. M.; Maclennan, J. E.; Wang, C.; Hexemer, A.; Clark, N. A. Resonant carbon K-edge soft X-ray scattering from latticefree heliconical molecular ordering: soft dilative elasticity of the twist-bend liquid crystal phase. Physical Review Letters 2016, 116, 147803.


Conjugated Polymers

177. Abberley, J.; Killah, R.; Walker, R.; Storey, J.; Imrie, C.; Salamonczyk, M.; Zhu, C.; Gorecka, E.; Pociecha, D. Heliconical smectic phases formed by achiral molecules. Nature Communications 2018, 9, 228. 178. Liu, T.; Hu, Q.; Wu, J.; Chen, K.; Zhao, L.; Liu, F.; Wang, C.; Lu, H.; Jia, S.; Russell, T.; Zhu, R.; Gong, Q. Mesoporous PbI2 scaffold for high-performance planar heterojunction perovskite solar cells. Advanced Energy Materials 2016, 6, 1501890. 179. Hansen, S. Calculation of Small-angle scattering profiles using Monte-Carlo simulation. Journal of Applied Crystallography 1990, 23, 344–346.

14 Morphology Evolution and Interfacial Design of Conjugated PolymerBased Photovoltaics Introduction...................................................................................................459 14.1 Polymer:fullerene-Based BHJs....................................................... 460 P3HT:fullerene System • PCPDTBT:fullerene System • DPP Polymer:fullerene System  •  BDT Polymer:fullerene System

14.2 Polymer:non-fullerene Acceptor-Based BHJs.............................. 468 Polymer:PDI Acceptor • Polymer:NDI Acceptor • Polymer:calamitic Shaped Acceptor

Yao Liu and Thomas P. Russell

14.3 Interfacial Design with Polymers....................................................472 14.4 Summary and Outlook.....................................................................475 References.......................................................................................................475

Introduction Polymer solar cells (PSCs) provide an avenue to inexpensive renewable energy by large area coating of lightweight and flexible organic semiconductors.1 –  7  In the past several decades, dramatic improvements in power conversion efficiencies (PCEs) of PSCs have been achieved.8  Tang first reported organic photovoltaics (OPVs) in 1986 with a bilayer device composed of vacuum-evaporated p-type and n-type organic semiconductors, with a PCE of ~1%.9  Later, Sariciftci et al. discovered ultrafast electron transfer from poly[​2-met​hoxy-​5-(2-​ethyl​hexyl​oxy)]​-1,4-​pheny​lenev​inyle​ne (MEH-PPV) to fullerene (C60 ) in 1992, ushering in the promise of conjugated polymers as electron donors and fullerenes as electron acceptors in PSCs.10 , 11  Three years later, Halls et al. and Yu et al. introduced the concept of the bulk heterojunction (BHJ) into PSCs, which could overcome the limitation of exciton (tightly bound electron and hole pair) diffusion length in the photoactive layer.12 , 13  This breakthrough is widely accepted as a milestone in OPVs, setting a standard for efficient OPV device design and optimization. Typically, the fabrication of BHJ active layers is achieved by blending donor and acceptor materials together to form a bi-continuous interpenetrating network with large domains for efficient exciton dissociation, which translates into an improved photo current in devices (Figure 14.1).14 , 15  However, the donor– acceptor blends in active layers that absorb the solar spectrum for electricity generation are complex soft matter structures. The thin film blends are usually cast from a single solvent or solvent mixtures, with or without subsequent post-treatment. Consequently, the resultant thin films are far 459


Conjugated Polymers

FIGURE 14.1  Typical structure of BHJ PSCs.

from equilibrium, which makes the understanding the development of the structure and morphology of the BHJ active layer challenging, since this depends upon multiple kinetic processes (solvent evaporation, phase separation, ordering and crystallization, and interfacial segregation). Due to the overlapping physical functions of the donor and acceptor in the blends, the structure– property relationship must integrate the specific materials, the processing conditions, and function. Three critical device metrics, open circuit voltage (V OC ), short circuit current density (J SC ), and fill factor (FF ) determine PCEs of photovoltaic devices.16  While V OC  is more related to intrinsic properties of active layer materials (e.g. energy band structures) and electrode contacts,17 - 19  J SC  and FF  are largely associated with material processing methods and post-treatments.20  It has been shown that the chemical structure of the materials, molecular weights, donor and acceptor blending ratio, solvent, additive, and annealing method (just to name a few parameters) can significantly influence the morphology and performance of BHJ active layers, thus dictating the resultant photovoltaic properties.2 , 3 , 21 , 22  The morphology of BHJ films includes lateral, vertical, and interfacial features of the films. Tremendous efforts have been devoted to establish a comprehensive understanding of the evolution of the morphology and optimization of the interface.14 , 15 , 23 Structural ordering of the materials, segmental interactions, and miscibility of the donors and acceptors, surface and vertical segregation, and phase separation have all been crucial in dictating the morphology of the BHJ active layers. In OPVs, the photo active layer, typically 100– 200 nm in thickness, is a multi-phased system with morphological parameters differing in the plane of and normal to the surface of the film. Therefore, it is essential to use instrumentation that can discern these differences and, since kinetic processes are critical, methods with sufficient power to observe morphological changes in real-time and in situ . These methods have been reviewed in detail elsewhere.15 The BHJ morphology can be characterized from three different perspectives: a) lateral phase separation (i.e. density correlations in the plane of the film); b) concentration variations normal to the film surface (i.e. vertical phase separation); and c) interfacial morphology and structural order. For each of these, characteristics such as crystallinity, crystal size and orientation, and polymer chain orientation must be considered. In this book chapter, we will review the significant advances that have been made recently in understanding and manipulating the BHJ film morphology and interfacial design of PSCs using several representative material systems.

14.1 Polymer:fullerene-Based BHJs 14.1.1 P3HT:fullerene System Poly (3-hexyltiophene) (P3HT) is a model conjugated polymer in organic electronics, including OPV research, due to its simple chemical structure and richness in morphology.24  The morphological evolution of P3HT:fullerene BHJ films strongly depends on the polymer quality and processing conditions.14  P3HT regioregularity,25  blending ratio,26 –  28  polymer molecular weight,29 , 30  processing solvent,31  annealing conditions,32 –  35  and solvent additives36 –  38  strongly influence the morphology of the BHJ. For example, when a P3HT:fullerene solution mixture is directly cast from chlorobenzene, no obvious phase

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 461

separation is evident, with the exception of a slight crystallization of P3HT.39  Yet a simple thermal treatment, either post- or pre-annealing, enhances the P3HT crystallinity, as evidenced by the increase in the absorption shoulder at ~605 nm,40  which drives a P3HT:fullerene phase separation, increases the length and connectivity of P3HT fibril structures, and induces a bi-continuous morphology.39  Moreover, as shown by small angle neutron scatting (SANS), thermal treatment coarsens the fullerene phase by increasing its agglomerate size and volume percentage.41  The correlation length increased from 0.5 nm to 4.6 nm after 5 s of annealing and reached 5.3 nm after 30 min of annealing. These were further evidenced by high resolution transmission electron microscopy (TEM); within a few seconds of annealing at 150° C, a bi-continuous network morphology was observed in P3HT:fullerene BHJ films with a characteristic length scale of ~10– 15 nm (Figure 14.2A). Grazing incidence X-ray diffraction (GIXD) showed that the P3HT crystal size along (100) and (010) crystal planes in the direction normal to the film surface are ~23 nm and ~12 nm, respectively, consistent with the SANS and TEM measurements (Figure 14.2B). In electron energy loss spectroscopy studies, one of the observed domains in the P3HT:fullerene BHJ films could be assigned to the P3HT crystals. Thus, P3HT nucleation and crystal growth drive the formation of the observed P3HT:fullerene BHJ film morphology.39  Vertical phase separation of P3HT:fullerene BHJ films also influences device performance markedly.42 –  45  Post-annealing, referring to thermal annealing after the cathode is evaporated onto the BHJ film, affords a much better photovoltaic performance than devices fabricated by pre-annealing, where the BHJ film is annealed prior the deposition of cathode. It was found that in the pre-annealing process, P3HT preferentially segregates to the film surface due to its lower surface energy. However, the interfacial energy at the cathode/active layer interface, rather than the surface energy of the materials, is of importance in the post-annealed samples. Therefore, the component concentrations near the cathode/ BHJ film interface are different under the two annealing conditions, though the bulk film morphologies, as shown by neutron and X-ray scattering and TEM, are essentially the same. For the pre-annealed samples, P3HT is more concentrated near the BHJ film surface, while a notable enhancement of fullerene concentration was observed near the cathode/BHJ film interface in the post-annealed samples. This concentration difference of donor and acceptor components normal to the film surface leads to differences in the performance of P3HT:fullerene BHJ films.39 

14.1.2 PCPDTBT:fullerene System Poly[​2,6-(​4,4-b​is(2-​ethyl​hexyl​)-4H-​c yclo​penta​[2,1-​b;3,4​-b′  ]d​ithio​phene​)-alt​-4,7-​(2,1,​3-ben​zothi​adiaz​ ole)]​(PCPDTBT) was first developed by Brabec and coworkers in an attempt to lower the band gap of

FIGURE 14.2  (A) Cross-sectional TEM of post-annealed P3HT:PCBM BHJ samples (30 min heating at 150 °  C). (B) GIXD curves of P3HT:PCBM blend films at different incident angles: as spun, pre-annealed 30 min, and postannealed 30 min. The insets represent the schemes of edge-on and face-on of P3HT chains. Reprinted with permission from Chen et al., Nano Lett.  2011, 11 , 561– 567. Copyright 2011 American Chemical Society.


Conjugated Polymers

conjugated polymers and thus enhance the absorption of the photoactive layer.46  This donor– acceptor (D– A) strategy proved to be an efficient route to extend the absorption of the polymer up to 800 nm with a band gap of 1.4 eV.47  Since then, PCPDTBT has become a representative low band gap polymer in OPV research. Yet, initially, PCPDTBT:fullerene BHJ films only gave a PCE of 3.18%.46  Later, various methods were used to optimize PCPDTBT:fullerene BHJ film morphology and Bazan and coworkers found that using a small amount of non-solvent additives, such as 1,8-octancedithiol (ODT) or 1,8-diiodooctane (DIO), could boost the device efficiency up to 5.5%.48 , 49  This approach was proven to be effective in systems involving many other low band gap polymers and became a standard device preparation technique for low band gap polymer-based OPV fabrication. The solvent additives usually have a boiling point higher than the major solvent. In addition, those solvent additives usually show selective solubility to active layer materials, commonly being bad solvents for the conjugated polymer and good solvents for fullerene. Generally, in the drying process of the BHJ films, the solvent additive deteriorates the solubility of the polymer leading to an ordering of the conjugated polymer, which promotes fibril formation and phase separation while keeping the fullerene dissolved, retarding the formation of large fullerene aggregates (Figure 14.3A).49 

FIGURE 14.3  (A) Schematic depiction of the role of the processing additive in the self-assembly of bulk heterojunction blend materials and structures of PCPDTBT, C71-PCBM, and additives. Reprinted with permission from Lee et al., J. Am. Chem. Soc.  2008, 130 , 3619– 3623. Copyright 2008 American Chemical Society. (B) Morphology for PCPDTBT/PCBM thin films processed with and without additives. Reprinted with permission from Gu et al., Adv. Energy Mater.  2012, 2 , 683– 690. Copyright 2012 John Wiley and Sons.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 463

PCPDTBT shows a low crystallinity when processed without solvent additives.50 –  52  The presence of fullerene further decreases the crystallinity of PCPDTBT, deteriorating the formation of a PCPDTBT percolation network for charge transport. With solvent additives, the crystallinity of PCPDTBT was improved, as evidenced by the appearance of the crystal diffraction peaks in GIXD profiles, as well as the red-shift of the UV– vis absorption. Later, the PCPDTBT polymer chains were found to prefer an edge-on orientation. Moreover, the phase separation between PCPDTBT and fullerene was optimized by use of solvent additives. Resonant soft x-ray scattering (RSoXS) measurements showed a reflection at a scattering vector that matched well with the mesh size of the continuous fibrillar networks observed by TEM. A sharp increase in scattering intensity was found in SANS studies, corresponding to a further phase separation in the inter-fibrillar region. These complementary scattering results indicate a multi-length scale morphology (Figure 14.3B), correlating well with the resultant device performance.50  Several groups also reported that using solvent additives enhanced PCPDTBT chain packing and phase separation with fullerene, leading to improved charge transport by the formation of percolated networks with reduced charge carrier losses.53 –  55  It was also found that different solvent additives had a dramatic influence on PCPDTBT:fullerene BHJ film morphologies and the resultant device performances.49 , 50  Structural modifications of PCPDTBT have also been intensively investigated over the last decades with optimizations mainly focusing on enhancing the π – π  stacking interactions of the polymer chains, which is beneficial for charge transport and photovoltaic properties.56 

14.1.3 DPP polymer:fullerene System Diketopyrrolopyrrole (DPP) is a strong electron-withdrawing unit that can be used to design conjugated polymers with much narrower band gaps when copolymerized with electron donor units.57  DPP polymers have been the focus of a tremendous amount of research in organic electronics (e.g. solar cells, transistors), due to their narrow energy gaps and superior charge transport properties.58  Up to now, single junction OPVs based on DPP polymers have achieved PCEs of > 9%.59  Generally, the aromatic substituents linked with the DPP unit are introduced during the formation of the DPP unit and strongly influence the planarity of the resultant DPP-based monomers.58  Thus, the conformation of the DPP polymer backbone has a significant effect on the π – π  interactions and crystallinity of the resultant polymers. Since 2008, DPP polymers have stimulated significant attention both in terms of comprehensive device physics and morphology characterization.  Wienk et al.60  synthesized a DPP-quaterthiophene polymer (pBBTDPP2 ) with an absorption extending to 900 nm. When processing the pBBTDPP2 :fullerene mixture from a single solvent, a large scale phase separation was found in the BHJ films. The size scale of the phase separation was significantly reduced using a chloroform (CF)/dichlorobenzene (DCB) solvent mixture, leading to relatively high device PCEs. Subsequently, polymers containing DPP and terthiophene or thiophene– benzene– t hiophene (TPT) units were synthesized that pushed the PCEs of DPP polymer-based solar cells to 5.5%. Again, DPP– TPT polymer:fullerene BHJ films coated from a single solvent (i.e. CF) showed large scale (> 200 nm) fullerene clusters and only had a low PCE of 2%. The widely used solvent additive DIO was found to be quite efficient in refi ning the morphology of DPP– TPT polymer:fullerene BHJ films. Only adding a small amount of DIO made the BHJ films more uniform and led to nanoscale fibrillar structures.61 , 62  Huo et al. designed and synthesized a series of DPP polymers with different electron donating units, with a PCE of 4% being achieved with no solvent additive added.63  Janssen and coworkers further combined thienothiophene and DPP and the resultant copolymer offered a high PCE of 5.4%.64  Fré chet and coworkers first incorporated furan units into DPP polymers. This molecular design dramatically improved the solubility of DPP polymers. When the DPP– f uran polymer:fullerene solution was spin coated from chlorobenzene (CB), one could observe macroscopic phase separation in the BHJ films similar to that observed with the DPP– TPT polymer. Interestingly, adding 9% of 1-chloronaphthalene into the BHJ CB solution induced a nanoscale phase separation in the BHJ films, affording a PCE of 5% in solar cell devices.65 


Conjugated Polymers

We designed copolymers containing DPP and fused ring thiophene. Optimizing the processing solvent mixture was found to be crucial for achieving high device PCE.66  Hou and coworkers systematically investigated the morphology of DPP– thiophene polymer:fullerene BHJ films prepared from multiple solvent mixtures. A low PCE of 4.9% was achieved when using a single processing solvent. Yet, a DCB/CF solvent mixture can push the PCE to 5.4%. Further optimization was performed by using a DCB/CF/DIO ternary solvent mixture, yielding a PCE of 6.7%. GIXD characterizations showed that the crystallinity of the BHJ films was continuously improved when the processing solvents were changed from a single solvent to a binary or a ternary solvent mixture. Also, the domain purity increased as evidenced by the RSoXS and total scattering intensity analysis, indicating that domain purity in the BHJ films is also an important factor that influences device performance.67  We also investigated the influence of processing solvent on the morphology of DPP– bithiophene polymer (pDPP):fullerene BHJ films.68 , 69  BHJ films prepared from a single solvent (i.e. CF) showed large scale phase separation due to fullerene aggregation. Adding a small amount of high boiling point solvent (i.e. 5% DCB) was found to yield nanowire mesh networks in the BHJ films with suppressed phase separation, which greatly enhanced the device performance. Systematic optimization showed that a CF/ DCB mixture ratio of 4:1 yielded the best device performance. However, we did not find that the addition of DCB changed the crystallinity of the BHJ films, as measured by GIXD. In contrast, the crystal size became smaller after adding DCB. These results indicated that using a solvent mixture affects DPP crystallization and the distribution of the fullerene in the BHJ films. Recalling the function of the solvent additive, since the solubility of pDPP in DCB is not as good as in CF, DCB can also be regarded as a solvent additive. It was observed that with the evaporation of CF, pDPP initiated crystallization with increasing DCB content. In the BHJ film drying process, more polymer crystals were formed and phase separation began at the same time (Stage I, II, in Figure 14.4). The nanofibrils of pDPP were “ frozen”  into mesh network structures (Stage III in Figure 14.4). The remaining amorphous polymer and fullerene aggregates were located within the mesh of the network (Stage IV in Figure 14.4). DCB is a good solvent for fullerene, leading to a more uniform distribution of fullerene within the BHJ films. Thus, DCB

FIGURE 14.4  Proposed mechanism for morphology evolution (the lines represent polymer chains; the dots represent PC71BM). Reprinted with permission from Liu et al., Adv. Mater.  2012, 24 , 3947– 3951. Copyright 2012 John Wiley and Sons.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 465

mainly works as a solvent additive to promote pDPP crystallization and a uniform fullerene deposition. The BHJ films are not in a thermodynamic equilibrium state and the morphology is dictated by multiple kinetic processes, as evidenced for example with the DIO solvent additive where a slower drying leads to fullerene aggregation. Li et al. found that the diameter of the DPP polymer fibrils is governed by polymer solubility and that the fibril width is crucial for determining the phase separation of the BHJ films.70 , 71  Specifically, the conjugated polymer backbone and the side chains attached to the DPP units significantly influence the solubility of the polymer and, hence, the fibril width. For example, integration of larger size π -conjugated segments can decrease the fibril size to a range that affords a suitable phase separation in the BHJ films (Figure 14.5A). Polymers with shorter side chains on the DPP unit formed smaller width fibrils, allowing more excitons to reach the donor– acceptor interface for charge generation. Moreover, introducing alternating short and long alkyls in DPP-based terpolymers can optimize domain spacing and phase purity in the resultant BHJ films, which further improves solar cell efficiency72 . The branching point of the alkyl chains on the DPP unit also plays an important role on the morphology of the BHJ and device performance. McCulloch and coworkers investigated thienothiophene-substituted DPP polymers with thiophene as a copolymerization unit.73  When moving the alkyl-chain branching position away from the polymer backbone (C1, C2, C3 in Figure 14.5B), the π − π  stacking distance of the polymer chains could be varied which impacted the polymer crystallinity and photovoltaic properties. From C1 to C2

FIGURE 14.5  (A) The influence of π -conjugated segments on the DPP polymer fibril size. Reprinted with permission from Li et al., J. Am. Chem. Soc.  2013, 135 , 18942– 18948. Copyright 2013 American Chemical Society. (B) Structures of DPP polymers with different branching points of the alkyl chains (molecular structures and GIXD patterns). Reprinted with permission from Meager et al., J. Am. Chem. Soc.  2013, 135 , 11537– 11540. Copyright 2013 American Chemical Society.


Conjugated Polymers

and C3, GIXD characterizations showed a small decrease in the π − π  stacking distance (3.59 to 3.52 Å ) and an increase in the degree of crystallinity.

14.1.4 BDT polymer:fullerene System Benzo[1,2-b:4,5-b’ ]dithiophene (BDT) polymers have invigorated the OPV research community due to their superior performance as electron donors. BDT is a planar aromatic unit, which can promote π – π  stacking, making BDT polymers of interest for organic electronics, such as transistors and solar cells.3  Yu and coworkers designed and synthesized a series of BDT– t hieno[3,4-b]thiophene copolymers (PTB1 to PTB7) by varying alkyl side chains and atom decoration of the monomer.74  These BDT polymers share typical quinoid structure with good stability, broad absorption, and deep highest occupied molecular orbital (HOMO) energy levels. In 2010, a world record PCE of PSCs (7.4%) was achieved with a PTB7:fullerene system, offering a bright future for OPVs.75  Interestingly, GIXD showed strong out-of-plane diffraction (010) in the BDT polymer films, indicating an edge-on orientation of the polymer chains (Figure 14.6), which is quite different from many high efficiency conjugated polymers reported previously.76 , 77  The excellent planarity of the BDT polymer backbones may contribute to this unique property. DeLongchamp and co-workers investigated the molecular orientation distribution of PTB7 by polarized light absorption spectroscopy and GIXD,78  showing that PTB7 crystallinity in the blend with

FIGURE 14.6  (A) The chemical structure of PTB7; (B) The GIXD pattern; and (C) the corresponding q x y  and q z  scans of the neat PTB7 thin films on a bare Si substrate. Reprinted with permission from Chen et al., Nano Lett.  2011, 11 , 3707– 3713. Copyright 2011 American Chemical Society.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 467

fullerene is lower than the neat PTB7 and only 20% of the polymer is ordered in the BHJ films. The use of a DIO additive in the casting solution was found to be crucial for making effi cient devices. Yet, the crystallinity and molecular orientation of PTB7 were not signifi cantly changed by solvent additive. However, the films processed with or without DIO showed dramatically different morphologies. Specifically, BHJ films prepared with DIO had a smooth surface and a remarkable decrease of domain size. A complex hierarchical morphology was further highlighted by RSoXS profiles, comprised of domains of PTB7 crystals, crystalline aggregates, and amorphous mixtures of PTB7 and fullerene domains (Figure 14.7), which was thought to contribute to device performance enhancement.77  Ade and coworkers also performed systematic soft X-ray measurements to understand the PTB7:fullerene BHJ film morphology.76  They found that the addition of a small amount of DIO into the blend solution dramatically reduced the domain size of fullerene aggregates in the BHJ films. The miscibility of PTB7 and fullerene is another important factor in dictating the fi nal morphology of the film. Russell and coworkers investigated the structure– property relationships of PTB7:fullerene-based OPVs.79  A multi-lengthscale morphology was observed with the use of a solvent additive. This multi-length-scale structure consists of a phase separated morphology with a characteristic length scale of ~30 nm, which is critical for generating high photocurrents in devices; a second length scale of ~130 nm arises from face-on PTB7 crystalline aggregates. The latter morphology characteristics were also found in films processed without using solvent additives. During the in situ  characterization of the structure formation of the BHJ films, the additive is observed to promote the formation of PTB7 ordered domains at an earlier stage of the solvent evaporation, which is crucial for the development of the final morphology. Moreover, a study of the PTB7:fullerene bilayer morphology evolution shows that the fullerene tends to diffuse into the PTB7 layer. However, the devices prepared by this method gave much lower PCE. This diffusion leads to a swelling of the PTB7 and a decline in the crystallinity, reflecting the miscibility of fullerene with PTB7. Finally, a single-length-scale morphology with slightly large phase separation was generated, leading to poor device performance.

FIGURE 14.7  Diagrammatic hypothesis of the refined BHJ model for PTB7/fullerene solar cells. Reprinted with permission from Chen et al., Nano Lett.  2011, 11 , 3707– 3713. Copyright 2011 American Chemical Society.


Conjugated Polymers

Szarko et al. characterized the detailed structural organization of the BDT polymer-based BHJ films using GIXD.80  They found that the position and shape of the side chain influenced the π – π  stacking distance and the resultant BHJ film morphology and related this to the device performance. Hou and coworkers developed a series of BDT polymers.81 , 82  When replacing the alkoxy groups with alkylthienyl groups, an enhancement in the device performance was found. Subsequently, they investigated the influence of the two different side chain groups on the morphology of the BDT-thienothiophene polymers.83  Introducing an alkylthienyl side chain, which contained a thiophene ring, extended the conjugation length in the side chain direction, increasing the conjugated plane and enhancing π – π  stacking. Thus, a very small π – π  stacking distance of 3.5 Å  was achieved, as confirmed by GIXD. This side chain modification also influences the phase separation of the BHJ films, as shown by grazing incidence small angle X-ray scatting (GISAXS) and TEM. Tiny fi brillar structures were observed in the BHJ films, providing an interpenetrating network that facilitated charge separation and transport, as evidenced by the simultaneous increase in the J SC , FF , and PCE. Fluorine atom decorated BDT-thienothiophene copolymers were also developed to further deepen the HOMO energy level of the materials.84  However, only the mono-fl uorinated polymer afforded the best device performance. Terfl uorination of the polymer backbone led to poor compatibility of the polymer with fullerene, a coarse phase separation, and poor device performance. You and coworkers synthesized BDT-benzothiadiazole-based copolymers and varied the number of fluorine substitutions on the benzothiadiazole units. More fluorine substitutions leads to a more face-on polymer crystallite orientation with respect to the substrate and larger and purer polymer/fullerene domains, which could contribute to the suppressed charge recombination in solar cells.85  They further replaced benzothiadiazole with fluorinated benzotriazole (FTAZ) in the copolymer and found molecular weight of the resultant copolymer significantly influenced the morphology and structural order of the BHJ films.86  When blending FTAZ with a kind of non-fullerene acceptor, this group can achieve efficient solar cells by blade coating in air from a non-halogenated solvent.87 

14.2 Polymer:non-fullerene Acceptor-Based BHJs Non-fullerene-based PSCs are invigorating the OPV community, overcoming some of the drawbacks of fullerene-based systems and offering a more diverse design space for both electron donors and acceptors in the BHJ active layers.88 , 89  Although fullerene acceptors have achieved great success in OPVs, their synthetic inflexibility has led to constraints in manipulating frontier energy levels, as well as poor absorption in the solar spectrum range, and an inherent tendency to undergo post fabrication crystallization, resulting in device instability.90  Non-fullerene acceptors hold the promise to overcome these limitations, providing more tunable absorption with high extinction coefficients, thus contributing to the photocurrent of the devices. Additionally, synthetic polymers afford more tunable solubility in different processing solvents, including non-halogenated solvents, and some other eco-friendly solvents, even water, which is promising for eco-friendly polymer solar cell fabrications.91 - 94 

14.2.1 Polymer:PDI Acceptor Perylene diimide (PDI)-based acceptors have been widely investigated in OPVs due to their desirable design features, including high electron mobility and high electron affinity (EA ; ca. 3.9 eV for the unmodifi ed PDI, similar to widely used fullerene acceptors). In solution processing, the strong intermolecular interactions of PDIs induce large (micron-sized) crystallites and therefore an undesirable coarse phase separation in the BHJ films. Yet these PDI aggregates with strong π − π  stacking are well-suited for electron transport. Therefore, suppressing the strong aggregation of PDIs without sacrifi cing their charge transport ability has been regarded as a critical molecular design strategy for PDI-based acceptors.88 , 95  Shivanna et al. performed imide functionalization on the PDI molecules to interfere with the PDI co-facial stacking. They found that the PDI dimer linked by hydrazine at the imide position had a

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 469

twisted structure, which suppressed the crystallinity of the molecule, yielding an average domain size of ~10 nm when blended with a polymer electron donor. This favorable BHJ film morphology enhanced device J SC  and PCE.96  Lateral position substitution of PDIs has been the focus of much research in nonfullerene acceptor design due to the synthetic accessibility and demonstrated success in suppressing PDI aggregation.97  For example, Wang and coworkers developed a bay-linked PDI dimer with an ~70°  angle between two PDI cores, yielding good device performance when used as acceptor due to its fl exibly twisted structure (Figure 14.8A).98  Further functionalizing the PDI dimer with sulfur bridges in the lateral positions can get a more twisted molecular confi g uration and a slightly lower electron affinity (Figure 14.8B). This molecular design strategy successfully suppressed PDI aggregation and provided a suitable phase separated BHJ film showing a high PCE of 7.16%.99  Linking PDI cores with an aromatic bridge has also been used to produce nonplanar PDI conformations. Yao and coworkers synthesized a thienyl-bridged PDI dimer with a dihedral angle of 50° − 60° 

FIGURE 14.8  Chemical structures of (A) SdiPBI and (B) SdiPBI-S and the corresponding side views of the optimized geometries obtained using density functional theory (DFT) calculations. Reprinted with permission from Sun et al., J. Am. Chem. Soc.  2015, 137 , 11156– 11162. Copyright 2015 American Chemical Society.


Conjugated Polymers

between the two PDI− t hienyl planes, showing dramatic reduction of the aggregation compared with its monomeric counterpart in BHJ films. Thus, the twisted dimer formed small phase separated domains, ∼ 30 nm in size, affording a much higher PCE than its monomeric counterpart.100  Liu et al. developed a PDI acceptor with PDIs linked by a tetraphenylethylene core. The twisted molecule afforded a small domain size of 20 nm when blended with a polymer donor.101  Marks and coworkers promoted a slipstacked crystalline motif to control the molecular stacking of PDIs, which suppressed face-to-face π − π  stacking of PDIs. A PDI acceptor with four phenyl groups substituted in the ortho-position showed a slip-stacked packing structure. These formed small crystalline acceptor grains (∼ 2− 5 nm) when blended with a polymer donor, affording a high device PCE.102  Zhan and coworkers pioneered the use of PDI-based copolymers as acceptors in OPVs.103 , 104  This group used binary additives to optimize the polymer:polymer BHJ morphology and achieved a PCE of 3.45% for all PSCs.105  Bao and coworkers reported a PCE of 4.4% for all PSCs based on an iso -indigo-based polymer donor with polystyrene side chains (PiI-2T-PS5) and a PDI– t hiophene copolymer.106  They further developed a flow-enhanced solution printing method to control phase separation of the all-polymer BHJ active layers.107 

14.2.2 Polymer:NDI Acceptor As another important rylene family member, naphthalene diimide (NDI) was used to develop nonfullerene acceptors. Jenekhe and co-workers developed a NDI– thiophene based small molecular acceptor (NDI– 3TH). They showed that the size and connectivity of the NDI– 3TH domains in the phase-separated P3HT:NDI– 3TH blends vary strongly with the concentration of the processing additive. The device PCE was enhanced ten-fold by using a very small amount of processing additive (0.2 vol%), far below the 2– 3 vol% optimum concentrations found in polymer:fullerene systems.108  Russell and coworkers synthesized a NDI dimer (BiNDI) by linking two NDI monomers with a vinyl donor moiety. Here, a small amount of DIO (i.e. 0.5%) in the blended fi lm facilitated the crystallization of BiNDI into fi brillar crystals, which is benefi cial for the improvement of device performance.109  McNeill and coworkers designed and synthesized NDI-base acceptors with a star-shaped structure: a triarylamine core flanked by three NDI moieties. GIXD measurements indicated the side-chain substitutional atom of the star-shaped molecules dramatically influenced molecular packing. RSoXS measurements showed the polymer:star-shaped acceptor blends do not phase-separate coarsely, with the average domain size for all three star-shaped acceptor blends typically being less than 100 nm.110  NDI-based conjugated polymers were also developed as electron acceptors in OPVs. For example, Mu et al. demonstrated highperformance all-PSCs based on a pair of crystalline low-bandgap polymers, NT:N2200 (Figure 14.9A). The polymer blend film is smooth with an average domain size of about 100 nm. GIXD indicates that NT can maintain its crystallinity with a preferential face-on orientation in the NT:N2200 blends, leading to a hole mobility that is reasonably balanced with the electron mobility of N2200.111  Kim et al. found a pentafluorobenzene-based additive (FPE) that improved the performance of all-PSCs based on PTB7Th:N2200 (Figure 14.9A, B). The polymer blend films prepared with FPE had a bi-continuous interpenetrating network morphology (Figure 14.9C) for efficient charge carrier extraction and transport across donor/acceptor interfaces. Additionally, processing with FPE enhanced the π – π  stacking with a face-on orientation in the bulk state of blend films (Figure 14.9D).112 

14.2.3 Polymer:calamitic Shaped Acceptor Small molecular acceptors with calamitic shape have recently received considerable attention due to the synthetic advantages that draw on the monomer design used in the polymer donor synthesis. Generally, these molecules show a discrete separation of the electron rich and deficient entities, forming a conjugated push− pull structure that narrows the optical bandgap by molecular orbital hybridization and offers control over the separation of the energy levels electron density in the molecule to facilitate charge transfer.88  Chabinyc and coworkers investigated the detailed morphology of BHJs of P3HT

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 471

FIGURE 14.9  (A) Chemical structures of NT and N2200; (B) Chemical structures of PTB7-Th and FPE; (C) Atomic force microscopy (AFM) height image of a PTB7-Th:N2200 blend film processed with CB + FPE; and (D) GIXD pattern of a PTB7-Th:N2200 blend film processed with CB + FPE. Reprinted with permission from Kim et al., Chem. Mater.  2017, 29 , 6793– 6798. Copyright 2017 American Chemical Society.

and a calamitic shaped small-molecule acceptor: 4,7-b​is(4-​(N-he​x ylph​t hali​mide)​v inyl​)benz​o[c]-​1,2,5​thia​diazo​le (HPIBT). GIXD, TEM, and atomic force microscopy (AFM) were used to investigate the kinetics of the phase transformation process of HPIBT during thermal annealing, which changes the BHJ film surface composition and improves charge extraction, leading to FF and PCE improvements. They further showed that HPIBT crystallizes to form micron-sized domains embedded within the film center and a donor rich capping layer near the cathode interface reducing efficient charge extraction.113  Chen and coworkers synthesized a simple small molecular acceptor (DICTF, Figure 14.10A). In the BHJ films, both DICTF and polymer donor exhibit a face-on orientation, which is favorable for effective charge transport in the photovoltaic devices. After thermal treatment, the diffraction peaks of the BHJ


Conjugated Polymers

FIGURE 14.10  Chemical structures of DICTF, F8-DPPTCN, PTB7-Th, PDBT-T1, and ITIC-Th.

films were enhanced, indicating improved molecule packing. DICTF was also miscible with the polymer donor, promoting an interpenetrating network with good morphological stability.114  Chen and coworkers115  developed a non-fullerene acceptor using fluorene as the core with arms of DPP having thiophene-2-carbonitrile as the terminal units (F8-DPPTCN, Figure 14.10). Morphology characterization showed that P3HT and F8-DPPTCN were kinetically trapped in a weakly separated state whereas thermal annealing led to the crystallization of P3HT and the formation of a network structure with a mesh-size of several hundred nanometers. When a solvent additive, DIO, was used and the mixture was thermally annealed, both P3HT and F8-DPPTCN crystallized and a multi-length scale network was formed.115  Zhan and coworkers recently developed an efficient fused-ring electron acceptor (ITIC-Th, Figure 14.10) based on indacenodithieno[3,2-b]thiophene core and thienyl side-chains. In the BHJs, both ITIC-Th and the polymer donor adopt a face-on orientation, beneficial for vertical charge transport. Using a small amount of solvent additive 1-chloronaphthalene (CN) improved the molecular packing in the BHJs, as evidenced by GIXD. RSoXS characterizations by Lin et al. further show that the addition of 1% CN leads to higher relative domain purity and suitable domain sizes, beneficial for exciton dissociation and charge transport in the devices.116  The same group also investigated the influence of electron deficient units on the performance of fused-ring electron acceptors. Interestingly, molecules with two fluorine substituents on the electron deficient units can co-crystallize with the polymer donor, affording BHJs with high crystallinity, as evidenced by the GIXD measurements.117  Hou and coworkers finely tuned the energy levels of small molecular acceptors by optimizing the alkyl substituents on the electron deficient units. They found that both highly ordered molecular packing of the polymer donor and the domain purity are crucial for efficient solar cell devices.118  Their further optimization of small molecule acceptors also showed that domain purity is critical for fullerene-free BHJs to achieve high solar cell efficiencies.119 

14.3 Interfacial Design with Polymers Solution-processible polymers have been recently investigated as efficient interfacial materials in organic electronic and optoelectronic devices.120  Polar and even charged electronically active polymers have been developed up to the point when they are not only capable of replacing the earlier generations of efficient inorganic interlayer materials but also leading to better device performance and stability.121  The morphology and structure of the interlayers plays a critical role in aligning the energy levels at a device interface. For example, conjugated polymer zwitterions (CPZs) have emerged as efficient interlayers in OPVs. From the range of different zwitterionic chemistries, sulfobetaines (SBs) have proven

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 473

most accessible for integration into CPZs.121  Perpendicular alignment of SB side chains to gold (Au) and indium tin oxide (ITO) surfaces was shown by carbon K-edge near-edge X-ray absorption fine structure (NEXAFS) measurements.122  Consistent with the strong dependence of dipole-metal surface interaction energy on distance in the electrostactic model and with routine observations of such dependence in the ultraviolet photoelectron spectroscopy (UPS) measurements (Figure 14.11).123 ,  124  On the other hand, no preferential orientation of SB-groups on the free surface of a zwitterionic interlayer film that was directly deposited onto the active layer, i.e. following the device fabrication sequence prior to metal deposition, was observed.124  This shows that the surface energy minimization due to the hydrophilic nature of side chains and hydrophobic nature of the polymer backbone is not a determining factor. For the orientation of the SB dipole at the interface, DeLongchamp confirmed a net perpendicular orientation of SB side chains to Au and ITO substrates using NEXAFS.125  Richter’ s vibrational resonant sum frequency generation (VR-SFG) measurements acquired on the SO3 −   group in polymer or fullerene zwitterion films on Au and ITO substrates also indicated a net orientation of − C− SO3 −   throughout the entire film with the SO3 −   groups directed towards the Au and ITO surfaces.125  Such dipole orientation was recognized to reduce the work function of metal electrodes, enabling the use of high work function and air stable metals as cathodes in OPVs. Conjugated polymer backbones were also optimized to investigate their performance as device interlayers. Emrick and coworkers developed CPZs, containing thiophene, DPP, and NDI backbones. These CPZs were intergrated into OPVs as interlayers between the photoactive layer and silver (Ag) cathode. The interlayer thickness had only a minor impact on the device performance for the DPP- and NDICPZs, a finding attributed to their electron-transport properties. NDI-based CPZ interlayers provide some of the best performing organic solar cells, and prove useful in conjunction with high-performing polymer-active layers and stable, high-work-function metal cathodes.126  This group further developed the synthesis of zwitterionic poly(phenylene vinylene)s (PPVs) in metal free aqueous polymerizations, using coupling strategies that afford polymers and copolymers having zwitterionic, anionic, and/or

FIGURE 14.11  Chemical structures of the representative CPZs and the alignment of their dipole side chains on metal surface. Reprinted with permission from Liu et al., Adv. Mater.  2013, 25 , 6868– 6873. Copyright 2013 John Wiley and Sons.


Conjugated Polymers

cationic side chains that open applications for PPVs as device interlayers. Interestingly, they found that both the ionic side chains and conjugated backbones exert significant influence on the interfacial energy level alignment and device performace.127  The PPV polymers containing cationic, zwitterionic, or anionic pendent groups were also used to fabricate charge transport layers with specific interfacial ionic functionalities, providing some insightful understanding of the influence of ionic functional groups on ion transport at lead halide perovskite interfaces.128  Cao and coworkers introduced mercury (Hg) into the backbones of amino-functionalized conjugated polymer interlayers. The strong intermolecular Hg– Hg interactions enhance the stacking of the resultant polymer, affording better charge transport.129  They also developed cross-linkable water-/alcohol-soluble conjugated polymer interlayers for high perforance inverted PSCs.130  Interfacial doping aims to modify the energy level alignment at the interface between electrodes and organic semiconductors (i.e. the hard– soft materials interface) to improve charge injection and extraction efficiency. Self-doping interlayer materials were also developed to improve device performance. Wu et al. synthesized NDI-based, self-doped, n-type water/alcohol-soluble conjugated polymers (WSCPs) that can be processed with a broad thickness range (from 5 to 100 nm) as efficient electron transporting layers (ETLs) for high-performance OPVs.131  Hu et al. developed PDI-based n-type polyelectrolytes as OPV device interlayers. They found that the doping behavior, optoelectronic, and morphological properties of the resulting polyelectrolytes can be optimized using different counter ions. The doping effect of counter ions on PDI units afforded the resulting polyelectrolytes with good electron-transporting properties, enabling these materials as thickness-insensitive electron-transporting layers in OPV devices.132  Emrick and coworkers covalently connect planar electronically active PDI-based units with polyelectrolyte dopants (Figure 14.12). They found that varying the PDI-to-cation ratio in these polymers allows chemical control over doping level and that the PDI-based ionene interlayers afford electronic devices with controllable interfacial doping and energy level alignment at the active layer/electrode interface.133  We further demonstrated that some self-doping organic semiconductors (e.g. fullerene derivatives, conjugated polyelectrolytes) hold the promise to stabilize lead halide perovskite solar cells.134 , 135  Wang et al. observed electron transfer in solution in self-doped PDI derivatives. They demonstrate the potential of n-type organic semiconductors with stable n-type doping capability and facile solution processibility for OPV devices.136  Sun et al. found that the contact between the n-type interlayer and the donor provides an extra interface for charge dissociation and the matching of energy levels between the n-type

FIGURE 14.12  Self-doping of PDI-based ionenes enables chemical and morphological control of interfacial doping and charge transport in OPVs. Reprinted with permission from Ref. 133. Copyright 2018 John Wiley and Sons.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 475

interlayer and the acceptor allows efficient electron extraction from the BHJ, which together leads to high solar cell performance.137 

14.4 Summary and Outlook We reviewed recent progress on the active layer morphologies of several typical material systems, the parameters that dictate the evolution of the morphologies, and described general methods to optimize the morphology for efficient solar cells. We highlighted some significant advances that have been made recently in understanding and manipulating interfacial design by using polymer interlayers, including the morphology and structure of interlayers and their influence on device performance. The breakthroughs on designing efficient non-fullerene acceptors in the last three years have boosted the PCEs of single junction PSCs over 14%138 ,  139  and tandem PSCs approaching 15%.140  This impressive PCE progress in such a short time not only makes OPVs more competitive among the third generation photovoltaic technologies, but also suggests much more work lies ahead to fully explore the strength of organic photovoltaic materials. Due to the more diverse material design in this system, it is more challenging to construct a clear picture to understand the structure-property relationship. Some efforts are critical: 1) a systematic investigation of the morphology of polymer:non-fullerene acceptor BHJs, including the effects of miscibility, molecular self-assembly, correlations between morphology and electronic structure, and the processing of the blend films; 2) more powerful techniques that can reveal the relationship between interfacial morphology and electronic structures at the device interface, including molecular orientation, dipole orientation, and their influence on energy level alignment at the interface.


1. Krebs, F. C.; Espinosa, N.; Hö sel, M.; Sø ndergaard, R. R.; Jø rgensen, M. 25th Anniversary Article: Rise to Power— OPV-Based Solar Parks. Adv. Mater.  2014, 26 , 29– 39. 2. Huang, Y.; Kramer, E. J.; Heeger, A. J.; Bazan, G. C. Bulk Heterojunction Solar Cells: Morphology and Performance Relationships. Chem. Rev.  2014, 114 , 7006– 7043. 3. Lu, L.; Zheng, T.; Wu, Q.; Schneider, A. M.; Zhao, D.; Yu, L. Recent Advances in Bulk Heterojunction Polymer Solar Cells. Chem. Rev.  2015, 115 , 12666– 12731. 4. Yan, C.; Barlow, S.; Wang, Z.; Yan, H.; Jen, A. K. Y.; Marder, S. R.; Zhan, X. Non-fullerene Acceptors for Organic Solar Cells. Nat. Rev. Mater.  2018, 3, 18003. 5. Brabec, C. J.; Heeney, M.; McCulloch, I.; Nelson, J. Influence of Blend Microstructure on Bulk Heterojunction Organic Photovoltaic Performance. Chem. Soc. Rev.  2011, 40 , 1185– 1199. 6. Hou, J.; Inganä s, O.; Friend, R. H.; Gao, F. Organic Solar Cells Based on Non-fullerene Acceptors. Nat. Mater.  2018, 17 , 119. 7. Li, G.; Zhu, R.; Yang, Y. Polymer Solar Cells. Nat. Photon.  2012, 6 , 153. 8. https​://ww ​w.nre​​/pv/a​ssets​/pdfs​/PV-e​ffici​encie​s-07-​17-20​18.pd​f 9. Tang, C. W. Two-Layer Organic Photovoltaic Cell. Appl. Phys. Lett.  1986, 48 , 183– 185. 10. Sariciftci, N. S.; Smilowitz, L.; Heeger, A. J.; Wudl, F. Photoinduced Electron Transfer from a Conducting Polymer to Buckminsterfullerene. Science  1992, 258 , 1474– 1476. 11. Kraabel, B.; Lee, C. H.; McBranch, D.; Moses, D.; Sariciftci, N. S.; Heeger, A. J. Ultrafast Photoinduced Electron Transfer in Conducting Polymer— Buckminsterfullerene Composites. Chem. Phys. Lett.  1993, 213 , 389– 394. 12. Halls, J. J. M.; Walsh, C. A.; Greenham, N. C.; Marseglia, E. A.; Friend, R. H.; Moratti, S. C.; Holmes, A. B. Efficient Photodiodes from Interpenetrating Polymer Networks. Nature  1995, 376 , 498. 13. Yu, G.; Gao, J.; Hummelen, J. C.; Wudl, F.; Heeger, A. J. Polymer Photovoltaic Cells: Enhanced Efficiencies via a Network of Internal Donor-Acceptor Heterojunctions. Science  1995, 270 , 1789– 1791.


Conjugated Polymers

14. Liu, F.; Gu, Y.; Jung, J. W.; Jo, W. H.; Russell, T. P. On the Morphology of Polymer-Based Photovoltaics. J. Polym. Sci. B: Polym. Phys.  2012, 50 , 1018– 1044. 15. Liu, F.; Gu, Y.; Shen, X.; Ferdous, S.; Wang, H.-W.; Russell, T. P. Characterization of the Morphology of Solution-Processed Bulk Heterojunction Organic Photovoltaics. Prog. Polym. Sci.  2013, 38 , 1990– 2052. 16. Kippelen, B.; Bredas, J.-L. Organic Photovoltaics. Energy Environ. Sci.  2009, 2 , 251– 261. 17. Yamamoto, S.; Orimo, A.; Ohkita, H.; Benten, H.; Ito, S. Molecular Understanding of the OpenCircuit Voltage of Polymer:Fullerene Solar Cells. Adv. Energy. Mater.  2012, 2, 229– 237. 18. Graham, K. R.; Erwin, P.; Nordlund, D.; Vandewal, K.; Li, R.; Ngongang Ndjawa, G. O.; Hoke, E. T.; Salleo, A.; Thompson, M. E.; McGehee, M. D.; Amassian, A. Re-evaluating the Role of Sterics and Electronic Coupling in Determining the Open-Circuit Voltage of Organic Solar Cells. Adv. Mater.  2013, 25 , 6076– 6082. 19. Vandewal, K.; Tvingstedt, K.; Gadisa, A.; Inganä s, O.; Manca, J. V. On the Origin of the OpenCircuit Voltage of Polymer– Fullerene Solar Cells. Nat. Mater.  2009, 8 , 904– 909. 20. Guo, X. G.; Zhou, N. J.; Lou, S. J.; Smith, J.; Tice, D. B.; Hennek, J. W.; Ortiz, R. P.; Navarrete, J. T. L.; Li, S. Y.; Strzalka, J.; Chen, L. X.; Chang, R. P. H.; Facchetti, A.; Marks, T. J. Polymer Solar Cells with Enhanced Fill Factors. Nat.  Photon.  2013, 7 , 825– 833. 21. Rivnay, J.; Mannsfeld, S. C. B.; Miller, C. E.; Salleo, A.; Toney, M. F. Quantitative Determination of Organic Semiconductor Microstructure from the Molecular to Device Scale. Chem. Rev.  2012, 112 , 5488– 5519. 22. Ruderer, M. A.; Muller-Buschbaum, P. Morphology of Polymer-Based Bulk Heterojunction Films for Organic Photovoltaics. Soft. Matter.  2011, 7 , 5482– 5493. 23. Richter, L. J.; DeLongchamp, D. M.; Amassian, A. Morphology Development in Solution-Processed Functional Organic Blend Films: An In Situ Viewpoint. Chem. Rev.  2017, 117 , 6332– 6366. 24. Dang, M. T.; Hirsch, L.; Wantz, G.; Wuest, J. D. Controlling the Morphology and Performance of Bulk Heterojunctions in Solar Cells. Lessons Learned from the Benchmark Poly(​3-hex​ylthi​ophen​ e):[6​,6]-P​henyl​-C61-​butyr​ic Acid Methyl Ester System. Chem. Rev.  2013, 113 , 3734– 3765. 25. Homyak, P. D.; Liu, Y.; Harris, J. D.; Liu, F.; Carter, K. R.; Russell, T. P.; Coughlin, E. B. Systematic Fluorination of P3HT: Synthesis of P(3HT-co-3H4FT)s by Direct Arylation Polymerization, Characterization, and Device Performance in OPVs. Macromolecules  2016, 49 , 3028– 3037. 26. Mü ller, C.; Ferenczi, T. A. M.; Campoy-Quiles, M.; Frost, J. M.; Bradley, D. D. C.; Smith, P.; Stingelin-Stutzmann, N.; Nelson, J. Binary Organic Photovoltaic Blends: A Simple Rationale for Optimum Compositions. Adv. Mater.  2008, 20 , 3510– 3515. 27. van Bavel, S. S.; Bä renklau, M.; de With, G.; Hoppe, H.; Loos, J. P3HT/PCBM Bulk Heterojunction Solar Cells: Impact of Blend Composition and 3D Morphology on Device Performance. Adv. Funct. Mater.  2010, 20 , 1458– 1463. 28. Sanyal, M.; Schmidt-Hansberg, B.; Klein, M. F. G.; Munuera, C.; Vorobiev, A.; Colsmann, A.; Scharfer, P.; Lemmer, U.; Schabel, W.; Dosch, H.; Barrena, E. Effect of Photovoltaic Polymer/ Fullerene Blend Composition Ratio on Microstructure Evolution during Film Solidification Investigated in Real Time by X-ray Diffraction. Macromolecules  2011, 44 , 3795– 3800. 29. Ma, W.; Kim, J. Y.; Lee, K.; Heeger, A. J. Effect of the Molecular Weight of Poly(3-hexylthiophene) on the Morphology and Performance of Polymer Bulk Heterojunction Solar Cells. Macromol. Rapid Commun.  2007, 28 , 1776– 1780. 30. Ballantyne, A. M.; Chen, L.; Dane, J.; Hammant, T.; Braun, F. M.; Heeney, M.; Duffy, W.; McCulloch, I.; Bradley, D. D. C.; Nelson, J. The Effect of Poly(3-hexylthiophene) Molecular Weight on Charge Transport and the Performance of Polymer:Fullerene Solar Cells. Adv. Funct. Mater.  2008, 18 , 2373– 2380. 31. Ruderer, M. A.; Guo, S.; Meier, R.; Chiang, H.-Y.; Kö rstgens, V.; Wiedersich, J.; Perlich, J.; Roth, S. V.; Mü ller-Buschbaum, P. Solvent-Induced Morphology in Polymer-Based Systems for Organic Photovoltaics. Adv. Funct. Mater.  2011, 21 , 3382– 3391.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 477

32. Li, G.; Yao, Y.; Yang, H.; Shrotriya, V.; Yang, G.; Yang, Y. “ Solvent Annealing”  Effect in Polymer Solar Cells Based on Poly(3-hexylthiophene) and Methanofullerenes. Adv. Funct. Mater.  2007, 17 , 1636– 1644. 33. Chen, F.-C.; Ko, C.-J.; Wu, J.-L.; Chen, W.-C. Morphological Study of P3HT:PCBM Blend Films Prepared Through Solvent Annealing for Solar Cell Applications. Sol. Energ. Mat. Sol. C  2010, 94 , 2426– 2430. 34. Park, J. H.; Kim, J. S.; Lee, J. H.; Lee, W. H.; Cho, K. Effect of Annealing Solvent Solubility on the Performance of Poly(3-hexylthiophene)/Methanofullerene Solar Cells. J. Phys. Chem. C  2009, 113 , 17579– 17584. 35. Li, H.; Tang, H.; Li, L.; Xu, W.; Zhao, X.; Yang, X. Solvent-Soaking Treatment Induced Morphology Evolution in P3HT/PCBM Composite Films. J. Mater. Chem.  2011, 21 , 6563– 6568. 36. Yao, Y.; Hou, J.; Xu, Z.; Li, G.; Yang, Y. Effects of Solvent Mixtures on the Nanoscale Phase Separation in Polymer Solar Cells. Adv. Funct. Mater.  2008, 18, 1783– 1789. 37. Moulé , A. J.; Meerholz, K. Controlling Morphology in Polymer– Fullerene Mixtures. Adv. Mater.  2008, 20 , 240– 245. 38. Chen, H.-Y.; Yang, H.; Yang, G.; Sista, S.; Zadoyan, R.; Li, G.; Yang, Y. Fast-Grown Interpenetrating Network in Poly(3-hexylthiophene): Methanofullerenes Solar Cells Processed with Additive. J. Phys. Chem. C  2009, 113 , 7946– 7953. 39. Chen, D.; Nakahara, A.; Wei, D.; Nordlund, D.; Russell, T. P. P3HT/PCBM Bulk Heterojunction Organic Photovoltaics: Correlating Efficiency and Morphology. Nano Lett.  2011, 11 , 561– 567. 40. Mihailetchi, V. D.; Xie, H. X.; de Boer, B.; Koster, L. J. A.; Blom, P. W. M. Charge Transport and Photocurrent Generation in Poly(3-hexylthiophene): Methanofullerene Bulk-Heterojunction Solar Cells. Adv. Funct. Mater.  2006, 16 , 699– 708. 41. Kiel, J. W.; Eberle, A. P. R.; Mackay, M. E. Nanoparticle Agglomeration in Polymer-Based Solar Cells. Phys. Rev. Lett.  2010, 105 , 168701. 42. Parnell, A. J.; Dunbar, A. D. F.; Pearson, A. J.; Staniec, P. A.; Dennison, A. J. C.; Hamamatsu, H.; Skoda, M. W. A.; Lidzey, D. G.; Jones, R. A. L. Depletion of PCBM at the Cathode Interface in P3HT/PCBM Thin Films as Quantified via Neutron Reflectivity Measurements. Adv. Mater.  2010, 22 , 2444– 2447. 43. Campoy-Quiles, M.; Ferenczi, T.; Agostinelli, T.; Etchegoin, P. G.; Kim, Y.; Anthopoulos, T. D.; Stavrinou, P. N.; Bradley, D. D. C.; Nelson, J. Morphology Evolution via Self-Organization and Lateral and Vertical Diffusion in Polymer:Fullerene Solar Cell Blends. Nat. Mater.  2008, 7 , 158. 44. Xu, Z.; Chen, L.-M.; Yang, G.; Huang, C.-H.; Hou, J.; Wu, Y.; Li, G.; Hsu, C.-S.; Yang, Y. Vertical Phase Separation in Poly(3-hexylthiophene): Fullerene Derivative Blends and Its Advantage for Inverted Structure Solar Cells. Adv. Funct. Mater.  2009, 19 , 1227– 1234. 45. Chen, D.; Liu, F.; Wang, C.; Nakahara, A.; Russell, T. P. Bulk Heterojunction Photovoltaic Active Layers via Bilayer Interdiffusion. Nano Lett.  2011, 11 , 2071– 2078. 46. Mü hlbacher, D.; Scharber, M.; Morana, M.; Zhu, Z.; Waller, D.; Gaudiana, R.; Brabec, C. High Photovoltaic Performance of a Low-Bandgap Polymer. Adv. Mater.  2006, 18 , 2884– 2889. 47. Zhu, Z.; Waller, D.; Gaudiana, R.; Morana, M.; Mü hlbacher, D.; Scharber, M.; Brabec, C. Panchromatic Conjugated Polymers Containing Alternating Donor/Acceptor Units for Photovoltaic Applications. Macromolecules  2007, 40, 1981– 1986. 48. Peet, J.; Kim, J. Y.; Coates, N. E.; Ma, W. L.; Moses, D.; Heeger, A. J.; Bazan, G. C. Efficiency Enhancement in Low-Bandgap Polymer Solar Cells by Processing with Alkane Dithiols. Nat.  Mater.  2007, 6 , 497. 49. Lee, J. K.; Ma, W. L.; Brabec, C. J.; Yuen, J.; Moon, J. S.; Kim, J. Y.; Lee, K.; Bazan, G. C.; Heeger, A. J. Processing Additives for Improved Efficiency from Bulk Heterojunction Solar Cells. J. Am. Chem. Soc.  2008, 130 , 3619– 3623. 50. Gu, Y.; Wang, C.; Russell, T. P. Multi-Length-Scale Morphologies in PCPDTBT/PCBM BulkHeterojunction Solar Cells. Adv. Energy Mater.  2012, 2 , 683– 690.


Conjugated Polymers

51. Agostinelli, T.; Ferenczi, T. A. M.; Pires, E.; Foster, S.; Maurano, A.; Mü ller, C.; Ballantyne, A.; Hampton, M.; Lilliu, S.; Campoy-Quiles, M.; Azimi, H.; Morana, M.; Bradley, D. D. C.; Durrant, J.; Macdonald, J. E.; Stingelin, N.; Nelson, J. The Role of Alkane Dithiols in Controlling Polymer Crystallization in Small Band Gap Polymer:Fullerene Solar Cells. J. Polym. Sci. B: Polym. Phys.  2011, 49 , 717– 724. 52. Rogers, J. T.; Schmidt, K.; Toney, M. F.; Kramer, E. J.; Bazan, G. C. Structural Order in Bulk Heterojunction Films Prepared with Solvent Additives. Adv. Mater.  2011, 23 , 2284– 2288. 53. Di Nuzzo, D.; Aguirre, A.; Shahid, M.; Gevaerts, V. S.; Meskers, S. C. J.; Janssen, R. A. J. Improved Film Morphology Reduces Charge Carrier Recombination into the Triplet Excited State in a Small Bandgap Polymer-Fullerene Photovoltaic Cell. Adv. Mater.  2010, 22 , 4321– 4324. 54. Etzold, F.; Howard, I. A.; Forler, N.; Cho, D. M.; Meister, M.; Mangold, H.; Shu, J.; Hansen, M. R.; Mü llen, K.; Laquai, F. The Effect of Solvent Additives on Morphology and Excited-State Dynamics in PCPDTBT:PCBM Photovoltaic Blends. J. Am. Chem. Soc.  2012, 134 , 10569– 10583. 55. Li, Z.; McNeill, C. R. Transient Photocurrent Measurements of PCDTBT:PC70BM and PCPDTBT:PC70BM Solar Cells: Evidence for Charge Trapping in Efficient Polymer/Fullerene Blends. J. Appl. Phys.  2011, 109 , 074513. 56. Schulz, G. L.; Fischer, F. S. U.; Trefz The PCPDTBT Family: Correlations between Chemical Structure, Polymorphism, and Device Performance. D.; Melnyk, A.; Hamidi-Sakr, A.; Brinkmann, M.; Andrienko, D.; Ludwigs, S. Macromolecules  2017, 50 , 1402– 1414. 57. Grzybowski, M.; Gryko, D. T. Diketopyrrolopyrroles: Synthesis, Reactivity, and Optical Properties. Adv. Optical Mater.  2015, 3, 280– 320. 58. Li, W.; Hendriks, K. H.; Wienk, M. M.; Janssen, R. A. J. Diketopyrrolopyrrole Polymers for Organic Solar Cells. Acc. Chem. Res.  2016, 49 , 78– 85. 59. Choi, H.; Ko, S.-J.; Kim, T.; Morin, P.-O.; Walker, B.; Lee, B. H.; Leclerc, M.; Kim, J. Y.; Heeger, A. J. Small-Bandgap Polymer Solar Cells with Unprecedented Short-Circuit Current Density and High Fill Factor. Adv. Mater.  2015, 27 , 3318– 3324. 60. Wienk, M. M.; Turbiez, M.; Gilot, J.; Janssen, R. A. J. Narrow-Bandgap Diketo-Pyrrolo-Pyrrole Polymer Solar Cells: The Effect of Processing on the Performance. Adv. Mater.  2008, 20 , 2556– 2560. 61. Bijleveld, J. C.; Gevaerts, V. S.; Di Nuzzo, D.; Turbiez, M.; Mathijssen, S. G. J.; de Leeuw, D. M.; Wienk, M. M.; Janssen, R. A. J. Efficient Solar Cells Based on an Easily Accessible Diketopyrrolopyrrole Polymer. Adv.  Mater.  2010, 22 , E242– E246. 62. Bijleveld, J. C.; Zoombelt, A. P.; Mathijssen, S. G. J.; Wienk, M. M.; Turbiez, M.; de Leeuw, D. M.; Janssen, R. A. J. Poly(diketopyrrolopyrrole− terthiophene) for Ambipolar Logic and Photovoltaics. J. Am. Chem.  Soc.  2009, 131 , 16616– 16617. 63. Huo, L.; Hou, J.; Chen, H.-Y.; Zhang, S.; Jiang, Y.; Chen, T. L.; Yang, Y. Bandgap and Molecular Level Control of the Low-Bandgap Polymers Based on 3,6-D​ithio​phen-​2-yl-​2,5-d​ihydr​opyrr​olo[3​ ,4-c]​pyrro​le-1,​4-dio​ne toward Highly Efficient Polymer Solar Cells. Macromolecules  2009, 42 , 6564– 6571. 64. Bronstein, H.; Chen, Z.; Ashraf, R. S.; Zhang, W.; Du, J.; Durrant, J. R.; Shakya Tuladhar, P.; Song, K.; Watkins, S. E.; Geerts, Y.; Wienk, M. M.; Janssen, R. A. J.; Anthopoulos, T.; Sirringhaus, H.; Heeney, M.; McCulloch, I. Thien​o[3,2​-b]th​iophe​ne−  Di​ketop​y rrol​opyrr​ole-C​ontai​ning Polymers for High-Performance Organic Field-Effect Transistors and Organic Photovoltaic Devices. J. Am. Chem.  Soc.  2011, 133 , 3272– 3275. 65. Woo, C. H.; Beaujuge, P. M.; Holcombe, T. W.; Lee, O. P.; Fré chet, J. M. J. Incorporation of Furan into Low Band-Gap Polymers for Efficient Solar Cells. J. Am. Chem.  Soc.  2010, 132 , 15547– 15549. 66. Jung, J. W.; Liu, F.; Russell, T. P.; Jo, W. H. A High Mobility Conjugated Polymer Based on Dithienothiophene and Diketopyrrolopyrrole for Organic Photovoltaics. Energy Environ. Sci.  2012, 5 , 6857– 6861. 67. Ye, L.; Zhang, S.; Ma, W.; Fan, B.; Guo, X.; Huang, Y.; Ade, H.; Hou, J. From Binary to Ternary Solvent: Morphology Fine-Tuning of D/A Blends in PDPP3T-Based Polymer Solar Cells. Adv. Mater.  2012, 24 , 6335– 6341.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 479

68. Liu, F.; Gu, Y.; Wang, C.; Zhao, W.; Chen, D.; Briseno, A. L.; Russell, T. P. Efficient Polymer Solar Cells Based on a Low Bandgap Semi-crystalline DPP Polymer-PCBM Blends. Adv. Mater.  2012, 24 , 3947– 3951. 69. Liu, F.; Wang, C.; Baral, J. K.; Zhang, L.; Watkins, J. J.; Briseno, A. L.; Russell, T. P. Relating Chemical Structure to Device Performance via Morphology Control in DiketopyrrolopyrroleBased Low Band Gap Polymers. J. Am. Chem. Soc.  2013, 135 , 19248– 19259. 70. Li, W.; Hendriks, K. H.; Furlan, A.; Roelofs, W. S. C.; Wienk, M. M.; Janssen, R. A. J.Universal Correlation between Fibril Width and Quantum Efficiency in Diketopyrrolopyrrole-Based Polymer Solar Cells. J. Am. Chem. Soc.  2013, 135 , 18942– 18948. 71. Li, W.; Hendriks, K. H.; Furlan, A.; Roelofs, W. S. C.; Meskers, S. C. J.; Wienk, M. M.; Janssen, R. A. J. Effect of the Fibrillar Microstructure on the Efficiency of High Molecular Weight Diketopyrrolopyrrole-Based Polymer Solar Cells. Adv. Mater.  2014, 26 , 1565– 1570. 72. Ye, L.; Jiao, X.; Zhang, S.; Yao, H.; Qin, Y.; Ade, H.; Hou, J. Control of Mesoscale Morphology and Photovoltaic Performance in Diketopyrrolopyrrole‐ Based Small Band Gap Terpolymers. Adv. Energy Mater.  2017, 7 , 1601138. 73. Meager, I.; Ashraf, R. S.; Mollinger, S.; Schroeder, B. C.; Bronstein, H.; Beatrup, D.; Vezie, M. S.; Kirchartz, T.; Salleo, A.; Nelson, J.; McCulloch, I. Photocurrent Enhancement from Diketopyrrolopyrrole Polymer Solar Cells through Alkyl-Chain Branching Point Manipulation. J. Am. Chem. Soc.  2013, 135 , 11537– 11540. 74. Liang, Y.; Yu, L. A New Class of Semiconducting Polymers for Bulk Heterojunction Solar Cells with Exceptionally High Performance. Acc. Chem. Res.  2010, 43 , 1227– 1236. 75. Liang, Y.; Xu, Z.; Xia, J.; Tsai, S.-T.; Wu, Y.; Li, G.; Ray, C.; Yu, L. For the Bright Future— Bulk Heterojunction Polymer Solar Cells with Power Conversion Efficiency of 7.4%. Adv. Mater.  2010, 22 , E135– E138. 76. Collins, B. A.; Li, Z.; Tumbleston, J. R.; Gann, E.; McNeill, C. R.; Ade, H. Absolute Measurement of Domain Composition and Nanoscale Size Distribution Explains Performance in PTB7:PC71BM Solar Cells. Adv. Energy Mater.  2013, 3 , 65– 74. 77. Chen, W.; Xu, T.; He, F.; Wang, W.; Wang, C.; Strzalka, J.; Liu, Y.; Wen, J.; Miller, D. J.; Chen, J.; Hong, K.; Yu, L.; Darling, S. B. Hierarchical Nanomorphologies Promote Exciton Dissociation in Polymer/Fullerene Bulk Heterojunction Solar Cells. Nano Lett.  2011, 11 , 3707– 3713. 78. Hammond, M. R.; Kline, R. J.; Herzing, A. A.; Richter, L. J.; Germack, D. S.; Ro, H.-W.; Soles, C. L.; Fischer, D. A.; Xu, T.; Yu, L.; Toney, M. F.; DeLongchamp, D. M. Molecular Order in HighEfficiency Polymer/Fullerene Bulk Heterojunction Solar Cells. ACS Nano  2011, 5 , 8248– 8257. 79. Liu, F.; Zhao, W.; Tumbleston, J. R.; Wang, C.; Gu, Y.; Wang, D.; Briseno, A. L.; Ade, H.; Russell, T. P. Understanding the Morphology of PTB7:PCBM Blends in Organic Photovoltaics. Adv. Energy Mater.  2014, 4 , 1301377. 80. Szarko, J. M.; Guo, J.; Liang, Y.; Lee, B.; Rolczynski, B. S.; Strzalka, J.; Xu, T.; Loser, S.; Marks, T. J.; Yu, L.; Chen, L. X. When Function Follows Form: Effects of Donor Copolymer Side Chains on Film Morphology and BHJ Solar Cell Performance. Adv. Mater.  2010, 22 , 5468– 5472. 81. Huo, L.; Zhang, S.; Guo, X.; Xu, F.; Li, Y.; Hou, J. Replacing Alkoxy Groups with Alkylthienyl Groups: A Feasible Approach To Improve the Properties of Photovoltaic Polymers. Angew. Chem. Int. Ed.  2011, 50 , 9697– 9702. 82. Hou, J.; Chen, H.-Y.; Zhang, S.; Chen, R. I.; Yang, Y.; Wu, Y.; Li, G. Synthesis of a Low Band Gap Polymer and Its Application in Highly Efficient Polymer Solar Cells. J. Am. Chem. Soc.  2009, 131 , 15586– 15587. 83. Huang, Y.; Guo, X.; Liu, F.; Huo, L.; Chen, Y.; Russell, T. P.; Han, C. C.; Li, Y.; Hou, J. Improving the Ordering and Photovoltaic Properties by Extending π – Conjugated Area of Electron-Donating Units in Polymers with D-A Structure. Adv. Mater.  2012, 24 , 3383– 3389. 84. Son, H. J.; Wang, W.; Xu, T.; Liang, Y.; Wu, Y.; Li, G.; Yu, L. Synthesis of Fluorinated Polythienothiophene-co-benzodithiophenes and Effect of Fluorination on the Photovoltaic Properties. J. Am. Chem. Soc.  2011, 133, 1885– 1894.


Conjugated Polymers

85. Stuart, A. C.; Tumbleston, J. R.; Zhou, H.; Li, W.; Liu, S.; Ade, H.; You, W. Fluorine Substituents Reduce Charge Recombination and Drive Structure and Morphology Development in Polymer Solar Cells. J. Am. Chem. Soc.  2013, 135, 1806– 1815. 86. Li, W.; Yang, L.; Tumbleston, J. R.; Yan, L.; Ade, H.; You, W. Controlling Molecular Weight of a High Efficiency Donor– Acceptor Conjugated Polymer and Understanding Its Significant Impact on Photovoltaic Properties. Adv. Mater.  2014, 26, 4456– 4462. 87. Ye, L.; Xiong, Y.; Zhang, Q.; Li, S. Wang, C.; Jiang, Z.; Hou, J. You, W.; Ade, H.; Surpassing 10% Efficiency Benchmark for Nonfullerene Organic Solar Cells by Scalable Coating in Air from Single Nonhalogenated Solvent. Adv.  Mater.  2018, 30 , 1705485. 88. Nielsen, C. B.; Holliday, S.; Chen, H.-Y.; Cryer, S. J.; McCulloch, I. Non-Fullerene Electron Acceptors for Use in Organic Solar Cells. Accounts Chem. Res.  2015, 48 , 2803– 2812. 89. Lin, Y.; Zhan, X. Non-fullerene Acceptors for Organic Photovoltaics: An Emerging Horizon. Mater. Horizons  2014, 1 , 470– 488. 90. Cheng, P.; Zhan, X. Stability of Organic Solar Cells: Challenges and Strategies. Chem. Soc. Rev.  2016, 45 , 2544– 2582. 91. Fan, B.; Ying, L.; Wang, Z.; He, B.; Jiang, X.-F.; Huang, F.; Cao, Y. Optimisation of Processing Solvent and Molecular Weight for the Production of Green-Solvent-Processed All-Polymer Solar Cells with a Power Conversion Efficiency over 9%. Energy Environ. Sci.  2017, 10 , 1243– 1251. 92. Ye, L.; Xiong, Y.; Li, S.; Ghasemi, M.; Balar, N.; Turner, J.; Gadisa, A.; Hou, J.; O' Connor, B. T.; Ade, H.; Precise Manipulation of Multilength Scale Morphology and Its Influence on Eco-Friendly Printed All-Polymer Solar Cells. Adv. Funct. Mater.  2017, 27 , 1702016. 93. Zhou, Y.; Gu, K. L.; Gu, X.; Kurosawa, T.; Yan, H.; Guo, Y.; Koleilat, G. I.; Zhao, D.; Toney, M. F.; Bao, Z. All-Polymer Solar Cells Employing Non-Halogenated Solvent and Additive. Chem. Mater.  2016, 28 , 5037– 5042. 94. Schmatz, B.; Yuan, Z.; Lang, A. W.; Hernandez, J. L.; Reichmanis, E.; Reynolds, J. R. Aqueous Processing for Printed Organic Electronics: Conjugated Polymers with Multistage Cleavable Side Chains. ACS Cent. Sci.  2017, 3 , 961– 967. 95. Chen, W.; Zhang, Q. Recent Progress in Non-fullerene Small Molecule Acceptors in Organic Solar Cells (OSCs). J. Mater. Chem. C  2017, 5 , 1275– 1302. 96. Shivanna, R.; Shoaee, S.; Dimitrov, S.; Kandappa, S. K.; Rajaram, S.; Durrant, J. R.; Narayan, K. S. Charge Generation and Transport in Efficient Organic Bulk Heterojunction Solar Cells with a Perylene Acceptor. Energy Environ. Sci.  2014, 7 , 435– 441. 97. Jiang, W.; Li, Y.; Wang, Z. Tailor-Made Rylene Arrays for High Performance n-Channel Semiconductors. Acc. Chem. Res.  2014, 47 , 3135– 3147. 98. Jiang, W.; Ye, L.; Li, X.; Xiao, C.; Tan, F.; Zhao, W.; Hou, J.; Wang, Z. Bay-Linked Perylene Bisimides as Promising Non-fullerene Acceptors for Organic Solar Cells. Chem. Commun.  2014, 50 , 1024– 1026. 99. Sun, D.; Meng, D.; Cai, Y.; Fan, B.; Li, Y.; Jiang, W.; Huo, L.; Sun, Y.; Wang, Z. Non-FullereneAcceptor-Based Bulk-Heterojunction Organic Solar Cells with Efficiency over 7%. J. Am. Chem. Soc.  2015, 137 , 11156– 11162. 100. Zhang, X.; Lu, Z.; Ye, L.; Zhan, C.; Hou, J.; Zhang, S.; Jiang, B.; Zhao, Y.; Huang, J.; Zhang, S.; Liu, Y.; Shi, Q.; Liu, Y.; Yao, J. A Potential Perylene Diimide Dimer-Based Acceptor Material for Highly Efficient Solution-Processed Non-fullerene Organic Solar Cells with 4.03% Efficiency. Adv. Mater.  2013, 25 , 5791– 5797. 101. Liu, Y.; Mu, C.; Jiang, K.; Zhao, J.; Li, Y.; Zhang, L.; Li, Z.; Lai, J. Y. L.; Hu, H.; Ma, T.; Hu, R.; Yu, D.; Huang, X.; Tang, B. Z.; Yan, H. A Tetraphenylethylene Core-Based 3D Structure Small Molecular Acceptor Enabling Efficient Non-fullerene Organic Solar Cells. Adv. Mater.  2015, 27 , 1015– 1020. 102. Hartnett, P. E.; Timalsina, A.; Matte, H. S. S. R.; Zhou, N.; Guo, X.; Zhao, W.; Facchetti, A.; Chang, R. P. H.; Hersam, M. C.; Wasielewski, M. R.; Marks, T. J. Slip-Stacked Perylenediimides as an Alternative Strategy for High Efficiency Nonfullerene Acceptors in Organic Photovoltaics. J. Am. Chem. Soc.  2014, 136 , 16345– 16356.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 481

103. Zhan, X.; Tan, Z. a.; Domercq, B.; An, Z.; Zhang, X.; Barlow, S.; Li, Y.; Zhu, D.; Kippelen, B.; Marder, S. R. A High-Mobility Electron-Transport Polymer with Broad Absorption and Its Use in Field-Effect Transistors and All-Polymer Solar Cells. J. Am. Chem. Soc.  2007, 129 , 7246– 7247. 104. Liu, Y.; Larsen-Olsen, T. T.; Zhao, X.; Andreasen, B.; Sø ndergaard, R. R.; Helgesen, M.; Norrman, K.; Jø rgensen, M.; Krebs, F. C.; Zhan, X. All Polymer Photovoltaics: From Small Inverted Devices to Large Roll-to-Roll Coated and Printed Solar Cells. Sol. Energ. Mat. Sol. C  2013, 112 , 157– 162. 105. Cheng, P.; Ye, L.; Zhao, X.; Hou, J.; Li, Y.; Zhan, X. Binary Additives Synergistically Boost the Efficiency of All-Polymer Solar Cells Up to 3.45%. Energy Environ. Sci.  2014, 7 , 1351– 1356. 106. Zhou, Y.; Kurosawa, T.; Ma, W.; Guo, Y.; Fang, L.; Vandewal, K.; Diao, Y.; Wang, C.; Yan, Q.; Reinspach, J.; Mei, J.; Appleton, A. L.; Koleilat, G. I.; Gao, Y.; Mannsfeld, S. C. B.; Salleo, A.; Ade, H.; Zhao, D.; Bao, Z. High Performance All-Polymer Solar Cell via Polymer Side-Chain Engineering. Adv. Mater.  2014, 26 , 3767– 3772. 107. Diao, Y.; Zhou, Y.; Kurosawa, T.; Shaw, L.; Wang, C.; Park, S.; Guo, Y.; Reinspach, J. A.; Gu, K.; Gu, X.; Tee, B. C. K.; Pang, C.; Yan, H.; Zhao, D.; Toney, M. F.; Mannsfeld, S. C. B.; Bao, Z. FlowEnhanced Solution Printing of All-Polymer Solar Cells. Nat. Commun.  2015, 6 , 7955. 108. Ren, G.; Ahmed, E.; Jenekhe, S. A. Non-Fullerene Acceptor-Based Bulk Heterojunction Polymer Solar Cells: Engineering the Nanomorphology via Processing Additives. Adv. Energy Mater.  2011, 1 , 946– 953. 109. Liu, Y.; Zhang, L.; Lee, H.; Wang, H.-W.; Santala, A.; Liu, F.; Diao, Y.; Briseno, A. L.; Russell, T. P. NDI-Based Small Molecule as Promising Nonfullerene Acceptor for Solution-Processed Organic Photovoltaics. Adv. Energy Mater.  2015, 5 , 1500195. 110. Rundel, K.; Maniam, S.; Deshmukh, K.; Gann, E.; Prasad, S. K. K.; Hodgkiss, J. M.; Langford, S. J.; McNeill, C. R. Naphthalene Diimide-Based Small Molecule Acceptors for Organic Solar Cells. J. Mater. Chem. A  2017, 5 , 12266– 12277. 111. Mu, C.; Liu, P.; Ma, W.; Jiang, K.; Zhao, J.; Zhang, K.; Chen, Z.; Wei, Z.; Yi, Y.; Wang, J.; Yang, S.; Huang, F.; Facchetti, A.; Ade, H.; Yan, H. High-Efficiency All-Polymer Solar Cells Based on a Pair of Crystalline Low-Bandgap Polymers. Adv. Mater.  2014, 26 , 7224– 7230. 112. Kim, H. I.; Kim, M.; Park, C. W.; Kim, H. U.; Lee, H.-K.; Park, T. Morphological Control of Donor/ Acceptor Interfaces in All-Polymer Solar Cells Using a Pentafluorobenzene-Based Additive. Chem. Mater.  2017, 29 , 6793– 6798. 113. O’ Hara, K. A.; Ostrowski, D. P.; Koldemir, U.; Takacs, C. J.; Shaheen, S. E.; Sellinger, A.; Chabinyc, M. L. Role of Crystallization in the Morphology of Polymer:Non-fullerene Acceptor Bulk Heterojunctions. ACS Appl.Mater. Interfaces  2017, 9 , 19021– 19029. 114. Li, M.; Liu, Y.; Ni, W.; Liu, F.; Feng, H.; Zhang, Y.; Liu, T.; Zhang, H.; Wan, X.; Kan, B.; Zhang, Q.; Russell, T. P.; Chen, Y. A Simple Small Molecule as an Acceptor for Fullerene-Free Organic Solar Cells with Efficiency Near 8%. J. Mater. Chem. A  2016, 4 , 10409– 10413. 115. Li, S.; Yan, J.; Li, C.-Z.; Liu, F.; Shi, M.; Chen, H.; Russell, T. P. A Non-fullerene Electron Acceptor Modified by Thiophene-2-Carbonitrile for Solution-Processed Organic Solar Cells. J. Mater. Chem. A  2016, 4 , 3777– 3783. 116. Lin, Y.; Zhao, F.; He, Q.; Huo, L.; Wu, Y.; Parker, T. C.; Ma, W.; Sun, Y.; Wang, C.; Zhu, D.; Heeger, A. J.; Marder, S. R.; Zhan, X. High-Performance Electron Acceptor with Thienyl Side Chains for Organic Photovoltaics. J. Am. Chem. Soc.  2016, 138 , 4955– 4961. 117. Dai, S.; Zhao, F.; Zhang, Q.; Lau, T.-K.; Li, T.; Liu, K.; Ling, Q.; Wang, C.; Lu, X.; You, W.; Zhan, X. Fused Nonacyclic Electron Acceptors for Efficient Polymer Solar Cells. J. Am. Chem. Soc.  2017, 139 , 1336– 1343. 118. Li, S.; Ye, L.; Zhao, W.; Zhang, S.; Mukherjee, S.; Ade, H.; Hou, J. Energy-Level Modulation of Small-Molecule Electron Acceptors to Achieve over 12% Efficiency in Polymer Solar Cells. Adv. Mater.  2016, 28 , 9423– 9429. 119. Yao, H.; Ye, L.; Hou, J.; Jang, B.; Han, G.; Cui, Y.; Su, G. M.; Wang, C.; Gao, B.; Yu, R.; Zhang, H.; Yi, Y.; Woo, H. Y.; Ade, H.; Hou, J. Achieving Highly Efficient Nonfullerene Organic Solar Cells with Improved Intermolecular Interaction and Open-Circuit Voltage. Adv. Mater.  2017, 29 , 1700254.


Conjugated Polymers

120. Hu, Z.; Zhang, K.; Huang, F.; Cao, Y. Water/Alcohol Soluble Conjugated Polymers for the Interface Engineering of Highly Efficient Polymer Light-Emitting Diodes and Polymer Solar Cells. Chem. Commun.  2015, 51 , 5572– 5585. 121. Liu, Y.; Duzhko, V. V.; Page, Z. A.; Emrick, T.; Russell, T. P. Conjugated Polymer Zwitterions: Efficient Interlayer Materials in Organic Electronics. Acc. Chem. Res.  2016, 49 , 2478– 2488. 122. Lee, H.; Puodziukynaite, E.; Zhang, Y.; Stephenson, J. C.; Richter, L. J.; Fischer, D. A.; DeLongchamp, D. M.; Emrick, T.; Briseno, A. L. Poly(sulfobetaine methacrylate)s as Electrode Modifiers for Inverted Organic Electronics. J. Am. Chem. Soc.  2015, 137 , 540– 549. 123. Page, Z. A.; Liu, Y.; Duzhko, V. V.; Russell, T. P.; Emrick, T. Fulleropyrrolidine Interlayers: Tailoring Electrodes to Raise Organic Solar Cell Efficiency. Science  2014, 346 , 441– 444. 124. Liu, F.; Page, Z. A.; Duzhko, V. V.; Russell, T. P.; Emrick, T. Conjugated Polymeric Zwitterions as Efficient Interlayers in Organic Solar Cells. Adv. Mater.  2013, 25 , 6868– 6873. 125. Lee, H.; Stephenson, J. C.; Richter, L. J.; McNeill, C. R.; Gann, E.; Thomsen, L.; Park, S.; Jeong, J.; Yi, Y.; DeLongchamp, D. M.; Page, Z. A.; Puodziukynaite, E.; Emrick, T.; Briseno, A. L. The Structural Origin of Electron Injection Enhancements with Fulleropyrrolidine Interlayers. Adv. Mater. Interfaces  2016, 3 , 1500852. 126. Liu, Y.; Page, Z. A.; Russell, T. P.; Emrick, T. Finely Tuned Polymer Interlayers Enhance Solar Cell Efficiency. Angew. Chem. Int. Ed.  2015, 54 , 11485– 11489. 127. Page, Z. A.; Liu, Y.; Puodziukynaite, E.; Russell, T. P.; Emrick, T. Hydrophilic Conjugated Polymers Prepared by Aqueous Horner– Wadsworth– Emmons Coupling. Macromolecules  2016, 49 , 2526– 2532. 128. Liu, Y.; Renna, L. A.; Thompson, H. B.; Page, Z. A.; Emrick, T.; Barnes, M. D.; Bag, M.; Venkataraman, D.; Russell, T. P. Role of Ionic Functional Groups on Ion Transport at Perovskite Interfaces. Adv. Energy Mater.  2017, 7 , 1701235. 129. Liu, S.; Zhang, K.; Lu, J.; Zhang, J.; Yip, H.-L.; Huang, F.; Cao, Y. High-Efficiency Polymer Solar Cells via the Incorporation of an Amino-Functionalized Conjugated Metallopolymer as a Cathode Interlayer. J. Am. Chem. Soc.  2013, 135 , 15326– 15329. 130. Zhang, K.; Zhong, C.; Liu, S.; Mu, C.; Li, Z.; Yan, H.; Huang, F.; Cao, Y. Highly Efficient Inverted Polymer Solar Cells Based on a Cross-Linkable Water-/Alcohol-Soluble Conjugated Polymer Interlayer. ACS Appl. Mater. Interfaces  2014, 6 , 10429– 10435. 131. Wu, Z.; Sun, C.; Dong, S.; Jiang, X.-F.; Wu, S.; Wu, H.; Yip, H.-L.; Huang, F.; Cao, Y. n-Type Water/Alcohol-Soluble Naphthalene Diimide-Based Conjugated Polymers for High-Performance Polymer Solar Cells. J. Am. Chem. Soc.  2016, 138, 2004– 2013. 132. Hu, Z.; Xu, R.; Dong, S.; Lin, K.; Liu, J.; Huang, F.; Cao, Y. Quaternisation-Polymerized N-type Polyelectrolytes: Synthesis, Characterisation and Application in High-Performance Polymer Solar Cells. Mater. Horizons  2017, 4 , 88– 97. 133. Liu, Y.; Cole, M. D.; Jiang, Y.; Kim, P. Y.; Nordlund, D.; Emrick, T.; Russell, T. P. Chemical and Morphological Control of Interfacial Self-Doping for Efficient Organic Electronics. Adv. Mater.  2018, 30 , 1705976. 134. Liu, Y.; Renna, L. A.; Page, Z. A.; Thompson, H. B.; Kim, P. Y.; Barnes, M. D.; Emrick, T.; Venkataraman, D.; Russell, T. P. A Polymer Hole Extraction Layer for Inverted Perovskite Solar Cells from Aqueous Solutions. Adv. Energy Mater.  2016, 6 , 1600664. 135. Liu, Y.; Page, Z. A.; Zhou, D.; Duzhko, V. V.; Kittilstved, K. R.; Emrick, T.; Russell, T. P. Chemical Stabilization of Perovskite Solar Cells with Functional Fulleropyrrolidines. ACS Cent. Sci.  2018, 4 , 216– 222. 136. Wang, Z.; Zheng, N.; Zhang, W.; Yan, H.; Xie, Z.; Ma, Y.; Huang, F.; Cao, Y. Self-Doped, n-Type Perylene Diimide Derivatives as Electron Transporting Layers for High-Efficiency Polymer Solar Cells. Adv. Energy Mater.  2017, 7 , 1700232.

Morphology Evolution and Interfacial Design of Conjugated Polymer-Based Photovoltaics 483

137. Sun, C.; Wu, Z.; Hu, Z.; Xiao, J.; Zhao, W.; Li, H.-W.; Li, Q.-Y.; Tsang, S.-W.; Xu, Y.-X.; Zhang, K.; Yip, H.-L.; Hou, J.; Huang, F.; Cao, Y. Interface Design for High-Efficiency Non-fullerene Polymer Solar Cells. Energy Environ. Sci.  2017, 10, 1784– 1791. 138. Xiao, Z.; Jia, X.; Ding, L. Ternary Organic Solar Cells Offer 14% Power Conversion Efficiency. Sci. Bull.  2017, 62 , 1562– 1564. 139. Zhang, S.; Qin, Y.; Zhu, J.; Hou, J. Over 14% Efficiency in Polymer Solar Cells Enabled by a Chlorinated Polymer Donor. Adv. Mater.  2018, 30 , 1800868. 140. Che, X.; Li, Y.; Qu, Y.; Forrest, S. R. High Fabrication Yield Organic Tandem Photovoltaics Combining Vacuum- and Solution-Processed Subcells with 15% Efficiency. Nat. Energy  2018, 3 , 422– 427.

15 The Relevance of Solubility and Miscibility for the Performance of Organic Solar Cells 15.1 Introduction.......................................................................................485 15.2 Principles of Mixing..........................................................................487 The Solubility Parameter Concept  •  Flory–Huggins Interaction Parameter • Spinodal Demixing

15.3 Computational Methods..................................................................492 Implicit Solvation Model: Conductor Like Screening Model  •  Prediction of HSP and Physical-Chemical Values

15.4 Solubility and Miscibility: Experimental Methods..................... 496 Solubility • Polymer–Solvent Miscibility • Solute–Solute Miscibility via Melting Point Depression

Stefan Langner, Jose Dario Perea Ospina, Chaohong Zhang, Ning Li, and Christoph J. Brabec

15.5 Miscibility and Phase-Stability in Organic Photovoltaics..........503 Introduction to OPV Stability  •  Microstructure Instabilities in Polymer-Fullerene Composites

15.6 Conclusion..........................................................................................508 15.7 Acknowledgments.............................................................................508 References.......................................................................................................509

15.1 Introduction Tremendous progress has been made in the field of organic photovoltaics (OPV) in the last decade, and excellent power conversion efficiencies (PCEs) of over 13% have been reported for solution-processed lab-scale organic solar cells (OSCs).1–9 The significant improvement in PCE can be attributed to the development of novel organic semiconductors with well-designed chemical structures as well as advanced device optimization methods.10–16 Along with their unique properties, such as low-cost, semitransparency, light-weight and large-scale manufacturing, OSCs are expected to serve as one of the most important sustainable energy sources for non-grid connected applications.17 In general, a module efficiency of > 10% in combination with an operational lifetime of > 10 years and a production cost ≪ 1 €/Wpeak is required to determine the success of a PV technology. The cost potential of roll-to-roll fabricated OSCs based on actual manufacturing data was comprehensively analyzed, and a very competitive production cost as low as 0.05 €/Wp was predicted for OSCs with a PCE of 10%.18 Since the PCE of OSCs 485


Conjugated Polymers

already surpassed the 10% benchmark, more efforts should be focused at the current stage on the processing reliability and on understanding degradation mechanisms of OSCs.19 The high-performance OSCs reported so far are based on the bulk-heterojunction (BHJ) structure, where the organic donor and acceptor are finely mixed in the nanometer regime to facilitate exciton dissociation at the donor/acceptor interface. This regime is frequently referred to as the amorphous regime. Pure donor and acceptor domains are also required to form bi-continuous pathways to sweep out the free charge carriers prior to bulk recombination. These are typically referred to as the crystalline or ordered regimes. The two micro-morphologies, the fine-mixed region and the phase separated region, have to be simultaneously optimized in the delicate BHJ structure to maximize the photovoltaic parameters, such as short circuit current density (JSC) and fill factor (FF). The JSC represents the number of photo-generated charge carriers extracted from an OSC, while the FF is influenced by many factors, and generally can be considered as the competition between extraction and recombination of the photo-generated charge carriers. Various processing strategies, such as thermal or solvent treatment, binary or ternary solvent formulation as well as thermally assisted film drying have been developed for solution-processed OSCs to attain an “ideal” BHJ microstructure.11,16,20 The device lifetime of OSCs is determined by the stability of the favorable but delicate BHJ micromorphology, which is required to be long-term stable under operational conditions. The well-known poly-3hexylthiophene (P3HT) : [6,6]-phenyl-C61-butyric acid methyl ester (PCBM) was the most studied absorber candidate for OSCs over the last decade. The promising reproducibility and the quite guaranteed PCE of 3–4% in combination with the large-quantity availability from commercial providers made P3HT that extremely successful in the first decade of the 21st century. Although thousands of organic semiconductors were developed in recent years to improve the PCE up to 13%, the processing properties and operational stability of state-of-the-art OPV materials are far worse than P3HT:PCBM. This is mainly due to the requirement for significant efforts in meticulous film formation to achieve the delicate BHJ micromorphology. These organic semiconductors in the form of thin films are extremely sensitive to the solvent choice, processing conditions and posttreatments, which lead to different microstructures, subsequently affecting performance. The final microstructure of a BHJ film is determined by the competition between thermodynamics (e.g. solubility, miscibility, etc.) and kinetics (e.g. solvent evaporation, material ordering, phase separation, etc.) during the drying process. The rapid growth of high-performance organic semiconductors is actually challenging the device community in terms of being capable to in depth analyze and characterize novel materials- and process-related performance variations, allowing the gaining of deep and comprehensive understanding. With the development and deployment of various advanced characterization methods, such as energy-filtered transmission electron microscopy (EF-TEM), grazing-incidence wide-angle and small angle X-ray scattering (GIWAXS and GISAXS), resonant soft X-ray scattering (R-SoXS), the microstructure morphology of OSCs was unraveled and well analyzed in recent years.21,22 However, it is worthwhile to note that all these characterization techniques only provide certain morphological information in a very limited area, and couldn’t give a complete picture about the BHJ microstructure morphology on the device level. In order to acquire detailed information on BHJ microstructures, additional characterization methods have to be implemented, such as energy-dispersive x-ray spectroscopy (EDX) for elemental distribution analysis, time-of-flight secondary ion mass spectrometry (ToF-SIMS) for vertical gradient analysis and differential scanning calorimetry (DSC) for thermal and mixing properties. In this book chapter, we will briefly review the stability of OSCs and discuss in more detail the role of solubility, miscibility and interaction parameter on the BHJ microstructure, and how to experimentally measure them. We will further introduce the approach to ab initio calculate the values based on computational methods, and a Figure of Merit (FoM) to predict the microstructure stability of BHJ composites.


The Relevance of Solubility and Miscibility

15.2 Principles of Mixing The mixtures of small non-polar molecular organic compounds are best represented by the change of the Gibbs free energy of mixing ΔGmix. Based on the changes of enthalpy and entropy within the regular solution, the following thermodynamic relationship describes the mixing of two compounds:

∆ Gmix  ∆ H mix  T ∆ Smix  0 (15.1)

where ΔHmix is the enthalpy of mixing and ΔSmix is the entropy of mixing.23 It can be stated that ΔGmix  0 and V EB  VEC/2) can, however, lead to a saturation mode, where ions are continuously injected into the junction. The effect of running IBJTs in saturation mode has not been well studied, but the resulting high ion concentration is believed to cause damage to the junction. A fourth mode, the reverse active mode, can be obtained by reversing both VEC and VEB. Bipolar membrane–based ion diodes and transistors have also been used to create all-ion-conducting logic gates and addressing circuits. Although typically slow compared to their electronic counterparts, the advantage of such circuits lies in the additional functionality that can be obtained, stored, or transmitted through the use of specific ions and biomolecules. Diode logics, e.g., AND and OR functions42 and addressing circuits,43 have been created using bipolar membrane diodes, and more advanced circuits, such as inverters and NAND-gates, have been realized with IBJTs.50


Conjugated Polymers

22.3.4 Ionic Diode Rectifiers to Circumvent Electrode Capacity Limitations In electrophoretic delivery applications, electrodes are used to drive the delivery. To avoid faradaic reactions with potentially toxic by-products, electrodes should be operated in the polarization regime, in which electric double layers are charged and discharged. The limited amount of charge that can be stored in electric double layers therefore constitutes a severe limitation on these kinds of delivery devices. One way around this limitation is to employ ionic diode rectifiers, in which ionic alternating current (AC) is converted into ionic direct current (DC).51 The ionic rectifier circuit is the equivalent of the conventional electronic diode bridge, but it rectifies ionic signals rather than electronic ones, which cannot be accomplished by the conventional electronic version. The design of an ionic rectifier is shown in Figure 22.8a–d. An alternating voltage is applied to the two driving electrodes so that they are operated in the polarization regime. The alternating ionic current is transformed into a direct ionic current by the diodes in the circuit.42, 43 The resulting ionic DC drives the delivery of (in this case) acetylcholine from the source to the target reservoir. The delivered amount of acetylcholine is decoupled from the electrode charging (Figure 22.8e), thus allowing for the delivery of large amounts of substance with relatively small electrodes without extensive faradaic reactions, i.e., without driving the electrodes to fully reduced or oxidized (and therefore irreversible) states.

22.3.5 Fast Delivery Circuits for Neurotransmitter Release Signal transmission in the synapse (interconnects) of biological neurons operates on the order of milliseconds.52 It is therefore desirable that delivery devices for neuroscience applications can match this time domain. This is a grand challenge, as fast delivery typically requires the substance to be stored close to the target, which without active measures results in high passive leakage. Theoretical calculations suggest that ion bipolar diodes can suppress passive diffusional leakage effectively when operated in reverse bias.53 In reverse bias, the diode junction is depleted of mobile ions and a high local electric field prevents leakage through the junction. One solution for a “fast delivery” ion delivery device comprises three terminals – source, waste, and target – where the waste terminal is used to induce a constant horizontal flow to replenish the substance at the point of delivery (Figure 22.9a).54 The delivery can be addressed by assigning

FIGURE 22.8  The four-diode, full-wave ionic current rectifier. (a) Schematic drawing of the diode rectifier showing the driving electrodes (In1, In2), source, and target. (b) The acetylcholine (ACh) molecules are transported from the source reservoir to the target reservoir. (c, d) Depending on the polarity of the applied signal, different ionic diodes are in forward and reverse bias. (e) The ionic AC input signal (bottom) is converted into a DC output signal. The alternating input signal is thereby converted into continuous delivery of ACh (middle). (Figure reproduced with permission from Ref. 51, copyright 2014, John Wiley and Sons.)

Organic Bioelectronics Based on Mixed Ion–Electron Conductors


FIGURE 22.9  Ionic diodes for fast release of neurotransmitters. (a) Schematic drawing of the delivery circuit. The horizontal channel is filled by transporting the substance from the source to the waste. (b) The PEDOT:PSS pad underneath the ion diode is addressed to initiate delivery. The diode is operated in forward bias that allows ions to pass through it. (c) To stop delivery and prevent passive leakage through the diode, it is operated in reverse bias. In this mode, the diode junction is depleted of mobile ions. (d) The diode current changes quickly with the applied voltage of the diode. (e) A measured amount of acetylcholine as a function of pulse length. The delivery starts following a delay of 10–50 ms after the voltage is applied. (Reproduced with permission from Ref. 54.)

individual poly(3,4-ethylenedioxythiophene) polystyrene sulfonate (PEDOT:PSS) electrodes underneath each ionic diode, thereby locally changing the electric potential upon delivery (Figure 22.9b, c, d). Calculations indicate that the time delay until physiological neurotransmitter concentrations are reached at the outlet can be below 10 ms. Indeed, actual devices show an onset of delivery of 10–50 ms (Figure 22.9e). The technology has the potential for creating addressable delivery arrays for spatiotemporally controlled delivery of neurotransmitters, which is attractive for the development of the next generation of brain–machine interfaces.

22.3.6 In Vitro Applications of Iontronics High spatiotemporal control of ion and biomolecule concentrations is attractive for a wide range of in vitro studies of biological systems. OEIPs can provide single-cell stimulation by using 10-µm-wide channel outlets.37 The small spatial extent of the concentration gradient allows for rapid changes in the concentration through changes in the driving voltage. Fast pulses of varying voltage can be used to probe the cellular response both in terms of magnitude and dynamics. As a step toward addressable chemical stimulation – e.g., a “chemical display” akin to matrix-addressed lighting displays – IBJTs have been used for the spatiotemporal control of cell signaling in neuronal cells.46 One application of such chemical displays could be to modulate the chemical microenvironment within tissue slices when studying neuronal signaling; another application could be to mimic spatially distributed neuromuscular junctions55 for future prosthetics. OEIPs can also be used with tissue slices (also known as ex vivo systems). Williamson et al. developed an OEIP system for controlling epileptiform activity in brain slices by ondemand delivery of molecules to only specific regions of the slice, demonstrating the OEIP’s potential


Conjugated Polymers

for therapeutic applications.56 The OEIP can also be used as a tool for controlled amyloid formation, in which the microenvironment can be controlled and thus the conditions for amyloid formation studied.38,57 The electronic control of the OEIP makes it ideal for closed-loop systems, in which sensory information is used to control the delivery of the OEIP. Simon et al. developed such a system, where a freestanding biosensor was used to measure the glutamate concentration, which in turn triggered the delivery of acetylcholine to neuronal cells.58 Jonsson et al. integrated both the sensor and delivery device into a sensor/delivery pixel, which could locally measure epileptic activity and deliver inhibitory neurotransmitters to the same site.59

22.3.7 Delivery of Therapeutic Substances In Vivo OEIP technology is suitable for implantation as it is purely electronically controlled, with no moving mechanical parts that are prone to failure or wear. Simon et al. converted the OEIP into a freestanding device with a sealed reservoir storing the substance to be delivered.39 The device was capable of delivering biomolecules in vivo in an acute setting: modulating hearing sensitivity in guinea pig cochlea. Jonsson et al. further developed the concept into an implantable device with several outlets (Figure 22.10).40 This OEIP was designed to be implanted into the spinal cord of rats, where it was used to treat neuropathic pain by delivering the inhibitory neurotransmitter γ-amino butyric acid (GABA). GABA would not have been possible to administer systemically to treat neuropathic pain due to a range of side effects, demonstrating the advantages of locally controlled administration of drugs. OEIP functionality can also be combined with microfluidic channels, which can supply the substance to be delivered to the vicinity of the delivery location of the OEIP. Uguz et al. developed such a combined device and demonstrated how it could be used to effect the neural activity of the cortex of a rat via GABA delivery.41 Freestanding OEIPs also open up the possibility to explore vastly different applications, as demonstrated by Poxson et al., where OEIPs were used to regulate plant physiology via delivery of the plant hormone auxin.60

FIGURE 22.10  Implantable delivery device. (a) The device comprises a reservoir with an electrode and a patterned delivery channel with four outlets. The channel is designed such that the substance is delivered evenly throughout the outlets. (b) The four outlets are aligned to where the sciatic nerve bundles enter the spinal cord. (c) The pain-relieving effect of GABA is studied by measuring the withdrawal threshold as a function of time for two different rates of GABA delivery and a control. The pain-relieving effect starts after 30 minutes. (d) The withdrawal threshold as a function of the delivered amount of GABA. (Reproduced with permission from Ref. 40.)

Organic Bioelectronics Based on Mixed Ion–Electron Conductors


22.4 Conclusion In this chapter, we have introduced the working principles for two bioelectronic actuator technologies based on organic mixed ion–electron conductors: surface switches and iontronic substance delivery. While these two case studies highlight the potential for such mixed conductors, they just scratch the surface of organic bioelectronics, and the general applications of mixed conductors. In addition to a range of other organic bioelectronic technologies9 including sensors, mechanical actuators, and energy harvesting, mixed-mode organic conductors are seeing application in areas such as thermoelectrics,6 supercapacitors,61,62 and next-generation functional paper.63 The next decade of research is sure to bring interesting intersections of these technologies, applied to biological and healthcare challenges, and beyond.

References 1. JM Leger. Organic electronics: The ions have it. Adv. Mater. 20, 837 (2008). 2. JM Leger, M Berggren, SA Carter. Iontronics: Ionic Carriers in Organic Electronic Materials and Devices (CRC Press: Boca Raton, FL, 2011). http:​//www​.worl​dcat.​org/t​itle/​iontr​onics​-ioni​c-car​riers​ -in-o​rgani​c-ele​ctron​ic-ma​teria​ls-an​d-dev​ices/​oclc/​65384​2654 3. Q Pei, G Yu, C Zhang, Y Yang, AJ Heeger. Polymer light-emitting electrochemical cells. Science 269, 1086 (1995). 4. SB Meier, D Tordera, A Pertegás, C Roldán-Carmona, E Ortí, HJ Bolink. Light-emitting electrochemical cells: Recent progress and future prospects. Mater. Today 17, 217 (2014). 5. H Wang, U Ail, R Gabrielsson, M Berggren, X Crispin. Ionic Seebeck effect in conducting polymers. Adv. Energy Mater. 5, 1500044 (2015). 6. D Zhao, S Fabiano, M Berggren, X Crispin. Ionic thermoelectric gating organic transistors. Nat. Commun. 8, 14214 (2017). 7. P Andersson Ersman, J Kawahara, M Berggren. Printed passive matrix addressed electrochromic displays. Org. Electron. 14, 3371 (2013). 8. P Andersson Ersman, D Nilsson, J Kawahara, G Gustafsson, M Berggren. Fast-switching allprinted organic electrochemical transistors. Org. Electron. 14, 1276 (2013). 9. DT Simon, EO Gabrielsson, K Tybrandt, M Berggren. Organic bioelectronics: Bridging the signaling gap between biology and technology. Chem. Rev. 116, 13009 (2016). 10. M Berggren, A Richter-Dahlfors. Organic bioelectronics. Adv. Mater. 19, 3201 (2007). 11. G Tarabella, FM Mohammadi, N Coppedè, F Barbero, S Iannotta, C Santato, F Cicoira. New opportunities for organic electronics and bioelectronics: Ions in action. Chem. Sci. 4, 1395 (2013). 12. T Someya, Z Bao, GG Malliaras. The rise of plastic bioelectronics. Nature 540, 379 (2016). 13. J Rivnay, RM Owens, GG Malliaras. The rise of organic bioelectronics. Chem. Mater. 26, 679 (2013). 14. J Isaksson, C Tengstedt, M Fahlman, N Robinson, M Berggren. A solid-state organic electronic wettability switch. Adv. Mater. 16, 316 (2004). 15. G Kossmehl, M Niemitz. Preparation and controlled wettability of poly (2,2′-bithienyl-5,5′-diyl) layers. Synth. Met. 41, 1065 (1991). 16. KM Yamada, K Olden. Fibronectins-adhesive glycoproteins of cell surface and blood. Nature 275, 178 (1978). 17. RO Hynes. Integrins: A family of cell surface receptors. Cell 48, 549 (1987). 18. JY Wong, R Langer, DE Ingber. Electrically conducting polymers can noninvasively control the shape and growth of mammalian cells. Proc. Natl. Acad. Sci. 91, 3201 (1994). 19. K Svennersten, MH Bolin, EWH Jager, M Berggren, A Richter-Dahlfors. Electrochemical modulation of epithelia formation using conducting polymers. Biomaterials 30, 6257 (2009).


Conjugated Polymers

20. AMD Wan, RM Schur, CK Ober, C Fischbach, D Gourdon, GG Malliaras. Electrical control of protein conformation. Adv. Mater. 24, 2501 (2012). 21. A Herland, KM Persson, V Lundin, M Fahlman, M Berggren, EWH Jager, AI Teixeira. Electrochemical control of growth factor presentation to steer neural stem cell differentiation. Angew. Chem. 50, 12529 (2011). 22. L Faxälv, MH Bolin, EWH Jager, TL Lindahl, M Berggren. Electronic control of platelet adhesion using conducting polymer microarrays. Lab. Chip 14, 3043 (2014). 23. J Rivnay, S Inal, A Salleo, RM Owens, M Berggren, GG Malliaras. Organic electrochemical transistors. Nat. Rev. Mater. 3, 17086 (2018). 24. MH Bolin, K Svennersten, D Nilsson, A Sawatdee, EWH Jager, A Richter-Dahlfors, M Berggren. Active control of epithelial cell-density gradients grown along the channel of an organic electrochemical transistor. Adv. Mater. 21, 4379 (2009). 25. B Winther-Jensen, K West. Vapor-phase polymerization of 3,4-ethylenedioxythiophene: A route to highly conducting polymer surface layers. Macromolecules 37, 4538 (2004). 26. MH Bolin, X Wang, IS Chronakis, EWH Jager, M Berggren. Nano-fiber scaffold electrodes based on PEDOT for cell stimulation. Sensors Actuators B Chem. 142, 451 (2009). 27. EWH Jager, E Smela, O Inganäs. Microfabricating conjugated polymer actuators. Science 290, 1540 (2000). 28. X Wang, M Berggren, O Inganäs. Dynamic control of surface energy and topography of microstructured conducting polymer films. Langmuir 24, 5942 (2008). 29. K Svennersten, M Berggren, A Richter-Dahlfors, EWH Jager. Mechanical stimulation of epithelial cells using polypyrrole microactuators. Lab. Chip 11, 3287 (2011). 30. G Zotti, S Zecchin, G Schiavon, L ‘Bert’ Groenendaal. Electrochemical and chemical synthesis and characterization of sulfonated poly (3,4-ethylenedioxythiophene): A novel water-soluble and highly conductive conjugated oligomer. Macromol. Chem. Phys. 203, 1958 (2002). 31. RH Karlsson, A Herland, M Hamedi, JA Wigenius, A Åslund, X Liu, M Fahlman, O Inganäs, P Konradsson. Iron-catalyzed polymerization of alkoxysulfonate-functionalized 3,4-ethylenedioxythiophene gives water-soluble poly(3,4-ethylenedioxythiophene) of high conductivity. Chem. Mater. 21, 1815 (2009). 32. KM Persson, R Karlsson, K Svennersten, S Löffler, EWH Jager, A Richter-Dahlfors, P Konradsson, M Berggren. Electronic control of cell detachment using a self-doped conducting polymer. Adv. Mater. 23, 4403 (2011). 33. KM Persson, R Gabrielsson, A Sawatdee, D Nilsson, P Konradsson, M Berggren. Electronic control over detachment of a self-doped water-soluble conjugated polyelectrolyte. Langmuir 30, 6257 (2014). 34. KM Persson, S Lönnqvist, K Tybrandt, R Gabrielsson, D Nilsson, G Kratz, M Berggren. Matrix addressing of an electronic surface switch based on a conjugated polyelectrolyte for cell sorting. Adv. Funct. Mater. 25, 7056 (2015). 35. T Arbring Sjöström, M Berggren, EO Gabrielsson, P Janson, DJ Poxson, M Seitanidou, DT Simon. A decade of iontronic delivery devices. Adv. Mater. Technol. 3, 1700360 (2018). 36. J Isaksson, P Kjäll, D Nilsson, N Robinson, M Berggren, A Richter-Dahlfors. Electronic control of Ca2+ signalling in neuronal cells using an organic electronic ion pump. Nat. Mater. 6, 673 (2007). 37. K Tybrandt, KC Larsson, S Kurup, DT Simon, P Kjäll, J Isaksson, M Sandberg, EWH Jager, A Richter-Dahlfors, M Berggren. Translating electronic currents to precise acetylcholine-induced neuronal signaling using an organic electrophoretic delivery device. Adv. Mater. 21, 4442 (2009). 38. EO Gabrielsson, K Tybrandt, P Hammarström, M Berggren, KPR Nilsson. Spatially controlled amyloid reactions using organic electronics. Small 6, 2153 (2010). 39. M Seitanidou, JF Franco-Gonzalez, TA Sjöström, I Zozoulenko, M Berggren, DT Simon. pH Dependence of γ-Aminobutyric Acid Iontronic Transport. J. Phys. Chem. B. 121, 7284 (2017) doi:10.1021/acs.jpcb.7b05218

Organic Bioelectronics Based on Mixed Ion–Electron Conductors


40. A Jonsson, Z Song, D Nilsson, BA Meyerson, DT Simon, B Linderoth, M Berggren. Therapy using implanted organic bioelectronics. Sci. Adv. 1, e1500039 (2015). 41. I Uguz, CM Proctor, VF Curto, AM Pappa, MJ Donahue, M Ferro, RM Owens, D Khodagholy, S Inal, GG Malliaras. A microfluidic ion pump for in vivo drug delivery. Adv. Mater. 29, 1701217 (2017). 42. J-H Han, KB Kim, HC Kim, TD Chung. Ionic circuits based on polyelectrolyte diodes on a microchip. Angew. Chem. 48, 3830 (2009). 43. EO Gabrielsson, K Tybrandt, M Berggren. Ion diode logics for pH control. Lab. Chip 12, 2507 (2012). 44. J-H Han, KB Kim, JH Bae, BJ Kim, CM Kang, HC Kim, TD Chung. Ion flow crossing over a polyelectrolyte diode on a microfluidic chip. Small 7, 2629 (2011). 45. L-J Cheng, H-C Chang. Microscale pH regulation by splitting water. Biomicrofluidics 5, 046502 (2011). 46. K Tybrandt, KC Larsson, A Richter-Dahlfors, M Berggren. Ion bipolar junction transistors. Proc. Natl. Acad. Sci. 107, 9929 (2010). 47. K Tybrandt, EO Gabrielsson, M Berggren. Toward complementary ionic circuits: The npn ion bipolar junction transistor. J. Am. Chem. Soc. 133, 10141 (2011). 48. EO Gabrielsson, K Tybrandt, M Berggren. Polyphosphonium-based ion bipolar junction transistors. Biomicrofluidics 8, 64116 (2014). 49. AV Volkov, K Tybrandt, M Berggren, IV Zozoulenko. Modeling of charge transport in ion bipolar junction transistors. Langmuir 30, 6999 (2014). 50. K Tybrandt, R Forchheimer, M Berggren. Logic gates based on ion transistors. Nat. Commun. 3, 871 (2012). 51. EO Gabrielsson, P Janson, K Tybrandt, DT Simon, M Berggren. A four-diode full-wave ionic current rectifier based on bipolar membranes: Overcoming the limit of electrode capacity. Adv. Mater. 26, 5143 (2014). 52. J-W Lin, DS Faber. Modulation of synaptic delay during synaptic plasticity. Trends Neurosci. 25, 449 (2002). 53. K Tybrandt. Exploring the potential of ionic bipolar diodes for chemical neural interfaces. Soft Matter 13, 8171 (2017). 54. A Jonsson, T Arbring Sjöström, K Tybrandt, M Berggren, DT Simon. Chemical delivery array with millisecond neurotransmitter release. Sci. Adv. 2, e1601340 (2016). 55. ZW Hall, JR Sanes. Synaptic structure and development: The neuromuscular junction. Cell 72, 99 (1993). 56. A Williamson, J Rivnay, L Kergoat, A Jonsson, S Inal, I Uguz, M Ferro, A Ivanov, TA Sjöström, DT Simon, M Berggren, GG Malliaras, C Bernard. Controlling epileptiform activity with organic electronic ion pumps. Adv. Mater. 27, 3138 (2015). 57. EO Gabrielsson, A Armgarth, P Hammarström, KPR Nilsson, M Berggren. Spatiotemporal control of amyloid-like Aβ plaque formation using a multichannel organic electronic device. Macromol. Mater. Eng. 301, 359 (2016). 58. DT Simon, KC Larsson, D Nilsson, G Burström, D Galter, M Berggren, A Richter-Dahlfors. An organic electronic biomimetic neuron enables auto-regulated neuromodulation. Biosens. Bioelectron. 71, 359 (2015). 59. A Jonsson, S Inal, I Uguz, AJ Williamson, L Kergoat, J Rivnay, D Khodagholy, M Berggren, C Bernard, GG Malliaras, DT Simon. Bioelectronic neural pixel: Chemical stimulation and electrical sensing at the same site. Proc. Natl. Acad. Sci. 113, 9440 (2016). 60. DJ Poxson, M Karady, R Gabrielsson, AY Alkattan, A Gustavsson, SM Doyle, S Robert, K Ljung, M Grebe, DT Simon, M Berggren. Regulating plant physiology with organic electronics. Proc. Natl. Acad. Sci. 114, 4597 (2017).


Conjugated Polymers

61. J Edberg, O Inganäs, I Engquist, M Berggren. Boosting the capacity of all-organic paper supercapacitors using wood derivatives. J. Mater. Chem. A 6, 145 (2018). 62. D Zhao, H Wang, ZU Khan, JC Chen, R Gabrielsson, MP Jonsson, M Berggren, X Crispin. Ionic thermoelectric supercapacitors. Energy Environ. Sci. 9, 1450 (2016). 63. A Malti, J Edberg, H Granberg, ZU Khan, JW Andreasen, X Liu, D Zhao, H Zhang, Y Yao, JW Brill, I Engquist, M Fahlman, L Wågberg, X Crispin, M Berggren. An organic mixed ion–electron conductor for power electronics. Adv. Sci. 3, 1500305 (2016).

23 Conducting and Conjugated Polymers for Biosensing Applications 23.1 Introduction.......................................................................................697 23.2 Optical Properties of Conjugated Polymers and Associated Transduction Mechanisms...............................................................699 23.3 Electronic Properties of Conjugated Polymers and Associated Transduction Mechanisms.......................................... 701 Polymer-Based Transistors

C. Pitsalidis, A.M. Pappa, C.M. Moysidou, D. Iandolo, and R.M. Owens

23.4 Biorecognition Element Immobilization/Integration on/ with Conjugated Polymers...............................................................704 23.5 Conjugated and Conducting Polymer Biosensor Applications Based on the Biorecognition Element.....................707 Nucleic Acid Sensors • Proteins • Lipids • Bacteria •  Cells  •  Toward More Biomimetic Systems for Biosensing

23.6 Perspectives........................................................................................722 References.......................................................................................................732

Although most of the currently available conjugated and conducting polymers were originally developed for non-bio applications, a growing number are now being utilized in a variety of transducer formats including electrodes, thin-film transistors and more, for biosensing. The advent of conducting polymer devices used for interfacing with biology is motivated by a desire to bridge the gap between the transducer and the biorecognition element by using carbon-based materials found in biological systems as the active material in the transducer. The general features of these materials, which render them attractive for biosensing, include optical transparency in a range useful for cell biology, decreased rigidity compared to traditional electrode materials, chemical tunability, ease of biofunctionalization and label-free electronic transduction. In this chapter, we will summarize recent advances in biosensing using conducting polymers, highlighting the novel uses of these highly functional materials for increasingly complex biological applications. We will further show the trend toward using conducting polymers in clinical applications. Finally, we will discuss novel conducting polymers that have been designed and synthesized specifically with biosensing in mind.

23.1 Introduction Today, the term biosensor refers to any analytical device that incorporates a biological component (or at least a derivative component) associated with a physicochemical transducer. In a typical biosensor setup, the presence of specific chemicals or end products of a given biochemical reaction will generate 697


Conjugated Polymers

a measurable signal, the magnitude of which will depend on the concentration of the target analyte.1 A century ago, Cremer et al. were the first to demonstrate the concentration dependence of an acid in a liquid with the electrical potential arising between parts of the fluid located on opposite sides of a glass membrane, today known as pH sensing. The term father of biosensors, however, was given 60 years later to Leland C. Clark Jr.2 who, along with his coworkers, introduced the enzyme electrode at the New York Academy of Science Symposiums in 1962. Their concept included the use of the enzyme glucose oxidase (GOx) entrapped in an oxygen (O2) electrode via a semipermeable membrane for the detection of β-Dglucose. They found that as GOx reacted with β-D-glucose, O2 was consumed, and its decrease was proportional to the concentration of β-D-glucose in the solution.3 This revolutionary idea of Clark et al. led to one of the most important patents (1965) in the biosensing industry for the use of specific enzymes to convert various substrates and subsequently electrochemically detect the consumption (i.e., O2) or the generation (i.e., hydrogen peroxide, H2O2) of electroactive species. A few years later, Yellow Springs Instruments (YSI) developed the first commercially available glucose “first-generation” biosensors for the direct measurement of glucose from a blood sample (Model 23, YSI, 1974).2 Drawbacks associated with those “first-generation” biosensors were mainly related to O2 concentration dependence and the need for the application of high potentials to electrochemically measure the generated H2O2. This led to the replacement of the natural co-substrate with an artificial redox compound, known as a mediator, facilitating the direct transfer of electrons between the immobilized biomolecules and the electrode.4 In fact, the use of ferricinium (the oxidized form of ferrocene) as an electron transfer mediator opened up the new era of “second-generation” glucose biosensors, also known as mediated amperometric glucose sensors (1984), with the first device reaching the market in 1987 and forming the basis of the devices that still dominate today’s biosensor market.4 Apart from the commercially dominating field of glucose sensing, biosensors targeting other physiological biomolecules (i.e., proteins, small metabolites, nucleic acids, lipids, etc.) and integrated with different transducers have been developed for diverse applications ranging from medical diagnostics, food safety and environmental monitoring, to defense and security applications.5 A typical biosensor setup, as represented in Figure 23.1, incorporates three key components; the interfacing biological element, known as the biorecognition element (i.e., enzymes, organelles, cells, tissues, antibodies, receptors and nucleic acids); the transducer (i.e., electrochemical, optical, acoustic and calorimetric); and the sample containing the target analytes. Among the different physicochemical signals measured by the various transducers, electrons and photons are mostly used. As such, optical and electrochemical transducers have dominated in biosensors to date. Optical methods are usually based on absorbance, reflectance or fluorescence emission, with

FIGURE 23.1  Main components of a typical biosensor setup.

Conducting and Conjugated Polymers for Biosensing Applications


the latter representing the most often used. Electrochemical methods typically rely on the generation of either a measurable current (amperometric) or a measurable potential (potentiometric) or alterations in the conductivity of a medium (conductometric) between electrodes. More recently, other techniques such as impedance spectroscopy,6 and field effect transistor technology which measures the current as a result of a potentiometric effect at a gate electrode,7 have emerged. Electrochemical techniques are generally thought to be more advantageous compared to optical techniques due to their direct and label-free character, their independence of solution color or turbidity, as well as their compatibility with miniaturization and microfabrication methods. The main challenge in electrochemical transducers lies in overcoming the often-inefficient electron transfer between the enzyme and the electrode surface due to the spatial separation of the donor–acceptor pair. This is primarily due to the inability of the enzyme to orient itself favorably with respect to the electrode surface. Other challenges in biosensor design are generally associated with non-specific binding and sometimes transducer fouling, often encountered in real-world sample matrices, limiting selectivity and sensitivity, respectively. In this respect, choosing the right interfacial materials to bridge the biological and electronic worlds is a very important step.8,9 Conjugated polymers (CPs), which are organic macromolecules with a backbone chain of alternating double and single bonds, possess unique optical and electrochemical characteristics. Their intriguing combination of electrical and optical properties close to those of metals and semiconductors, their versatility in processing as well as their mechanical properties have stimulated their use, especially as functional bio-transducers.10 They display signal amplification (compared to their small molecule counterparts) and their structures can be easily tailored to adjust the solubility, absorption/emission wavelengths, energy offsets for excited state electron transfer and/or for use in solution or in the solid state. Their sensitivity and selectivity for biosensing are primarily determined by the difference in their optical or electronic properties prior and post exposure to the target analyte. For example, the presence of the target analyte can change the number of excited states and/or the mobility of charge carriers, changing their optical spectrum and/or their conductivity, respectively. Alongside the direct transduction schemes that those materials offer, their combination with state-of-the-art processing technologies including printing and microfabrication allows miniaturization, portability and high-throughput analysis, thereby rendering them a very attractive class of materials at the forefront of biosensor research.11 In this chapter, we review state-of-the-art bio-transducers based on conjugated and conducting polymer technologies, focusing on optical and electrical signal transduction schemes. In addition to the aforementioned classical biosensing using biorecognition events, we also show key examples of biological sensing (i.e., cell activity or integrity) from the past decade, to provide a broader perspective and highlight the great potential of such technologies where chemistry and transduction are tightly linked. In particular, Sections 23.2 and 23.3 introduce the various transduction mechanisms in conjugated and conducting polymer– based biosensors associated with their optical and electronic intrinsic properties, respectively. Section 23.4 discusses the different biofunctionalization schemes for conjugated and conducting polymers, to introduce the biorecognition elements of choice. We thus provide state-of-the-art examples and compare the various functionalization strategies with respect to the application of interest. Finally, Section 23.5 summarizes the state of the art in conjugated and conducting polymer–based biosensing categorized by the type of biorecognition element used (i.e., proteins, nucleic acids, lipids, cells and microorganisms).

23.2 Optical Properties of Conjugated Polymers and Associated Transduction Mechanisms Optical sensing mechanisms are based on changes in light refraction or propagation and fluorescence, luminescence or colorimetric changes. The working principle of CP-based optical sensors relies on changes in the optical properties of the polymer when it interacts with an analyte of interest, which induces oxidation/reduction or protonation/deprotonation of the polymer. Responses to these changes have been exploited for the development of several sensing mechanisms.11,12


Conjugated Polymers

The optical properties of CPs are associated with their delocalized electronic structure, as are the electrochemical properties, as we will see in Section 23.3. The alternating single and double bonds between the atoms on the polymer backbone create p-orbital interactions, forming a “molecular wire” that is responsible for the efficient ultraviolet (UV)-visible light absorption or emission.11–15 In particular, the width of the polymer’s bandgap is responsible for its color.12 Hence, changes in the chemical nature, redox or protonation state, solubility, absorption/emission wavelength or intramolecular conformation can strongly affect the intrinsic structure of the polymer and subsequently its absorption and fluorescent properties. Therefore, CPs can be highly sensitive transducers of biological events, such as the absorption of an analyte, affecting their aforementioned properties.11,12,15 In particular, the ability of CPs to translate a binding/dissociation event into an easily detectable and measurable optical response, the versatility in tailoring their structures and the large signal amplification that they exhibit, due to the collective response of the repeated CP units, favors their implementation in optical sensing platforms.13 As such, several CP biosensors have been developed, which in the presence of the molecule of interest, undergo either a change in fluorescence or a change in color.14 In such optical biosensors, emission signal amplification or quenching phenomena are based on exciton migration along the conjugated polymeric chain. The binding of an analyte to a specific receptor locally on the CP chain creates a signal that travels across the entire chain. Therefore, the recognition event results from the collective response of many conjugated units, as opposed to small molecule– based sensors. This is particularly important for sensing tasks in which the molecule of interest is present in extremely low concentrations.11,13 CP-based optical biosensors rely on three basic detection mechanisms: a fluorescence-enhancing approach (turn-on), a fluorescence-quenching approach (turn-off) and a fluorescence or visible color change.11,14 The simplest mechanism in CP-based optical detection is based on visible colorimetric changes. In this case, the binding event alters the conjugation length of the CP backbone, resulting in the alteration of the wavelength at which the polymer absorbs light (Figure 23.2).11 These sensors allow for signal detection with the naked eye. The latter detection mechanism was recently reported by Lee et al. Colorimetric imidazolium and imidazole-derived polydiacetylene (PDA)-based biosensors were shown not only to act as probes, but also to exhibit antibacterial effects. When bacterial strains, including methicillinresistant Staphylococcus aureus and extended-spectrum β-lactamase-producing Escherichia coli, are in contact with the biosensor, the electrostatic interaction between the positively charged polymer and the negatively charged bacterial membrane leads to the disruption of the membrane and the death of bacterial cells. This event was easily detectable by a rapid colorimetric change in the solution, i.e., from blue to red.16

FIGURE 23.2  Illustration of the various optical transduction schemes of conjugated polymers benefitting from their intrinsic optical properties.

Conducting and Conjugated Polymers for Biosensing Applications


As previously discussed, the fluorescent properties of CPs result from their characteristic delocalized electronic structure. Due to this structure, CPs can exhibit strong luminescence that can be altered significantly under perturbation of their electronic network. Hence, under photoexcitation, electron–hole pairs are generated and move along the polymer backbone. This exciton migration forms the basis for the signal amplification or quenching phenomena, described as the “molecular wire effect” by Swager in 1998.17 Based on these properties, highly sensitive fluorescence biosensing systems have been developed, where a minor interaction between the CP and an analyte of interest disrupts the electronic network of the polymeric chain or alters the electronic density within the polymer, leading to amplification of the fluorescence response of the polymer.11,13,18–20 Common terms for CPs with such behavior include amplifying fluorescent polymers and amplified quenching (or super-quenching) polymers.14,21,22 Also, semiconducting CPs such as polypyrrole (PPy), poly-(3,4-ethylenedioxythiophene) (PEDOT) and polyaniline (PANI) functionalized with specific biological molecules have been used as fluorescent transducers for the detection of DNA, proteins, enzymes and other biomolecules, as thoroughly discussed in Section 23.5. The working principle of fluorescent biosensors is based on two different modes of signal transduction. The first detection mechanism, known as the “turn-on” mode, involves binding of the molecule of interest at the functional receptor(s) on the CP chain, which interrupts the electron density of the CP chain or induces conformational changes, resulting in fluorescence of the polymer.11,14 By contrast, in the “turn-off” mechanism, binding of the target molecule to the receptor induces quenching of the CP fluorescence through electron or energy transfer (Figure 23.2).11,13 Both mechanisms are shown in the biosensing scheme developed by Cao et al. The authors used the cationic CP polyfluorenylene phenylene to screen α-glucosidase inhibitors (AGIs): in the presence of para-nitrophenyl-α-D-glucopyranoside and α-glucosidase, the fluorescence of the CP is quenched; however, when AGIs are added, the fluorescence of the polymer is enhanced.23 In addition, in both modes (on and off) the signal can be enhanced by using fluorescence resonance energy transfer (FRET). This approach involves energy transfer between two chromophores (from a CP to a fluorophore or to a quenching molecule).11 In such sensing tasks, ratiometric fluorescence measurements can be easily obtained from significant alterations in emission profiles.14 Such a biosensing scheme was proposed in 2010 by Feng et al. for DNA methylation detection in cancerous cells, using an optically amplifying cationic CP (CCP; poly(​(1,4-​pheny​lene)​-2,7-​[9,9-​bis(6​ʹ-N,N​,N-tr​imeth​yl ammonium)-hexyl fluorene] dibromide)). The DNA obtained from the cancerous cells was treated with a methylationsensitive restriction endonuclease to exclude unmethylated DNA sequences and then polymerase chain reaction (PCR) was performed to incorporate fluorescein-labelled deoxynucleotide triphosphates into the methylated DNA sequence of interest. The detection of this gene’s methylation is based on triggering the FRET signal from the CCP to the fluorescein that is incorporated into the DNA.24 Finally, it is worth noting that optical CP–based detection can occur both in the solid state as well as in solution, enabling direct integration with the biological milieu, given the fact that most biological targets require aqueous environments to maintain their functions and structural integrity. Overall, CPs optical detection schemes have been shown to be simple and convenient, in many cases allowing direct and label-free sensing of molecules.

23.3 Electronic Properties of Conjugated Polymers and Associated Transduction Mechanisms Intrinsically conducting polymers are mainly used as active materials for electrical/electrochemical biosensing. Their electronic conductivity favors their use for electrical transduction leading to high sensitivity to small perturbations caused by the binding event between the polymer and the biorecognition element. Further, their often mixed ionic and electronic conductivity allows for direct transduction of biological signals. All these features enable the development of biosensors that provide highly sensitive, accurate, fast, in-line, label-free measurements.25,26


Conjugated Polymers

As in any conventional biosensor, when the CP is in contact with a sample, including the bioanalyte of interest (e.g., cells, metabolites, proteins, lipids and DNA), the biorecognition element/receptor binds to the molecule of interest, inducing changes in the electrical properties of the conducting polymer (Figure 23.3). These changes are then converted into electrical signals via the transducer and can be monitored via several techniques. These techniques depend on the type of readout measurements: potential differences – potentiometry, current measurements – voltammetry/amperometry and impedance measurements – electrochemical impedance spectroscopy (EIS).25,27 In turn, the measurement mode depends on the type of conjugated/conducting polymer, and particularly on its reduction/oxidation potentials. The working principle of biosensors that are based on current measurements (amperometry) relies on the application of a potential between two electrodes (working and reference) and the measurement of the current on the working electrode, as a result of the electrochemical reaction taking place on the surface of this electrode.25 In the case of voltammetry, a range of potentials, defined by the reduction/oxidation current of the CP used, is applied and the responding current is measured as a plateau, which is proportional to the amount of bioanalyte present in the sample. In amperometric biosensors, a constant potential is applied for a period of time to the working electrode and the alterations in current induced by the electrochemical reduction/oxidation are measured.25,27,28 For example, Shrestha et al. developed

FIGURE 23.3  Electronic transduction schemes in conjugated polymers benefitting from their intrinsic electronic properties. (Reproduced from25.)

Conducting and Conjugated Polymers for Biosensing Applications


an amperometric glucose biosensor based on a hybrid film of PPy-Nafion-functionalized multiwalled carbon nanotubes (MWCNTs), on the surface of which they immobilized a chitosan–GOx complex. By applying a potential difference between their working CP electrode and the counter electrode, they observed a current increase as glucose concentration increased, as a result of the electrocatalytic enzymatic reaction of glucose.29 Potentiometric biosensors work under equilibrium conditions, at zero current, monitoring the amount of charge accumulated as a result of the binding event of the bioanalyte of interest to the biorecognition element (e.g., protein, enzyme and antibody) associated with or bound to the working electrode.28 EIS is widely used for the characterization of the electrical properties of interfaces, such as the resistance and capacitance of semiconductors and the transepithelial resistance of barrier-forming cells. In biosensing applications based on this technique, an alternating potential is applied to the biosensor, forcing current to flow through the transducer, receptor and bioanalyte. An alternating current can afford advantages where biorecognition elements are sensitive to a prolonged direct current or electric field. This event alters the electrical properties of the conducting polymer, such as the charge transfer resistance and capacitance, which are monitored as electrical readouts. These measurements are then utilized to analyze the biorecognition event.25,27 For example, EIS is commonly used in DNA biosensors, as we will see in more detail in Section 23.5.1, for monitoring the hybridization process between the probe DNA and the target sequence reflected by changes in the charge transfer resistance.30

23.3.1 Polymer-Based Transistors The use of a polymeric material as the active component of a transistor was first established by Wrighton et al., who, in the early 1980s, developed the so-called organic electrochemical transistor (OECT).31 However, at the time no specific application was proposed. Since then, polymer transistor technology has entered the biomedical arena as an especially promising tool due to its miniaturization and facile integration into portable electronic devices. Additionally, as the main advantage of a transistor is its ability to amplify and control the input signal, transistors have proven to improve sensitivity in biosensing compared to passive electrodes. Electrolyte-gated (polymeric) transistors (EGPTs) have so far dominated in CP biomedical applications, due to their fundamental mode of operation that uses the electrolyte as a component of the gate electrode. EGPTs are three-terminal devices in which two electrodes, the source and the drain, are connected via a conducting polymer and the third electrode, the gate, is separated from the polymer by an electrolyte, directly determining the current that flows in the channel of the transistor.7,32 Because an aqueous electrolyte solution is the natural environment for biological receptors, such polymeric transistors are compatible with biological reactions and components. To date, two main types of EGPTs exist, those that rely on the capacitive double-layer formation at the electrolyte–channel interface (field effect) and the OECTs that rely on the bulk interactions between the electrolyte ions and the transistor channel (Figure 23.4).7 In field-effect EGPTs, primarily used as immunosensors, the sensing mechanism generally relies on alterations in their threshold voltage as a result of a biorecognition event. Several models have been proposed, taking into account the nature of the biorecognition event and the type of polymer channel. These include the introduction of charge carriers in the polymeric channel that can act as traps or dopants or the formation of a surface dipole leading to capacitive coupling phenomena.33–35 OECTs, on the other hand, predominantly fabricated with PEDOT:polystyrene sulfonate (PSS) as the active material in the channel, exhibit record amplification values (transconductance) as a result of the higher capacitance of the volume of the channel, allowing for higher sensitivity in quasi-static biosensing events.36 OECTs have been mostly used as an alternative to conventional electrochemical transducers with a wealth of applications in catalytic sensors (enzyme-based sensors), cell activity and tissue integrity monitoring as well as some applications in immunosensing.37 The particular applications of these two types of polymer transistors for biosensing will be thoroughly discussed in Section 23.5.


Conjugated Polymers

FIGURE 23.4  Different types of electrolyte-gated polymer transistors. Red circles indicate where the biosensing event takes place. (Reproduced from 7.)

23.4 Biorecognition Element Immobilization/ Integration on/with Conjugated Polymers The functionalization of CPs can largely boost their intrinsic properties (described in Sections 23.2 and 23.3), providing the opportunity to tailor/adjust them to various applications. For example, improving CPs’ solubility/wettability in an aqueous environment can enhance their performance in a physiological milieu. Similarly, functionalization strategies can be applied to improve cell–surface interactions, to reduce (bio)fouling, to induce antibacterial properties and to provide drug delivery capabilities when used in in vivo or in vivo–like environments. Readers who are interested in the general functionalization strategies of CPs in order to improve the polymer’s biocompatibility, and the biotic–abiotic interface are referred to recent comprehensive reviews.38,39 Of particular interest here is the “biofunctionalization” of CPs, which involves the incorporation of suitable biomolecules or functional groups (the biorecognition element) to introduce selectivity and sensitivity to the CP-based biosensors. Such biocomponents can be nucleotides, proteins (i.e., enzymes or antibodies), microorganisms, cells, tissues, lipids, etc.40 Generally, within a biosensor platform, the biorecognition element is vulnerable to the extreme physiological conditions in biological milieu such as the temperature, the pH and the ionic strength. Given the typically short lifetime of biomolecules when they are in the liquid phase, in order to maintain their activity, they have to be preserved in an appropriate matrix or surface. A number of techniques such as physical adsorption, cross-linking, gel entrapment, covalent coupling, etc., have been used to immobilize biological molecules on the surface of CPs, while various matrices have been used as carriers of the biomolecules including membranes, gels, carbon, graphite, silica, polymeric films, etc.38,39 Examples of different biofunctionalization schemes applied for biorecognition element incorporation within a conjugated and conducting polymer system are given in Figure 23.5. As can be seen, CPs have attracted much interest as a suitable matrix themselves for the entrapment of biomolecules such as enzymes. Hence, typically, integration of the biorecognition element into the CP system mainly involves two distinct strategies, both benefiting from the CP’s versatility in processing either during or post-synthesis: (i) the incorporation of the biorecognition element during the chemical or electrochemical synthesis of the CP (in many cases acting as a primary or secondary dopant) and (ii) the

Conducting and Conjugated Polymers for Biosensing Applications


FIGURE 23.5  Examples of modification strategies used with conjugated and conducting polymers. CPs can be modified chemically and physically using a number of approaches as a means to change conductivity, bioactivity, and physical topography/geometry. A few common modification approaches are showcased, including non-covalent chemical methods (molecule entrapment and adsorption) and covalent chemical conjugation. (Reproduced from 54,59–62).

immobilization/attachment of the biorecognition element post-synthesis using conventional methods. Physical adsorption is the simplest method of surface immobilization and one of the first and most widely used for enzyme and DNA incorporation in CP-based biosensors.41,42 Despite the apparent simplicity of this method, lack of control over the amount of adsorbed molecules, stability issues due to the weak non-covalent forces involved as well as the decreased sensitivity of the sensors because of the spatial limitations of the adsorbed monolayer have prompted the development of physical entrapment during electropolymerization. The choice of biofunctionalization strategy and its parameters highly depends on factors such as the accessibility of the biorecognition elements when buried in the CP film or the risk of them leaching out in the first place and in the second place, on their long-term stability and their orientation and conformation when attached to the CP surface. From a material/device functionality perspective, it is equally important that the biofunctionalization method used does not deteriorate or alter their performance. The incorporation of biomolecules into electrodeposited CP films43 during polymerization is a widely used strategy since it allows the localization of the biomolecules on electrodes of any size or geometry. GOx has been incorporated in CPs during electropolymerization, representing a one-step fabrication thanks to co-electropolymerization from an aqueous solution containing the enzyme.44–49 The chemical activity of the immobilized enzyme can be indirectly evaluated by monitoring glucose oxidation using the electropolymerized CP enzyme-modified electrode. As expected, the incorporation of biological components, especially in the case of enzymes, can result in decreased activity depending on the parameters and the methodology chosen for the immobilization. Hence, a relatively new concept in biosensing involves the use of synthetic molecules that act as biomimetics of the biorecognition element, also known as molecularly imprinted polymers (MIPs), selectively binding to the analytes of interest. For example, in a recent study, MIP-based protein-sensing films were developed by substrate-guided dopant immobilization with subsequent CP (PPy) film formation. Indeed, using sequential interactions of the dopant


Conjugated Polymers

and pyrrole with the substrate, followed by electrochemical polymerization, the obtained films showed good selectivity and sensitivity for the protein substrate, ricin toxin chain A.50 Biomolecule incorporation during CP synthesis has also been attained via chemical cross-linking. Based on this strategy, coimmobilization of GOx in polymer acid–templated PANI with water-soluble diepoxide resulted in an electron-conducting, glucose-permeable, redox hydrogel that electrically wired the enzyme, allowing the electrocatalytic oxidation of glucose.51 Alongside their versatility in synthesis, CPs exhibit flexibility in the available chemical structure, which can be easily modified by conventional surface functionalization techniques to allow the immobilization of the desired biomolecule post synthesis. For example, our group reported on the incorporation of hydroxyl groups on the surface of PEDOT:PSS (PEDOT doped PSS) either by mixing the CP with the hydroxyl-abundant polymer, PVA,52 or by applying mild O2 plasma treatment in order to attain the direct attachment (by condensation) of surface-assembled monolayers (SAMs). SAMs bearing various functional groups can act as versatile linkers for the incorporation of biomolecules such as enzymes,52 oligonucleotides (ONs),53 etc. In addition, PEDOT:PSS exhibits a negatively charged surface due to the negatively charged PSS dopant, facilitating the functionalization of the CP based on electrostatic interactions. In line with this, we recently showed the formation of layer-by-layer (LbL) assemblies based on polycationic (poly-l-lysine, PLL) and polyanionic (PSS) polyelectrolytes on top of PEDOT:PSS films as a flexible approach for the incorporation of charged bioactive molecules such as ONs.54 Generally, the practicality and facile nature of this procedure has allowed the fabrication of multilayer and multi-enzyme biosensing devices using a variety of CPs.40 In another approach, polythiophene (PTh)-based films with a variety of side chains, end groups and secondary polymer chains were deposited via inkjet printing, forming active sensing films on a single chip of a chemiresistor-type sensor array. This multipurpose sensor allowed vapors of multiple polar and non-polar volatile compounds to be detected using a compound analysis technique. This work, though in the strict definition not classified as a biosensor, features the versatility of CPs in chemical structure engineering and film processing, allowing the development of multisensor arrays for highthroughput analysis in a single chip.55 Ultimately, the concept of functional group or molecular attachment on the surface of CPs can be further extended by using synthetic macromolecular chemistry, to graft polymer chains onto the CP backbone. Polymer chains, such as polymer brushes bearing the appropriate biocomponents of functional groups, can be grafted directly onto the CP by side chain engineering allowing facile biofunctionalization. Indeed, polypeptide brushes have already been used to covalently bind with enzymes.56 An interesting example of this methodology was reported by Kesik et al., who attached poly (L-Boc) to an amino-functionalized bis-EDOT, and copolymerized them electrochemically with ferrocene imidazole derivatives of dithiophene. Alcohol oxidase was then covalently immobilized at the end of the polypeptide chain, and the resulting electrode was used as an ethanol sensor.57 Similarly, Akbulut et al. engineered the surface of a PTh film with amine- or PEG-functionalized units alternating with unfunctionalized “spacer” thiophene units to immobilize laccase (via the amine groups) for catechol detection.58 The PEG brushes significantly improved the activity of the PTh backbone in aqueous solutions, facilitating electron transfer between the electrode and the solution. Similarly, a facile route to conducting polymer brushes via cyclopentadiene–maleimide Diels–Alder ligation was reported by Yameen et al., and further combined with a biomimetic polydopamine-assisted surface functionalization, as a general biofunctionalization strategy of choice for poly-3-hexylthiophene (P3HT) films.59 The choice of biofunctionalization strategy ultimately depends on the requirements of each sensing scheme, the complexity of the system and the envisaged application. A more thorough understanding of the effect of biomolecules and the (bio)functionalization steps on the intrinsic properties of CPs61 will definitely allow us to better harness the properties of CPs and further improve current technologies in the area of biosensing by enhancing signal transduction and generating closed-loop feedback systems.62

Conducting and Conjugated Polymers for Biosensing Applications


23.5 Conjugated and Conducting Polymer Biosensor Applications Based on the Biorecognition Element 23.5.1 Nucleic Acid Sensors Nucleic acids are key molecules for the transfer of genetic instructions that determine the unique characteristics of living organisms. The primary molecule responsible for this function in humans and most organisms is DNA, and in some viruses RNA. Besides DNA and RNA, other types of nucleotides exist, and are found in varying functions such as energy storage (adenosine triphosphate: ATP) and energy transfer (nicotinamide adenine dinucleotide : NADH; nicotinamide adenine dinucleotide phosphate: NADPH). The detection of nucleic acids has been the subject of extensive research for over two decades in various disciplines such as medicine, disease diagnosis, drug development and food technology.63–65 One of the most commonly used techniques for the recognition of DNA sequences is hybridization. Typically, this process involves the immobilization of a single-stranded DNA (ssDNA) sequence (probe) on top of a (functionalized) support layer, which is used to detect a complementary ON sequence (target) due to base-pairing interactions (Figure 23.6a).66 The interface between the functionalized layer and the polynucleotide formations plays a crucial role as it affects the generated recognition signal and further the precision and the performance of the device. Depending on the signal transduction method, the detection of DNA assembly can be optical67,68 or electrical.69–72 The major challenges faced when using DNA hybridization approaches involve the immobilization process on the target surfaces, the selectivity, the use of chemical labels as well as the reproducibility of the method. To date, a broad range of CPs has been used for nucleic acid biosensors such as PANI, PPy, PEDOT:PSS, PTh, polyphenyleneethynylene (PPE), etc. (Figure 23.6b).74,76–78 The immobilization of DNA probes on top of the support layer can be typically achieved by adsorption or by covalent bonding. Another approach uses non-covalent bonding onto the electroactive surface using the avidin/

FIGURE 23.6  (a) Schematic illustration of the hybridization process. (b) Immobilization of NH2-DNA and peptide nucleic acid on polyaniline film.73,74 (c) Schematic depiction of the interaction between cationic polymers and single-stranded DNA (ssDNA), double-stranded DNA (ds DNA), single-stranded PNA and PNA–DNA duplex and the corresponding fluorometric detection and fluorescence intensity graphs.75


Conjugated Polymers

streptavidin–biotin interaction. In the case of adsorption, the attachment occurs by electrostatic interaction between the negatively charged phosphate groups of DNA and the positively charged species of the activated substrate. Various cationic polymeric materials have been used for passivation of the hosting surfaces, such as PPy,79 PLL,80 polyethyleneimine (PEI),81 benzothiadiazole-fluorene copolymer derivatives or chitosan75 (Figure 23.6c). Adsorption can be achieved by electrochemical interactions using an electric field applied on a host electrode. This method allows for the electrodeposition of the DNA probes and composites without pretreatment and functionalization of the substrate. Velusamy et al. showed that the combination of electrostatic and electrochemical adsorption can enhance the immobilization and the stability of the DNA probes on PPy/Au electrodes.82 In the covalent binding technique, a synthetic DNA probe derivatized with functional groups (amines, thiols, etc.) can covalently attach to the surface of the support layer. This approach offers high specificity and stability to the resulting DNA layer. Typically, the process involves the use of thiol-modified DNA probes that can be covalently immobilized on Au surfaces due to the high binding strength of the Au–S bonds.83 Alternatively, metal surfaces (Au, carbon, indium tin oxide [ITO], graphene [Gr], etc.) can be coated with functionalized Au nanoparticles (NPs) in order to achieve more efficient DNA immobilization.84 For covalent attachment on polymeric surfaces, DNA probes are usually functionalized with carboxyl or amino groups and bind on the surface through peptide bonds using carbodiimide coupling chemistry.78,85 Recently, Galán et al. immobilized an acetylene-terminated ON probe, complementary to a hepatitis C virus target sequence, onto an azido-PEDOT polymer by covalent binding using “click chemistry”.86 Approaches that involve the binding of DNA probes onto functionalized CNTs and surfaces modified with CNT/CP composites have also been reported.77,87 In another study by Yang et al., the immobilization of DNA probes was carried out using PANI nanofiber and MWCNT hybrid nanocomposites. In the case of avidin/streptavidin–biotin, a non-covalent interaction that leads to the immobilization of the DNA probe occurs. Avidin/streptavidin is a tetrameric protein that contains four identical biotin-binding sites that can form tetravalent avidin/streptavidin–biotin bonds capable of promoting DNA immobilization on top of solid support layers. Various strategies have been reported for the avidin/streptavidin functionalization of CP electrodes toward the immobilization of biotinylated DNA probes.88,89 In addition to the hybridization step, the transduction mechanism is crucial for the conversion of the recognition event to signal. The analyte detection can be quantified as an alteration in the impedance, a change in the potential, in the current or in the conductivity of the system. CPs offer great advantages for the latter, as they exhibit low resistance and high redox capacitance, which results in better signal amplification and stability.90,91 PANI, PPy and PEDOT-based electrode devices have been extensively investigated for DNA sensing by means of EIS and voltammetry techniques (cyclic voltammetry [CV], differential pulse voltammetry [DPV]). PANI and its derivatives was one of the first CPs investigated for the realization of organic semiconducting devices due to its simple synthesis, good conducting properties (when doped), acid-doping/base-de-doping chemistry and its chemical/structural flexibility to bind biomolecules. Saberi et al. fabricated an electrochemical DNA sensor based on PANI/Au nanocomposite electrodes and they studied its sensing response using EIS and CV measurements. The use of Au NPs enhanced the performance of PANI in a neutral pH environment while simultaneously promoting the immobilization of thiolated ONs, resulting in higher hybridization efficiency.92 More recently, Zheng et al. used PANI/Gr composites for the fabrication of a free-label electrochemical DNA sensor with a detection capability ranging from 0.01 pM to 1 μM. By changing the composition ratio of the PANI/Gr, they investigated the effects on the hybridization of the DNA probe using EIS. Moreover, they found that the increase in the response of charge transfer resistance is due to double-stranded (ds) DNA remaining on the surface in the presence of PANI, while the decrease in the response is associated with the dsDNA being released from the sensor when Gr is dominant.94 Recently, Shoaie et al. used PANI and Au NPs on top of screen-printed carbon electrodes for the immobilization of biotinylated nucleic acid. Using a digoxigenin-labelled detector probe, anti-digoxigenin antibody conjugated to horseradish peroxidase (HRP), they were able to obtain stable detection in the range of 1000–0.001 pM with a limit of detection

Conducting and Conjugated Polymers for Biosensing Applications


(LOD) at 0.01 fM.93 PPy and its derivatives is one of the first CPs used for the realization of nucleic acid sensors, as it offers great biocompatibility and stability in water, as well as the ability to be syntesized at neutral pH and deposited by various techniques (electropolymerization, wet techniques, etc.). Livache et al. first used a copolymerized PPy–DNA composite for the detection of DNA hybridization on CP electrode supports.94 More recently, Booth et al. developed a PPy–pyrrolylacrylic acid (PAA) copolymer film-based sensor to detect targeted DNA sequences. In this work, two different dopants were used (PSS and LiClO4) in order to study the sensing response as a result of DNA adsorption and stability.95 Other researchers used electrochemically synthesized PPy nanowires obtained using a potentiostatic technique. The DNA in this study was immobilized to the PPy nanowire through linkage between the NH groups and the phosphate groups of the probe DNA. The sensor was capable of detecting concentrations as low as 0.1 nM of E. coli DNA while the response time of the sensing platform was 10 sec.96 In an attempt to improve the mechanical properties of PPy, Esmaeili et al. synthesized PPy-kappacarrageenan (KC) composites. To facilitate the immobilization of the thiol-modified ssDNA probe, Au NPs were electrodeposited onto the polymer composites. The use of the KC-PPy matrix containing Au NPs showed enhancement in the sensing response, while the sensors exhibited a wide detection range (from 5 × 10−18 to 5 × 10−12 M). The biosensor was used for the gender classification of Arowana fish.97 Electroactive poly(ε-caprolactone) (PCL)/PPy nanofibers functionalized with ssDNA were fabricated by Guler et al. The DNA immobilization and the interfacial characteristics of the nanofibers and hybridization were observed by EIS measurements. The authors compared the effects on the transfer resistance between DNA immobilized nanofibers and non-complementary DNA, while they quantified their observations by simulating the EIS spectra and extracting the parameters from the equivalent circuit.98 Electrochemical DNA hybridization sensing has also been performed for PEDOT derivatives and PEDOT composites. Typically, PEDOT films exhibit optimum morphological and surface characteristics that aid in the bioconjugation of functional groups and promote the immobilization and hybridization of DNA. Compared to other functionalized layers (SAM, polymer, etc.), they are more uniform and can be easily tuned manufactured. Luo et al. demonstrated the great versatility of PEDOT biointerfaces for fluorescent, quartz crystal microbalance (QCM) and electrochemical DNA detection. Using ON-grafted PEDOT thin films fabricated by electropolymerization, they immobilized the capture probe by a bioconjugation reaction. By varying the molar fraction of the monomer, the authors were able to synthesize PEDOT biointerfaces with different CP densities and optimize the output signal. The resulting detection was found to be in the range between 10 and 1000 nM.99 Radhakrishnan et al. described the use of PPy-PEDOT nanotubes functionalized with Ag NPs. This route promoted the immobilization of labelled DNA and allowed efficient hybridization. The dynamic detection range was 10−11 to 10−14 M with the lower detection limit (3 s/b) of 5.4 × 10−15 M. Interestingly, surfaces modified with PPy–PEDOT–Ag–S-ssDNA exhibited high stability while still maintaining a sensitivity of 91.5%, the initial value after 7 days in storage.100 In a recent study by Tao et al., an OECT based on PEDOT:PSS was used as a peptide nucleic acid (PNA) sensor. A gold-coated porous anodic aluminum oxide was used as a gate electrode in order to facilitate DNA sensing. The sensing mechanism of the OECT platform can be attributed to the modulation of the surface potential of the gate electrode induced by the PNADNA hybridization process, which induces changes in the measured source-drain current.101 In another OECT approach by our group, LbL assembly on PEDOT:PSS-based OECTs was used for the electrical monitoring of a nucleic acid. This platform was based on the adsorption of messenger RNA onto the positively charged PLL layer of the LbL assembly. The OECT sensor was capable of detecting the binding event of mRNA, which was quantified by the amperometric response for various concentrations.102 Due to the photosensitivity of CPs, photoelectrochemical techniques have been used to transduce the photocurrent response for the detection of DNA hybridization. In some of the first works by Lassalle et al., they studied the PPy copolymer modified by grafted ONs. The complex was exposed to complementary and non-complementary ON in order to determine the effects on the photocurrent spectra. The photocurrent in the case of hybridization was significantly lower, which was attributed to an alteration in the photosensitivity of the films due to the presence of ON.103 More recent works have used CCPs for


Conjugated Polymers

optical signal amplification of fluorescent DNA detection. Huang et al. used CCPs with different backbones and side chains and correlated their structure with the signal amplification of the FRET-based fluorescent DNA assays. Specifically, they showed that twisted conjugated backbones resulted in the most sensitive optical reporter to the difference between ssDNA and dsDNA in charge density, while the linear-conjugated backbone was more sensitive to the difference in the hydrophobic properties between ssDNA and dsDNA.104 DNA sensors using transistor technologies have been realized, benefiting from their inherent amplification properties. Kergoat et al. described a water-gated organic field-effect transistor based on poly [3-(5-carboxypentyl)thiophene-2,5-diyl] (P3PT–COOH) onto which DNA probes were covalently grafted using NHS/EDC chemistry. In an attempt to explore the Debye length effects, they used both low and high ionic strength electrolytes. As such, upon hybridization a decrease in the off current of the device was observed when using deionized (DI) water whereas no significant change occurred when using saline solutions. This was associated with the steric hindrance of DNA, which inhibits ions interacting with the semiconductor channel.105 In another approach, OECTs based on PEDOT:PSS active layers were employed in a flexible microfluidic system toward the realization of a label-free DNA sensor (Figure 23.7). In this work, the sensing mechanism relied on the modulation of the surface potential of the gate electrode by the immobilized ssDNA probe. The device was able to detect complementary DNA targets at concentrations as low as 10 pM

FIGURE 23.7  (a) Schematic diagram of an OECT integrated into a flexible microfluidic system, which is characterized before and after the modification and hybridization of DNA on the surface of an Au gate electrode. (b) Photographs of a device bent to both sides. (c) Transistor characteristics and output characteristics of an OECT measured at different bending status. The microfluidic channels filled with PBS solution. (d) Time-dependent channel current of an OECT measured after applying different gate voltages.106

Conducting and Conjugated Polymers for Biosensing Applications


using pulse-enhanced hybridization of DNA. Here, no effect due to Debye limitations was observed, presumably because the thickness of the DNA layer was much thinner than the thickness of the double layer above the gate electrode of the device.106

23.5.2 Proteins Catalytic Biosensors Catalytic biosensors rely on enzymes as recognition elements. These biosensor designs are especially attractive due to the variety of measurable signals arising from the specific catalytic reactions, which include protons, electrons, light and heat generation. Of all enzyme-based biosensors, glucose biosensors (based on the enzyme GOx) are the most studied and widely praised sensor success story, due to the rising clinical relevance of diabetes.107 As already mentioned, in enzymatic sensors the electrical communication between the enzyme and the electrode is the most challenging aspect and primarily dictates the sensitivity of the sensor. Electron transfer is thus typically mediated by including redox active molecules in the sensor architecture that act as electron relays.108 As with conventional enzymatic biosensors, also in CP-based biosensors the biosensor generations have evolved from the most standard approach of using hydrogen peroxide (the enzymatic reaction by-product) detection to using electron transfer mediators typically immobilized in the CP matrix, to the most direct approach based on direct electron transfer (Figure 23.8a–d). Indeed, in many cases, CPs, thanks to their ability to support electronic charge transport along their backbones as well as ion transport through their bulk,109 can themselves act as electron relays. The majority of the early work in this area consists of biosensors where GOx, along with the corresponding electron transfer mediators, have been successfully entrapped in PPy

FIGURE 23.8  Enzyme-based catalytic amperometric detection using CPs. (a) Schematic showing three different generations of bioelectrocatalytic enzymatic sensors and key examples using CP-based enzymatic detection of glucose based on (b) hydrogen peroxide–catalyzed oxidation using a conducting polymer/Pt hybrid network,60 (c) mediator-aided electron transfer from the enzyme catalytic site to the electrode based on mediator and enzyme entrapment on PEDOT:PSS119 and (d) direct electron transfer of glucose oxidase on PANI microtube electrodes.121


Conjugated Polymers

films.110,111 However, the first demonstration of a CP acting as a molecular wire for the enzyme without the need for mediators was shown in 1996, using doped PPy.112 Since then, several studies on CP-based enzymatic sensing have followed,113 with PEDOT showing great potential as a more electrochemically stable alternative to PPy.60,114,102 PANI has also been extensively employed in enzymatic sensing, due to the presence of two redox couples in its structure with the appropriate electrochemical potential facilitating the enzyme–polymer charge transfer processes.115 Improving the electron transfer network in PANI as well as its intrinsic conductivity by the development of three-dimensional porous networks and by the addition of inorganic composites (often in the form of nanomaterials), respectively, has resulted in the enhanced stability and improved sensitivity of PANI-based biosensors.116,117 In this context, Xian et al. reported a CP-hybrid glucose biosensor based on gold NP-conjugated PANI nanofiber system that exhibited excellent conductivity and many micron-sized gaps to easily immobilize GOx, allowing rapid electron transfer and enhanced interactions toward various glucose concentrations.118 Direct electron transfer from hydrogen peroxide was proposed using PANI microtubes and their strong electrostatic interaction with the negatively charged GOx (Figure 23.8d). The large geometric surface of PANI microtubes facilitated enzyme loading and efficient electrocatalytic activity allowing direct electron transfer from GOx.119 In a very recent work by our group, an electron transporting (n-type) CP bearing polar side chains was shown to promote interactions with the enzyme lactate oxidase, and allowed for direct electrical communication between the film and the enzyme. The mediator-free lactate detection was greatly improved when employed in a transistor configuration, due to inherent signal amplification and design flexibility. The benefits of transistor-based biosensors are described in more detail in the following section.120 In the past few years, transistor technology has been increasingly employed due to its inherent amplification function as well as its ease of integration. A glucose sensor based on a field-effect EGPT composed of P3HT was demonstrated by Bartic et al.122 They used an H+-selective membrane as a dielectric layer on top of P3HT and immobilized GOx on its surface. The device could detect local pH changes near the channel due to the presence of gluconic acid, the product of the enzymatic reaction with glucose. The majority of enzymatic, transistor-based sensors has been realized using OECT technology. The first demonstration of an OECT-based glucose sensor is attributed to Contractor et al.,123 who used PANI functionalized with GOx as the active component of the transistor. The pH-dependent conductivity of PANI allowed the indirect determination of glucose in this case too. Subsequently, OECTs for enzymatic sensing were optimized124 and investigated as sensitive hydrogen peroxide sensors39 as well as coupled with electron mediators for the direct detection of the analyte.125 The operation principle relies, in the first case, on the catalyzed oxidation of hydrogen peroxide by the gate electrode (usually bearing Pt nanostructures) that increases the gating efficiency of the transistor, hence the doping state of the channel, to an extent proportional to the analyte concentration. In the second case, the O2/H2O2 couple is replaced by a fastredox couple, such as the ferrocene/ferricenium ion couple that wires electron(s) generated to the gate electrode. The potential use of OECTs as enzymatic sensors of critical metabolites has since been applied/ exploited/investigated in biological environments such as sweat126 and saliva127 as well as in human breath (Figure 23.9a)128 paving the way for point of care, even wearable electronic technologies.129 Recently, we showed the development of multi-enzyme functionalized OECT arrays integrated with microfluidics for the simultaneous detection of different metabolites including glucose lactate and cholesterol from one drop of saliva (Figure 23.9b).130 OECT circuits for the detection of lactate from tumor cells131 were also developed by our group, as a more integrated reference-based biosensor for improved sensitivity in the highly complex cell culture media. In all those platforms, miniaturization, system integration and enhanced sensitivity are facilitated by the use of organic transistor technology. Affinity Biosensors With the notable exception of the glucose sensor, the majority of diagnostic tools employ antibodies for the recognition, identification and quantification of target analytes. Antibody use was sparked by the advent of monoclonal antibody (mAb) technology132 that used cell clones that could specifically

Conducting and Conjugated Polymers for Biosensing Applications


FIGURE 23.9  Examples of polymer-based transistor biosensors employing proteins as biorecognition elements. Catalytic OECT biosensors using enzymes for (a) alcohol detection in breath128 and (b) multiple analyte detection using an array of transistors and microfluidics.130 Affinity field-effect transistor-based biosensors able to detect (c) proteins beyond Debye limitations138; (d) enantiomers binding to odor-binding proteins140; (e) the concept of floating-gate transistor technology for affinity-based sensing141; and (f) OECT immunosensors for the catalyzed detection of cancer biomarkers.144

produce mAbs of choice. Antibody-based sensors make use of the sensitivity and specificity of antibody–antigen interactions. This selective protein-binding interaction (usually immobilized on a surface) elicits an electronic or optical signal related to the amount of analyte present in the sample. Lee et al. developed single PANI nanowires for the detection of immunoglobulin and myoglobin. This system showed great promise for cardiac marker detection.133 Sun et al. developed an immunosensor for detecting chlorpyrifos by using Au NPs and PANI/MWCNT-CS nanocomposites, with the Au NPs greatly enhancing the electrochemical signal as well as the adsorption capacity of the antibodies.134 A conductimetric reagentless immunosensor using antibodies has been fabricated using the CP PEDOT as both the immobilization matrix and the transducer. The sensor was able to detect either the antigen or the antibody.135 Mouffouk et al. described a method to functionalize an EDOT derivative with


Conjugated Polymers

biotin by physical entrapment. By optimizing the electrodeposition parameters, they were able to confine the CP film in a 10 μm diameter Pt disk microelectrode and detect avidin at the 10−13 M level.136 Field-effect polymeric transistors have been successfully used as biosensors to detect binding events. In 2013, Suspene et al. reported on the sensing of streptavidin with an EGPT based on biotinylated P3HT as the active sensing and semiconducting material.137 The biotin–streptavidin couple was chosen as a model binding system given its low dissociation constant (high affinity). In an attempt to elucidate the transduction mechanisms of such devices, Palazzo et al.138 investigated the sensitivity of a P3HT EGPT as a function of the Debye length, the receptor charge and the distance of the binding event using the model system biotin–avidin as well as antigen–antibody interactions (C-reactive protein CRP and antiCRP, respectively). The sensor was shown to successfully detect binding events occurring at distances 30 times the Debye length value from the transistor’s channel even at high salt concentrations. The sensing mechanism was attributed to capacitive changes at the channel–electrolyte interface. Indeed, it has been ascribed to the formation of Donnan equilibria within the protein layer, resulting in an extra capacitance in series to the gating system, which is insensitive to the Debye length value (Figure 23.9c). In a slightly different sensing approach, Casalini et al.139 reported a P3HT-based EGOFET biosensor for dopamine detection where a gold gate was used as the sensing area instead of the semiconductor. They modified the gate–electrolyte interface with a self-assembled monolayer of cysteamine and 4-formylphenyl boronic acid, enabling the selective covalent binding of dopamine that, in turn, modulated both the work function of the gate electrode and the capacitance of the electrode–electrolyte double layer. More recently, Mulla et al.140 reported on the binding of (S)-(+)- and (R)-(−)-carvone enantiomers to an odorant-binding protein (OBP) mutant, by using a back-gated EGPT based on P3HT. With this approach, the authors were able to discriminate the binding of two molecules for which the difference in binding energy differs, the OBP functionalized Au gate being as low as 1.1 ± 0.5  kJ/mol (Figure 23.9d). Electrolyte-gated transistors that incorporate floating gates were recently reported by the Frisbie group as a sensitive platform for the label-free electronic detection of proteins. These devices have proven sensitivity on both capacitance and work function changes on the floating gate, associated with chemical binding events.141 Using this technology, the same group demonstrated quantitative detection of ricin from complex food matrices,142 with improved performance due to noise suppression from the reference device integration and AC testing modality (Figure 23.9e). OECTs have been used only sparsely for protein sensing, using electrostatic interactions between proteins and the organic semiconductor channel. Kanungo et al. successfully fabricated OECT-based immunosensors. An electrochemically polymerized PEDOT layer served as both the active layer of the transistor and the immobilization layer. The devices were characterized by measuring the conductance as a function of time after exposing them to the antigen solution.135 Changes in the conformation of the CP, following the antibody∶antigen binding, was suggested as the sensing mechanism despite the earlier studies by Swager and coworkers suggesting that there may not be a significant conformational change in the CP upon binding of the antigen to the surface-confined antibody.143 More recently, Fu et al. developed an OECT immunosensor targeted to the cancer biomarker, human epidermal growth factor 2, having the surface of the gate electrode functionalized with catalytic nanoprobes (Au NPs that are functionalized with specific antibodies and the electrochemically active enzyme HRP). The devices were operated by detecting electrochemical activity on the gate electrodes, which is dependent on the concentrations of the target proteins detected by the catalytic nanoprobes. As such, the sensing mechanism of the device was attributed to the electrochemical reaction catalyzed by the nanoprobes on the gate, resulting in high sensitivity144 (Figure 23.9f).

23.5.3 Lipids One of the key roles of lipids is to contribute to compartmentalization in biological systems. This is an essential process in biology needed in the first instance for protection, but also to define environments with appropriate physical conditions (i.e., for discrete processes such as a defined enzymatic process).

Conducting and Conjugated Polymers for Biosensing Applications


In all cells, the lipid bilayer consists of amphipathic lipids with a hydrophilic headgroup and a hydrophobic tail. Generally, the lipids are phospholipids and have a two-tailed structure. Commonly found (phospho)lipids include phosphatidylethanolamine, phosphatidylserine and sphingomyelin. Phospholipids have a number of conserved properties including the fact that they have one saturated aliphatic tail and a second tail with an unsaturated bond. This results in a particular structure that makes lipids capable of forming a sealed compartment rather than a micelle (e.g., single-tailed phospholipids). Phospholipids give rise to fluid and dynamic structures where proteins can be docked. This fluidity is very important for biological function. An extremely important property of lipid bilayers is their impermeability, with only small molecules allowed to freely cross them. The passage of water and other polar molecules is extremely limited and the passage of ions is highly regulated. This is governed by ion channels, protein complexes specific for certain ions, being able to discriminate ions not only according to their charge but also according to their size. This is essential for biological function, as ion gradients are maintained across biological membranes and control a wide variety of processes including neuronal action potentials and energy generation. The impermeability or modulation of the ion flow is a property that is often exploited for lipid-based electrical detection mechanisms. Noh et al. developed a biosensor for the detection of phthalate esters (PE) based on functionalized phospholipids bonded to the CP and including a microfluidic component needed for sample concentration with an electrochemical biosensor.145 As reported in Figure 23.10a, the electrode is composed of a conductive layer obtained by electropolymerizing the terthiophene monomer TTBA (2,2′:5′,2″-terthiophene-3′(p-benzoic acid)) onto gold NPs. PolyTTBA was functionalized using toluidine blue O (TBO) along with 1,2-d​ioleo​yl-sn​-glyc​ero-3​-phos​phoet​hanol​amine​ (DOPE) in order to obtain a hydrophobic and positively charged biomimetic surface layer to interact with negatively charged target molecules. The samples were analyzed by electropherograms that allowed to discriminate among the different components and to quantitatively assess their abundance in the samples. The developed device guaranteed very high selectivity and sensitivity with the ability to distinguish among the different PEs extracted from real-life samples (i.e., manicure and PVC samples) and to detect PE levels between ∼12.5 and ∼35.2 pM. Importantly, the biosensor also proved to be applicable for the analysis of cell extracts when mammalian kidney epithelial cells (Vero) were exposed to the analyzed PEs. Only after approximately 45 continuous measurements did the biosensor lose just 2.7% of its initial response, proving it to be highly stable. In 2006, Kwon and collaborators developed an approach based on the functionalization of the electrode surface with a lipid layer, which was aimed at reducing interference while sensing a specific analyte by excluding hydrophilic electroactive materials from the detecting surface.149 The electrode was applied to sense superoxide and it comprised a monolayer of 1,2-dioleoyl-sn-glycero-3-succinate (DGS) covalently bonded to poly-(3,4-diamiono-2,2:5,2-terthiophene) (DATT) electrochemically grown onto gold electrodes. The conductive polymer served both as a transducer and as an anchoring point for the lipid layer. In turn, the first lipid layer was functionalized in order to obtain a biomimetic membrane, depositing either 1-pal​mitoy​l-2-o​leoyl​-sngl​ycero​-3-ph​ospha​te (POPA) or cardiolipin (CL) by the Langmuir–Blodgett technique. The lipid bilayer was used as a docking site for cytochrome c to sense the superoxide anion radical. Interestingly, the sensitivity of the Au/poly-DATT/DGS/CL/cyt c-modified electrode was about two times higher than that of the Au/poly-DATT/DGS/POPA/cyt c electrode. The approach explored by Kwon was further developed by Gurudatt et al. and exploited for direct electrochemical sensing of leukemia cells.150 Both normal and cancer cells display receptors on their surfaces for folate acid, a molecule that is internalized and used by cells to produce nucleotides, the building blocks in both DNA and RNA synthesis. Highly proliferative cells, such as the cells of myeloid leukemia, steadily overexpress folate acid receptors (FR) on their surface and they are undetectable with conventional methods such as immunohistochemistry.151 In the work by Gurudat and colleagues, the amine functionalized organic conducting polymer [2,2′:5′,2″-terthiophene]-3′,4′-diamine (DATT) was used as an electrode material and as a substrate for both drug and lipid immobilization. Differently charged lipid molecules were tested to check their impact on the voltammetric response of the sensor surface with the


Conjugated Polymers

FIGURE 23.10  Schematic representation of the electrode structure and fabrication and structure of the phthalate ester molecules and of the microfluidic device. (a) Structures of the three modified calixarene molecules (receptors)145; bottom panel: schematic structure of the receptor/phospholipid/PDA assemblies (depicting 2 as an example). Key: blue, PDA; black, phospholipids; red, receptor units; green, protein. (b) Calixarene molecules decorated with basic or acidic amino acid residues were embedded into PDA vesicles containing dimyristoylphosphatidylcholine (DMPC) as lipid component.146 The integration of lipid membranes with OECTs: the gating of the OECT can be controlled using (c) suspended lipid membranes (BLMs) on a Teflon support147 and (d) supported lipid bilayers on top of the conducting polymer (PEDOT:PSS).148

positively charged phosphatidylcholine showing a better CV response along with decreased electrode fouling. Five different drugs targeting FR, such as raltitrexed (Rtx), pemetrexed, folinic acid, folic acid and methotrexate, have been tested, with Rtx showing the best performance for cancer cell detection. Both chronoamperometric and EIS measurements were performed to evaluate electrode specificity and sensitivity. EIS measurements showed a linear range in between 1.0 × 103 and 2.5 × 105 cells/mL with a detection limit of 68 ± 75 cells/mL. Interestingly, EIS showed a different response when the electrode was exposed to the same number of normal and cancerous cells, providing a tool for the detection of human acute T leukemia cells. It is important to highlight that this result was not achieved without lipid conjugation, highlighting the role played by these molecules in the recognition event. Another interesting example of lipids playing a pivotal role in the biosensor architecture has been reported in a work where the CP PDA has been used in conjunction with phospholipids for the colorimetric detection of water-soluble proteins.146 The ability of PDA materials to respond to changes in environmental parameters together with the ease to form self-assembled structures are two key advantages of these materials that have prompted their use in the development of biosensors.152 PDA has an intense blue color (650 nm) due to its absorption of the ene–yne cross-linked framework in the visible region. Structural changes introduced in

Conducting and Conjugated Polymers for Biosensing Applications


the conjugated backbone of the polymer result in the modification of the delocalized conjugated electronic networks, thereby inducing a chromatic shift from blue to red (550 and 500 nm). These changes can be due to a wide range of phenomena and, interestingly, surface perturbations in PDA, due to binding events, give rise to the blue to red shift. The ability of PDA to also experience this chromatic transition when assembled into vesicles has been exploited in a recent paper by Kolusheva and colleagues.146 In this project, calixarene molecules decorated with basic or acidic amino acid residues were embedded in PDA vesicles containing dimyristoylphosphatidylcholine (DMPC) as the lipid component (Figure 23.10b). The binding of three different proteins (pepsin, histone and albumin) to the embedded molecules produced a pattern of vesicle colors that was used to create a fingerprint for protein recognition. Overall, the developed technology allows the detection of a non-membrane protein at micromolar concentrations with the naked eye. Modified PDAs have been designed to address the recognition of lipopolysaccharides (LPS) from different bacteria.153 LPS are complex glycolipids embedded within the outer membrane of gram-negative bacteria. Rangin and Basu reported a biosensor based on PDA functionalized with tryptophan and tyrosine residues to make contact with the glycidic part of LPS. In this way, they were able to discriminate five different LPS types corresponding to five different bacteria. Interestingly, although further research needs to be performed to improve the sensor arrays, the fingerprints obtained were specific enough to allow the unequivocal identification of the five investigated LPS in a blind test. Lipids are essential components of biological membranes and have been used to obtain isolated bilayers providing a means to study transmembrane protein functions in terms of ionic flux, the binding of signaling molecules or their interaction with pathogens. For the study of ion channel recording and pharmaceutical screening, the integration of electrode materials with either suspended lipid bilayers (e.g., black lipid membranes [BLMs]) or supported lipid bilayers (SLB) has been widely reported. Attempts have been made to interface lipid bilayers with conducting polymers using the latter for signal transduction in the development of biosensors.147,148,154 The integration of OECTs with lipid bilayers was first shown by Bernards et al., who developed BLMs on a Teflon support suspended between the gate and channel of an OECT, showing that gating can be fully suppressed by the insulating BLM, and then subsequently restored by the incorporation of gramicidin ion channels, which permeabilize the membrane. The valence-dependent permeability of gramicidin enables these devices to discriminate between monovalent and divalent ions (Figure 23.10c).147 Our group later reported on the use of vesicle fusion to form an SLB on the surface of PEDOT:PSS films.54 The ability to study SLB properties using OECTs was further explored. SLB formation was induced by osmotic shock using 1,2-d​iphyt​anoyl​-sn-g​lycer​o-3-p​hosph​ochol​ine (DPhPC) and 1,2-d​ iphyt​anoyl​-sn-g​lycer​o-3-p​hosph​oetha​nolam​ine (DPhPE) isolated from archea (commonly used in the formation of BLMs for the study of ion channels). Their branched chains, being more disordered than lipids of other origin, result in more fluid membrane structures (Figure 23.10d). As we will see in the next section, PEDOT:PSS-based OECTs have been used among other applications to monitor barrier tissue layer integrity. This is because the cell layer modulates the ion flux, which is transduced via a decrease in the drain current. The ionic flux thus determines the speed at which the drain current reaches steady state. In this study, the OECT was used to detect the presence and integrity (insulating nature) of the SLB as well as the insertion of the membrane protein α-hemolysin, which resulted in a net increase of ions flowing through the SLB into the transistor channel, changing the transconductance value. This device configuration is currently being exploited by us for further study of transmembrane proteins within SLBs. Tarabella and colleagues suggested the use of a PEDOT:PSS-based OECT for monitoring micelle formation. In their study, the cationic surfactant hexadecyltrimethylammonium bromide (cetyltrimethylammonium bromide, CTAB) was used.155 It is reported that above a concentration of 1 × 10−3 M at 298 K, known as the critical micellar concentration (CMC), CTAB forms micelles. The OECT modulation was measured as a function of CTAB concentration. OECT modulation corresponds to (I – I0)/I0, where I is the off source-drain current (for gate voltages, Vgs ≠ 0 V) and I0 is the on source-drain current (for Vg = 0 V), varying from about 0.45 M for concentrations below 1 × 10−4 M up to about 0.9 at 1 × 10−3 and 1 × 10−2 M. The authors demonstrated that the transistor responds to an increase in CTAB concentration only when it is above the CMC for CTAB, proving useful for monitoring micelle formation.


Conjugated Polymers

The use of phospholipids was described by Magliulo and colleagues who developed a BIOEGOFET.156 In their study, a layer of streptavidin-functionalized phospholipids was deposited on the gate surface of an electrolyte-gated organic field-effect transistor (EGOFET) with a P3HT channel. This device presented an LOD of 10 nM with a dynamic range of 10 nM–1 μM. The advantages of the developed system, in comparison to other available solutions, lay in the wide dynamic range as well as the possibility of working in a high ionic strength electrolyte solution. The mode of action can be explained taking into account the binding of negative species/anions to the surface of a dielectric layer (the phospholipid bilayer) at the interface with a p-type material (i.e., P3HT) when a negative bias is applied at the gate. This, in turn, results in the accumulation of positive charges on the surface of the conducting polymer. The binding of the negatively charged streptavidin molecules to biotin on the phospholipid layer induces more holes in the transistor channel, registered as an increase in the measured current. Control experiments were run in order to assess the specificity of the binding event and its mechanism.

23.5.4 Bacteria A new generation of biosensors use bacterial membranes as recognition elements for the analysis of various biological substrates. For example, such membranes display surface enzyme complexes, allowing for enzyme-based detection of biomolecules. Also, when the bacterial cells are the target analyte, a wealth of biorecognition elements is available, including antibodies, bacteriophages and lectins.157 Another well-studied bacterial sensing approach involves quorum sensing (QS), the bacterial cell–cell communication process that encompasses the production, detection and response of bacteria to extracellular signaling molecules.158 Along these lines, Zhang et al. demonstrated the use of a water-soluble CP (PFPG2) as a means to interact with bacteria and form aggregates through electrostatic interactions. The authors showed that PFP-G2 coated with bacteria aggregates not only stimulated the bacteria QS system but also prolonged the duration of QS signal molecules.159 In another approach, Tuncagil and colleagues reported on the use of graphite electrodes coated with the CP poly(​1-(4-​nitro​pheny​l)-2,​5-di(​2-thi​enyl)​-I H-pyrrole) [SNS(NO(2))])160 for microbial sensing. A suspension of the two bacteria Gluconobacter oxydans and Pseudomonas fluorescens was spread on the electrodes after the electropolymerization step and the obtained devices were wrapped in a dialysis tube to prevent leakage of bacterial cells. The two electrodes displayed specificity for the microorganism type adsorbed to their surface, as evidenced by the recorded current densities following oxygen consumption at −0.7 V as a result of the metabolic activity. In an interesting approach, Lee et al. showed the first example of a PDA-based system for both sensing and killing resistant bacteria. Contact of imidazolium and imidazole-derived PDA with various bacterial strains, including methicillin-resistant S. aureus and extended-spectrum β-lactamase-producing E. coli, resulted in a distinct blue-to-red colorimetric change in the solution followed by rapid disruption of the bacterial membrane. The antibacterial activity of the PDA solution was attributed to electrostatic interactions between the negatively charged bacterial surface and the positively charged polymers, highlighting the great potential of such conjugated polymeric systems as both a probe of bacteria and potentially a bactericidal agent.16

23.5.5 Cells A key characteristic of the OECT is its ability to act as an ion-to-electron converter that can be operated at low voltages, thus being compatible with biological systems and aqueous environments. In the last decade, PEDOT:PSS-based OECTs have been used at the interface with both electrogenic and nonelectrogenic cells (i.e., barrier cells vs cardiac cells) given its reported biocompatibility.37,161–165 In the case of electrically inactive (non-electrogenic) cells, electronic measurements can be used to estimate surface coverage and cell viability as well as their differentiation. This technique is based on the electrical impedance spectroscopy developed by Giaever and Keese.166 This technique has also been adapted to evaluate the integrity of a specific type of tissue (composed of certain epithelial and

Conducting and Conjugated Polymers for Biosensing Applications


endothelial cell types) that in vivo serve as physical barriers by tightly controlling ionic flux. Ion transport in between cells (paracellular ion flux) is regulated by protein structures known as tight junctions, the state of their functionality providing information about barrier tissue and indicative of certain disease states. Thus, by measuring the electrical resistance (transepithelial/endothelial resistance) of a specific cell layer upon its treatment with specific drugs/toxins, one can derive knowledge on the effect of the applied compounds on the investigated barrier tissues. Interestingly, these kinds of measurements have significant advantages over traditional assays, such as a lactate dehydrogenase assay used to evaluate toxicity, thanks to improved temporal resolution, dynamic measurements and enhanced sensitivity. In 2010, Lin and collaborators suggested the idea of using the OECT as a cell-based biosensor.167 In their work, the presence of a cell layer of either human esophageal squamous epithelial cancer cell lines (KYSE30) or fibroblast cell lines (HFF1) on top of PEDOT:PSS OECTs was investigated for its effect on the transistor functioning and characteristics. Indeed, cell monolayers acted as a barrier to the ionic flux influencing the number of ions reaching the OECT channel. In turn, the channel current served as a measure of the barrier function of the cells. The use of the OECT as a direct measure of ionic flux was proposed by Yao and collaborators.170 The authors anticipated the use of a cation-sensitive polymer such as PEDOT:PSS for recording transepithelial ion transport of human airway epithelial cells. A human airway epithelial cell model, Calu-3 cell line, was used to study transepithelial current dynamics on the OECT device. The specificity of the response by the OECT was investigated by evaluating its behavior when a molecule such as forskolin was used in the study, known to activate the membrane transporter specific for chloride. Cells responded to this event by transporting sodium ions to the apical area, thereby reducing the sodium concentration in the intercellular space of the gap region at the basolateral side facing the PEDOT:PSS channel. This event translated into negative effective gate voltage changes. In a similar approach, our group adopted a human colon cancer cell line, the Caco-2 cell line, to develop a model of the gastrointestinal epithelium (Figure 23.11a).168 Cells were grown on permeable Transwell filters until fully differentiated. These were then integrated with a PEDOT:PSS-based OECT by means of a top, external Ag/AgCl gate electrode. The authors showed for the first time the possibility of monitoring cell layer integrity by interfacing it with an OECT. Using this approach, the authors were able to assess the effect of two cytotoxic compounds (ethanol and H2O2) on Caco-2 cells in vitro using the PEDOT:PSS-based OECT. We further adapted the coupling of the Transwell culturing system with OECT monitoring to investigate the effects of the enteric pathogen Salmonella typhimurium on gastrointestinal epithelia.169 In order to overcome the drawbacks of top gate configuration, related to device fabrication and optical imaging, we went on to develop a device configuration in which a coplanar gate electrode was introduced for the simultaneous recording of the electrical behavior and optical imaging of the established barrier tissue.162,164 The ability of the transistors to detect early and subtle changes in the transepithelial ion flow was demonstrated with the development of a highly sensitive biosensor of epithelial integrity. Interestingly, the device was able to monitor stages in cell layer coverage and differentiation (Figure 23.11b).170 The use of OECT for impedance monitoring of MDCK-I (Madin–Darby canine kidney) cells was further investigated in our group by Rivnay et al.161 MDCK-I cells form well-characterized epithelial-type monolayers when cultured on planar surfaces. These monolayers are considered “tight” as they have a relatively high resistance to ion flow through the paracellular pathway, with an additional cleft resistance thanks to the adhesion complexes formed on the substrate. The authors established a new protocol for OECT operation for cell barrier characterization by combining simultaneous measurements of both drain and gate currents. This optimized device was applied in a proof-of-concept experiment for monitoring the disruption of functional MDCK-I barrier tissue grown on the planar OECT by using the lytic enzyme trypsin. A strategy to detect cancer cells based on the functionalization of thiophene-based CPs CP with boronic acid molecules was developed by Dervisevic and collaborators.173 Boronic acid is able to bind to sialic acid molecules, overexpressed on the surface of cancer cells. The authors developed a biosensor based on a film of boronic acid–functionalized thiophene that was obtained by electropolymerization on the surface of pencil graphite electrodes (PGE). The binding was characterized by electrochemical impedance measurements, and the sensors were shown to be able to detect in the range 10–106 cells/mL. In this work, a solution


Conjugated Polymers

FIGURE 23.11  Examples of cell-based sensing using transistor technology for both electrogenic and non-electrogenic cells. (Reproduced from (a)168, (b-d)170–172.)

of K3[Fe(CN)6]/K4[Fe(CN)6] (1:1) was used as a redox probe. After having immersed the functionalized electrode in the solution containing human caucasian gastric adenocarcinoma (AGS) cancer cells, it was then rinsed and moved into a solution of the redox mediator to run the actual electrochemical impedance measurement. The developed electrode displayed high sensitivity toward AGS cells in comparison with human embryonic kidney 293 (HEK293) normal cells and bone marrow mesenchymal stem cells (BM-hMSCs). CP devices have also been successfully used for monitoring electrogenic cells.174 Rigid and flexible OECT arrays were successfully implemented for monitoring cardiac cell activity. The measured action potentials exhibited excellent signal to noise ratios, typically greater than 4 (Figure 23.11c).172 More recently, an OECT array was developed and used for electrophysiological recording by Hempel and collaborators.165 The major challenge for OECTs is that their speed should overlap with that of cellular action potentials (2–5 ms). A reliable method to produce OECT arrays with high reproducibility was developed and the transistors were shown to be highly stable under cell culture conditions with the net gm experiencing an average increase of 4.7%. The OECT arrays were used for HL-1 cardiomyocytes monitoring, and spontaneous signals were recorded after 3 days of cell culture. Spanu et al. reported on a novel type of polymeric thin-film transistor called an organic charge–modulated FET (OCMFET), as a flexible, transparent, referenceless transducer of extracellular electrical activity. The referenceless, low-cost device with high-spatial and high-temporal resolution was able to record signals from primary cultures of cardiocytes as well as estimate the propagation speed of the electrical signal (Figure 23.11d).171

Conducting and Conjugated Polymers for Biosensing Applications


When it comes to the use of cells as the biorecognition element in biosensors, an interesting approach was developed in laboratories led by Weiss and Penner,175–178 where the use of the M13 filamentous bacteriophage to induce nanostructuring in the conductive polymer poly(3,4-ethylenedioxythiophene)177 was established. The virus particles were incorporated into the polymeric backbone of PEDOT during electropolymerization using lithographically patterned nanowires as electrodes. Donovan and collaborators further expanded upon this technology by developing an impedentiometric sensor for the detection of the M13 bacteriophage.178 Their paper showed the possibility of reaching high sensitivity (6 nM) in the specific detection of the bacteriophage by using a PEDOT film embedding M13 bacteriophages obtained by electropolymerization on gold substrates as the working electrode. By running electrochemical impedance measurements, the authors were able to detect the binding of specific antibodies to the targeted microorganisms. This platform was further exploited by Arter et al. and Mohan et al. for the detection of the prostate-specific membrane antigen (PSMA).175,179 In prostate tissues, the levels of this molecule are proportional to the aggressiveness of the tumor growth. In healthy patients, PSMA levels are as low as 0.5 nM, whereas in patients with prostate cancer it rises to 6 nM. In these papers, the M13 bacteriophage was used to express a synthetic polymer that specifically recognizes and binds to PSMA on its surface. Arter et al. were able to detect levels of PSMA as low as 56 nM by monitoring changes in the electrical resistance of the developed virus–PEDOT nanowires.179 A further improvement in the detection limits was brought about by the work done by Mohan and collaborators.175 The authors reported on a combined approach in which the bacteriophage to be co-electropolymerized with PEDOT to form a film is first genetically modified to express the PSMA-binding molecule and then “functionalized” with a synthetic peptide that wraps around the bacteriophage. This combined binding event, by inducing an apparent higher affinity for the target PSMA molecules, leads to increased sensor sensitivity with detection levels as low as 100 pM, revealed by an increase in the resistance of the electrochemical impedance of the virus-PEDOT film upon binding of PSMA. This technology was also exploited for the development of a miniaturized, portable biosensor for the detection of human serum albumin (HSA).180 Indeed, by selecting a phage displaying a peptide specifically binding to HSA in a mega random peptide library, they were able to selectively detect the biding of HSA molecules. The selected phage was then mixed with EDOT monomers and used to coat the sensor electrodes with a film of PEDOT-phage as previously reported. Impedance measurements for HSA detection were run both in phosphate-buffered saline (PBS) and synthetic urine, without the need for either reference or counter external electrodes. In synthetic urine, the sensitivity of the electrode decreased and this was explained by a de-doping effect exerted by the two amine groups in urea. The sensor was shown to be useful in the physiologically relevant range for HSA (100 nM−5 µM), going beyond state-of-the-art technology such as dipstick tests being able to detect only macroalbuminuria (HAS levels above 200 mg/ mL corresponding to 3 µM), a stage at which the disease has already progressed through kidney failure.

23.5.6 Toward More Biomimetic Systems for Biosensing The integration of devices with living systems such as cells is a very promising approach in biosensing and paves the way for complex organ-on-chip platforms. The challenge, however, lies in creating a fully biomimetic environment for living cells integration and subsequent tissue formation. Along these lines, we developed an in vitro, in-line, sensing platform for monitoring cells, using OECT technology and integrated with µ-fluidics that allowed for media flow with physiological parameters. The physiologically relevant shear stress induced was important for physiologically relevant cell differentiation. Multiparametric cell sensing was shown by combining the OECT-based detection of glucose downstream of the impedance monitoring microfluidic device. Moreover, this platform was used for the assessment of wound healing of the fully confluent cell layer (Figure 23.12a).181 In a very recent approach, Pas et al. showed an enhanced electrophysiological recording yield by varying cell densities on PEDOTPSS​-coat​ed microelectrode arrays (MEAs). The high cell densities (900 cells/mm2) of primary cortical cells resulted in 3D clusters known as neurospheres significantly increasing the single-unit activity


Conjugated Polymers

FIGURE 23.12  The trend in in vitro bioelectronics systems. From (a) 2D, flat biology181 to (b) more physiologically relevant biological systems such as spheroids or cell clusters cultured on planar microelectronic transducers182 to (c) more complex 3D biological systems, fully integrated with adapted electronic elements (i.e., scaffolds) that serve the dual role of hosting and monitoring biological systems.183

recordings as well as the overall yield. This sensitive and biomimetic MEA system bearing more in vivo– like 3D cultured systems paves the way for studying pharmacology-based effects on neural networks (Figure 23.12b).182 An alternative approach was reported by Inal and collaborators who developed an electroactive 3D scaffold based on the conducting polymer PEDOT:PSS for supporting and monitoring cell proliferation.183 In this case, instead of a continuous 2D film, a 3D macroporous matrix was developed displaying pores with diameters compatible with cell infiltration and colonization. Media exchange was enhanced by using a perfusion system. Scaffold colonization and cell proliferation were assessed by running electrochemical impedance measurements (Figure 23.12c). In this case, the conducting polymer performs the dual function of housing cells within a 3D biomimetic “soft” electrode allowing at the same time for monitoring cells proliferation.

23.6 Perspectives Biosensing using CPs has seen enormous growth over the last decade. Since the initial CP biosensing applications more than 20 years ago, increasing sophistication in materials use, increased complexity in modality and tailormade applications have meant that CP biosensing devices are now competing and even outstripping the traditional materials used in biosensing devices. Our understanding of traditional healthcare diagnosis and related tools has started to evolve through an easy-to-use and decentralized diagnosis perspective that offers concepts and devices including miniaturized, wearable and implantable biosensors. However, achieving these paradigm shifts requires significant progress and research of new materials, interfaces, circuit designs, data processing and business models. The next decades will see increased integration of purpose-designed materials in biosensing devices. The numerous advantages of CPs used in biosensing applications illustrated in this chapter will be further exploited to develop future generations of biosensors with increased sensitivity, specificity and user-friendly operation, and open doors to as yet unexplored applications thanks to the “smart” nature of these materials.


Type of Biosensor

S1 nuclease and hydroxyl radicals

Streptavidin Bacterial strains (MRSA and ESBL-EC)

DNA of hepatitis B virus (HBV)

Oligonucleotides (ODNs) (modified with dye)


Re-biotin Monomers 1 and 2

ssDNA strands

d Poly[​9,9-b​is(6,​6-(N,​N,N-t​rimet​hylam​moniu​ m)-fl​uoren​e)-2,​7-yle​nevin​ylene​ -co-alt-2,5dicyano-1,4-phenylene)] (PFVCN) and tungsten disulfide (WS2) nanosheets Poly[​5-met​hoxy-​2-(3-​sulfo​poxy)​-1,4-​pheny​ lenev​inyle​ne] (MPS-PPV) Imidazolium and imidazole-derived polydiacetylenes (PDAs) (monomer 1 [PCDA-Im-ethanol] and monomer 2 [PCDA-Im]) Poly(ethylenimine) (PEI) modified upconversion nanoparticles (NH2-UCNPs)

ATP on plasma membrane of HeLa cells Human prion protein PrPC (103–231) DNA

Target DNA (oligonucleotides) Heparin DNA methylation of cancerous cells

Quorum sensing bacteria

Analyte of Interest

Poly(​2,5-b​is[et​hyl-7​-hept​anoat​e]-p-​pheny​ lenev​inyle​ne-al​t-phe​nylen​eviny​lene)​ (PBEH), bound to amino functionalized magnetic beads

Quaternary ammonium salt groups on the side chains DNA aptamers

HPF-Ir4 Fluorescein-labeled dNTP

Several conjugated polymers Poly(​(1,4-​pheny​lene)​-2,7-​[9,9-​bis(6​′-N,N​,N-tr​ imeth​yl ammonium)-hexyl fluorene] dibromide) n Poly(p-phenylene ethynylene terthiophene) backbone and side chain N2 (PPET3-N2) Polypyrrole (PPy), modified by streptavidin

Polyaniline (PANI)

Terminal amino or quaternary ammonium moieties (provide polyvalent interactions with bacteria) Probe DNA (oligonucleotides)

(Bio)recognition Element

Dendronized poly-(fluorene)s with positively charged amine groups (PFP-G2)

Conducting Polymer

TABLE 23.1  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Fluorescence (quenching through FRET)

Fluorescence quenching (turn-off) Colorimetric and fluorescence

(Continued )







Surface plasmon resonance FRET (turn-on) or superquenching (turn-off), depending on the ODN Fluorescence (quenching through FRET)


185 24




Fluorescence turn-on

Fluorescence (turn-off fluorescence quenching) Colorimetric Fluorescence resonance energy transfer (FRET)


Sensing Mechanism/ Technique

Conducting and Conjugated Polymers for Biosensing Applications 723

Poly (9,9-bis(6′-N,N,N-triethylammonium) hexyl) fluorenylene phenylene dibromide] (PFP)

Polydiacetylene vesicles (PDA) Poly{​[9,9-​bis(6​′-(N,​N,N-d​iethy​lmeth​ylamm​ onium​)hexy​l)-2,​7-flu​oreny​lene ethynylene]alt-co-[​2,5-b​is(3′​-(N,N​,N-di​ethyl​methy​ lammo​nium)​-1′-o​xapro​pyl)-​1,4-p​henyl​ene]}​ tetraiodide (PFEP) Poly (9,9-bis (6′-N,N,N-trimethylammonium) hexyl) fluorenylene phenylene, with DFPPM-Ir-bpy-CO2H)PF6 Ir(III) complex

Poly(fluorenylene phenylene) (PFP)

Poly{​[9,9-​bis(6​′-(N,​N,N-d​iethy​lmeth​ylamm​ onium​)hexy​l)-2,​7-flu​oreny​lene ethynylene]alt-co-[2,5​-bis(​3′-(N​,N,N-​dieth​ylmet​hylam​ moniu​m)-1′​-oxap​ropyl​)-1,4​-phen​ylene​] tetraiodide} (PFEP) Poly(​9,9-b​is(6′​-N,N,​N-tri​methy​lammo​nium)​ hexyl​)fluo​rine phenylene (PFP) Poly(2,5-bis(3- sulfo​natop​ropox​y)-1,​4-phe​nylet​ hynyl​eneal​t-1,4​-poly​(phen​ylene​ ethynylene)) (PPESO3)

Conducting Polymer

Adenosine deaminase (ADA) Catecholamine (catechol, dopamine DA, adrenaline AD, and norepinephrine NE) α-Glucosidase inhibitors

Adenosine–aptamer complex

Fluorescence Fluorescence (quenching through FRET)

Fluorescence (amplifying FRET signal of the complex CCP-Ir(III) and aptamer) Fluorescence (quenching through FRET)

Pathogenic bacterial rRNA CD44 (cancer-mediated protein)




Fluorescence (turn-on/ turn-off)

Fluorescence (amplifying through FRET) Fluorescence (quenching –turn-off)

Fluorescence (quenching through FRET)

Sensing Mechanism/ Technique

The fluorescence of cationic poly(fluorenylene phenylene) (PFP) was quenched in the presence of para-nitrophenyl-α-dglucopyranoside and α-glucosidase, and turned on upon addition of AGIs DNA probe Fluoresceinamine-hyaluronan (FA-HA). HA is an anionic natural glycosaminoglycan that can specifically bind to the overexpressed CD44 on various kinds of cancer cells Aptamer



Analyte of Interest


(Bio)recognition Element

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor

(Continued )



194 195






724 Conjugated Polymers


Streptavidin (SA)

AChE and ChOx

Biotin-linked DNA



Cholesterol oxidase (ChOx)

Probe DNA (ssDNA)




Fluorescence (ratiometric fluorescent response)

Fluorescence (quenching through FRET)

Ca2+-induced conformation changes of calmodulin Choline and acetylcholine

Fluorescence quenching Fluorescence quenching

Fluorescence quenching

Fluorescence (amplification through target-mediated fluorescence resonance energy transfer [TMFRET]) Fluorescence (quenching through FRET)

Sensing Mechanism/ Technique

Bisphenol A (BPA) Hyaluronidase (HAase)


Phenylboronic acid (PBA) (tags on the nanoparticle surface that act as recognition elements for dopamine) HRP Hyaluronan (HA) and anticancer drug doxorubicin (Dox)

PPESO3 Poly{​[9,9-​bis(6​′-(N,​N,N-d​iethy​lmeth​ylamm​ onium​)hexy​l)-2,​7-flu​oreny​lene ethyn​ylene​ ]-alt​-co-[​2,5-b​is(3′​-(N,N​,N-di​ethyl​methy​ lammo​nium)​-1′-o​xapro​pyl)-​1,4-p​henyl​ene]}​ tetraiodide (PFEP) Poly[​(9,9-​bis(6​′-N,N​,N-tr​imeth​ylamm​onium​) hexy​l)-fl​uoren​ylene​ phenylene dibromide], with graphene oxide Poly(fluorene-co-phenylene) derivative (PFP-FB) modified with boronate-protected fluorescein (peroxyfluor-1) via PEG linker Poly(​9,9-b​is(6′​-N,N,​N-tri​methy​lammo​nium)​ hexyl​)-flu​oreny​lene phenylene dibromide (PFP), with GO with TPSMLD system Polymer of (Z)-4-(4-(9H-carbazol-9-yl) benzylidene)-2-(4-nitrophenyl) oxazol5(4H)-one (CBNP) Polyaniline (PANI) mixed with oxidized graphene


Glucose oxidase (GOx)

Poly (9,9-bis (6′-N,N,N-trimethylammonium) hexyl) fluorenylene phenylene (PF), with 3-mercaptopropionic acid (MPA)-capped CdTe/CdS QDs PFPBA nanoparticles

Sequence-specific DNA-binding proteins

Analyte of Interest

Ds DNA probe (fluorophore labeled, bearing a nuclear factor-kappaB [NF-κB]-binding site)

(Bio)recognition Element

Poly{​9,9-b​is[6′​-(N,N​-diet​hylam​ino)h​exyl]​ -2,7-​fluor​enyle​neeth​ynyle​ne}-a​lt-{2​,5-bi​s [3′-(​ N,N-d​iethy​lamin​o)-1′​-oxop​ropyl​]-1,4​-phen​ ylene​}

Conducting Polymer

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor

(Continued )






201 202





Conducting and Conjugated Polymers for Biosensing Applications 725

Poly(3,4-ethylenedioxythiophene) (PEDOT), with bovine serum albumin (BSA) and gold nanoparticles (Au NPs) Overoxidized polypyrrole (PPy), with graphene (PPyox/GR)

[2,2:5,2-terthiophene-3-(p-benzoic acid)] (pTTBA), with gold nanoparticles (Au NPs)

Vinyl-substituted polyaniline (VS-PANI), with polyacrylic acid (PAA) hydrogel, with reduced grapheme oxide (rGO) and lutetium phthalocyanine (LuPc2) Poly(3,4-ethylenedioxythiophene) (PEDOT), with NBR mats, grafted with polyacrylic acid Poly (pyrrole propionic acid) (pPPA), with carbon nanotubes (CNTs) Poly(​5-hyd​roxy-​1,4-n​aphth​oquin​one-c​o-hyd​ roxy-​2-thi​oacet​ic acid-1,4-naphthoquinone) Pyrrole (Py), functionalized with 3-(N-hydroxyphthalimidyl ester) (PyNHP) and 1-(ph​thali​midyl​butan​oate)​-1'-(​N-(3-​butyl​ pyrro​le)bu​tanam​ide) ferrocene (PyFcNHP) and PAMAM G4 dendrimers Polyaniline (PANI)

Poly-pyrrole and poly(3,4ethylenedioxythiophene), (PPy-PEDOT) mixed with silver (Ag) Poly-pyrrole (PPy), with gold nanoparticles

Conducting Polymer

Adenine, guanine

Linear sweep voltammetry

Cyclic voltammetry


Hypoxia inducible factor 1 alpha (HIF1α) in cancerous cells

Antibody (anti-HIF1α) and a nano-bioconjugate with hydrazine and a secondary antibody of HIF1α (sec-Ab2) attached to Au NPs Horseradish peroxidase (HRP)

Differential pulse voltammetry

Electrochemical impedance spectroscopy Differential pulse voltammetry Square wave voltammetry

Electrochemical impedance spectroscopy Amperometry

DNA (non-Hodgkin’s lymphoma gene) Hepatitis B surface antigen (HBsAg) Antibody (anti-ovalbumin) Human prion protein PrPC (103–231)


Cyclic voltammetry and electrochemical impedance spectroscopy Amperometry

Sensing Mechanism/ Technique

Protein psoriasin (S100A7)

Antibodies to protein psoriasin


Alkaline phosphatase (ALP)conjugated secondary antibodies Antigen (ovalbumin)

Probe DNA (ssDNA)

S. aureus ss-DNA

Probe DNA, hybridized with horseradish peroxidase Glucose oxidase (GOx) Glucose

DNA hybridization

Analyte of Interest

6-Mercapto-1-hexhane (HS-ssDNA)

(Bio)recognition Element

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor

(Continued )













726 Conjugated Polymers

Target DNA (sequence for bacterium Bacillus cereus) Target DNA for hepatitis C virus Target DNA (ssDNA strand) for HIV virus

DNA probe DNA probe


Polypyrrole (PPy), with nitrilotriacetic acid

Ricin toxin chain A


DNA probe


Glucose oxidase (GOx)



Cholesterol oxidase (ChOx)

Laccase enzyme


Polyphenol oxidase (PPO)




Alcohol oxidase


Electrochemical impedance spectroscopy Differential pulse voltammetry (DPV) Electrochemical impedance spectroscopy


Electrochemical impedance spectroscopy Amperometry





Double-stranded DNA (dsDNA) Co(phen)33+ Enzymes

Cyclic voltammetry/ differential pulse voltammetry Amperometry

Catechol DNA damage

Chitosan-glucose oxidase (CH-GOx)

Sensing Mechanism/ Technique Electrochemical impedance spectroscopy Amperometry

Analyte of Interest DNA detection (applied for GMO) Glucose


(Bio)recognition Element

bis-EDOT, functionalized with N-Boc-L-lysine and ferrocene Polythiophene (2,5-dibromothiophenes), with PEG Polypyrrole (PPy)

Poly(3,4-ethylene dioxythiophene):poly(styren esulfonate) (PEDOT:PSS) Poly(1,5-diaminonaphthalne) (poly-DAN), with graphene oxide (GO) and gold nanoparticles (GO/Au NPs) Poly(3,4-ethylene dioxythiophene), with graphene oxide (GO) and Fe2O3 nanoparticles Polypyrrole (PPy), with sodium dodecylbenzene sulphonate (DBS) Poly(3,4-ethylene dioxythiophene), with nanoporous gold Polypyrrole (PPy)

Polypyrrole (PPy), with multiwalled carbon nanotubes (MWCNTs) Polypyrrole (PPy), with Nafion (Nf) and multiwalled carbon nanotubes (fMWCNTs) Polydiphenylamine-4-sulfonic acid (PDPASA)

Conducting Polymer

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor

(Continued )
















Conducting and Conjugated Polymers for Biosensing Applications 727


Polyaniline (PANI), with multiwalled carbon nanotubes and chitosan Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​ly(st​ yrene​sulfonate) (PEDOT:PSS)


DNA probe (ssDNA)

Polypyrrole (PPy), with poly(ε-caprolactone) (PCL) Poly(3,4-ethylenedioxythiophene)

(Lactate oxidase/chitosan-ferrocene)

DNA probe (thiol-modified ssDNA of the Arowana fish)

Polypyrrole (PPy), with gold nanoparticles and kappa-carrageenan

Immunoglobulin G (IgG) and myoglobin (Myo) Chlorpyrifos

DNA probe

Monoclonal antibodies (mAbs) of IgG or Myo Anti-chlorpyrifos antibody

Target oligonucleotides

DNA probe

PPy–pyrrolylacrylic acid (PAA) copolymer, doped with either PSS or LiClO4 Polypyrrole (PPy)

Target DNA

Target DNA (for the classification of the Arowana fish gender) Target DNA

Biotin/avidin/digoxigenin-labelled DNA probe

Polyaniline (PANI), with gold nanoparticles

DNA probe

Target DNA

DNA probe

Polyaniline (PANI), with graphene nanosheets

Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​lysty​rene sulfonate (PEDOT:PSS), with porous anodic aluminum oxide (AAO) Polyaniline (PANI)

Target DNA (biotinylated and coupled with streptavidin-alkaline phosphatase conjugate enzymes) Target DNA (and single-nucleotide polymorphisms [SNPs]) Anti-digoxigenin antibody conjugated to horseradish peroxidase (HRP) Target DNA

Analyte of Interest

DNA probe (17-mer thiol-tethered DNA and spacer thiol: 6-mercapto-1hexanol [MCH])

(Bio)recognition Element

Polyaniline (PANI), with gold nanoparticles

Conducting Polymer

Transistor-based electrochemical



Electrochemical impedance spectroscopy Amperometry (and fluorescence) Transistor-based electrochemical

Electrochemical impedance spectroscopy Electrochemical impedance spectroscopy Differential pulse voltammetry (DPV)

Cyclic voltammetry

Electrochemical impedance spectroscopy

Differential pulse voltammetry (DPV)

Sensing Mechanism/ Technique

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor

(Continued )














728 Conjugated Polymers


Detection of the effect of human cytomegalovirus (CMV) on cells Detect the presence of liposomes and liposomebased nanoparticles in solution Detection of antibiotics (ampicillin or kanamycin A) Uric acid, cholesterol, and triglyceride

Glucose oxidase (GOx)

Enzymes uricase (UOx), cholesterol esterase/cholesterol oxidase (ChEt/ ChOx), and lipase/glycerol kinase/ glycerol-3-phosphate oxidase (LIP/ GK/GPO)

Aptamer probes (with affinity to ampicillin or kanamycin A)


Glucose oxidase (GOx)

Poly(3,4-ethylenedioxythiophene) (PEDOT:TsO) and the hydroxymethyl derivative PEDOT-OH:TsO PANI hydrogel/PtNPs hybrid electrodes

Salmonella typhimurium

Human acute leukemia T cells Cardiomyocytes

Analyte of Interest

Drug/lipid probe (raltitrexed/ phosphatidylcholine) –

(Bio)recognition Element

Poly(3,4-ethylenedioxythiophene) doped with poly(styrene sulfonate), PEDOT:PSS, with ethylene glycol and DBSA

Poly-(3,4-diamiono-2,2:5,2-terthiophene) (DATT) Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​ly(st​ yrene​sulfonate) (PEDOT:PSS), with ethylene glycol Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​ly(st​ yrene​sulfonate) (PEDOT:PSS), with ethylene glycol and DBSA Poly(​9,9-d​i-(2-​ethyl​hexyl​)-flu​oreny​l-2,7​-diyl​) end capped with 2,5-diphenyl-1,2,4oxadiazole (PFLO) Poly[​9,9-d​i-(2-​ethyl​hexyl​)-flu​oreny​l-2,7​-diyl​] end capped with N,N-bis(4-methylphenyl)-4aniline (PFLA), with multiwalled carbon nanotubes (MWCNTs) and zinc phthalocyanine (ZnPc) Poly(3,4-ethylenedioxythiophene) doped with tosylate (PEDOT:TsO)

Conducting Polymer


Electrochemical impedance spectroscopy

Transistor-based electrochemical

Electrochemical impedance spectroscopy



(Continued )









Transistor-based electrochemical Transistor-based electrochemical




Sensing Mechanism/ Technique

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor

Conducting and Conjugated Polymers for Biosensing Applications 729

Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​ly(st​ yrene​sulfonate) (PEDOT:PSS)

Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​ly(st​ yrene​sulfonate) (PEDOT:PSS)–poly(vinyl alcohol) (PVA) and Pt nanoparticles (NPs) P3HT

Poly(EDOT-co-EDOTPC) film, dopant NaClO4 Poly(3,4-ethylene dioxythiophene):poly(styrene sulfonate) (PEDOT:PSS), with gate electrode modified with GO and PANI/Nafion-graphene bilayer films Poly(3,4-ethylene dioxythiophene):poly(styrene sulfonate) (PEDOT:PSS), with Nafion/nanomaterial (graphene flakes [Gr] or graphene oxide [GO] or single-walled carbon nanotubes [SWNTs]) modified Pt gate electrodes Carboxylated polypyrrole nanoparticles (CPPyNPs) Poly(3-thiophene acetic acid) (P3) on ITO surface Polymers of 2-(2-​(2-(4​H-dit​hieno​[3,2-​b:2′,​ 3′-d]​pyrro​l-4 yl)ethoxy)ethoxy)ethanamine (poly[DTP-alkoxy-NH2]) Poly(3,4-ethylene dioxythiophene):poly(styrene sulfonate) (PEDOT:PSS) PEDOT:PSS

Conducting Polymer


Dopamine (DA)

Glucose, lactate, and cholesterol Glucose and lactate

Glucose oxidase (GOx)

GOx, LOx, ChOx

PEG-SH blocking layer and aptamers that specifically capture ricin GOx cells

Glucose cell viability


Electrochemical impedance spectroscopy Amperometry


Glucose oxidase (GOx) Lactate oxidase (LOx)

Transistor-based field effect


Parathyroid hormone receptor (hPTHR) Antibody–anti-TNFa

Transistor-based electrochemical

Transistor-based field effect

Transistor-based electrochemical Transistor-based electrochemical

Transistor-based electrochemical

Transistor-based electrochemical


Differential pulse voltammetry (DPV) Transistor-based electrochemical

Sensing Mechanism/ Technique

Uricase (UOx)

Analyte of Interest Human C-reactive protein (CRP) Uric acid (UA)

(Bio)recognition Element

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor

(Continued )













730 Conjugated Polymers

Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​ly(st​ yrene​sulfonate) (PEDOT:PSS) P3HT

Poly(​3,4-e​thyle​nedio​xythi​ophen​e):po​ly(st​ yrene​sulfonate) (PEDOT:PSS)

Conducting Polymer

Cell barrier integrity Streptavidin (SA)

Biotinylated phospholipids

Cell barrier function

Analyte of Interest

Human esophageal squamous epithelial cancer cells (KYSE30) fibroblasts (HFF1) Caco-2 cells

(Bio)recognition Element

Transistor-based electrochemical Transistor-based field-effect transistor

Transistor-based electrochemical

Sensing Mechanism/ Technique

TABLE 23.1 (CONTINUED)  Summary of CP Biosensing Devices Describing the CP Type, Biorecognition Element, Analyte and Sensing Mechanism

Type of Biosensor





Conducting and Conjugated Polymers for Biosensing Applications 731


Conjugated Polymers

References 1. Bhalla, N., Jolly, P., Formisano, N. & Estrela, P. Introduction to biosensors. Essays Biochem. 60, 1–8 (2016). 2. Heineman, W. R. & Jensen, W. B. Leland C. Clark Jr. (1918–2005). Biosens. Bioelectron. 21, 1403–1404 (2006). doi:10.1016/j.bios.2005.12.005 3. Clark, L. C. & Lyons, C. Electrode systems for continuous monitoring in cardiovascular surgery. Ann. N. Y. Acad. Sci. 102, 29–45 (2006). 4. Scheller, F. W. et al. Second generation biosensors. Biosens. Bioelectron. 6, 245–253 (1991). 5. Turner, A. P. F. Chemical society reviews biosensors: Sense and sensibility. Chem. Soc. Rev. Chem. Soc. Rev 42, 3175–3196 (2013). 6. Daniels, J. S. & Pourmand, N. Label-free impedance biosensors: Opportunities and challenges. Electroanalysis 19, 1239–1257 (2007). 7. Kergoat, L., Piro, B., Berggren, M., Horowitz, G. & Pham, M.-C. Advances in organic transistorbased biosensors: From organic electrochemical transistors to electrolyte-gated organic fieldeffect transistors. Anal. Bioanal. Chem. 402, 1813–1826 (2012). 8. Sin, M. L., Mach, K. E., Wong, P. K. & Liao, J. C. Advances and challenges in biosensor-based diagnosis of infectious diseases. Expert Rev. Mol. Diagn. 14, 225–244 (2014). 9. Palchetti, I., Hansen, P.-D. & Barcelo ́, D. Past, Present and Future Challenges of Biosensors and Bioanalytical Tools in Analytical Chemistry: A Tribute to Professor Marco Mascini (Amsterdam: Elsevier, 2017). 10. Facchetti, A. π-Conjugated polymers for organic electronics and photovoltaic cell applications. Chem. Mater. 23, 733–758 (2011). 11. Lee, K., Povlich, L. K. & Kim, J. Recent advances in fluorescent and colorimetric conjugated polymer-based biosensors. Analyst 135, 2179 (2010). 12. Korotchenkov, G. S. Chemical Sensors Fundamentals of Sensing Materials. Volume 2 (Momentum Press, New York, 2010). 13. Rochat, S. & Swager, T. M. Conjugated amplifying polymers for optical sensing applications. ACS Appl. Mater. Interfaces 5, 4488–4502 (2013). 14. Alvarez, A., Costa-Fernández, J. M., Pereiro, R., Sanz-Medel, A. & Salinas-Castillo, A. Fluorescent conjugated polymers for chemical and biochemical sensing. Trends Anal. Chem. 30, 1513–1525 (2011). 15. Leclerc, M. Optical and electrochemical transducers based on functionalized conjugated polymers. Adv. Mater. 11, 1491–1498 (1999). 16. Lee, S. et al. Sensing and antibacterial activity of imidazolium-based conjugated polydiacetylenes. Biosens. Bioelectron. 77, 1016–1019 (2016). 17. Swager, T. M. The molecular wire approach to sensory signal amplification. Acc. Chem. Res. 31, 201–207 (1998). 18. Chen, Z. et al. Sensitive conjugated-polymer-based fluorescent ATP probes and their application in cell imaging. ACS Appl. Mater. Interfaces 8 , 3567–3574 (2015). doi:10.1021/acsami.5b06935 19. Zhu, C., Liu, L., Yang, Q., Lv, F. & Wang, S. Water-soluble conjugated polymers for imaging, diagnosis, and therapy. Chem. Rev. 112, 4687–4735 (2012). doi:10.1021/cr200263w 20. Feng, X., Liu, L., Wang, S. & Zhu, D. Water-soluble fluorescent conjugated polymers and their interactions with biomacromolecules for sensitive biosensors. Chem. Soc. Rev. 39, 2411–2419 (2010). 21. Tan, C. et al. Amplified quenching of a conjugated polyelectrolyte by cyanine dyes. J. Am. Chem. Soc. 126, 13685–13694. doi:10.1021/ja046856b 22. Zhou, Q. & Swager, T. M. Fluorescent chemosensors based on energy migration in conjugated polymers: The molecular wire approach to increased sensitivity. J. Am. Chem. Soc. 117, 12593–12602 (1995).

Conducting and Conjugated Polymers for Biosensing Applications


23. Cao, A., Tang, Y. & Liu, Y. Novel fluorescent biosensor for α-glucosidase inhibitor screening based on cationic conjugated polymers. ACS Appl. Mater. Interfaces 4, 3773–3778 (2012). 24. Feng, F., Liu, L. & Wang, S. Fluorescent conjugated polymer-based FRET technique for detection of DNA methylation of cancer cells. Nat. Protoc. 5, 1255–1264 (2010). 25. Aydemir, N., Malmström, J. & Travas-Sejdic, J. Conducting polymer based electrochemical biosensors. Phys. Chem. Chem. Phys. 18, 8264–8277 (2016). 26. Owens, R. M. & Malliaras, G. G. Organic electronics at the interface with biology. MRS Bull. 35, 449–456 (2010). 27. Moon, J.-M., Thapliyal, N., Hussain, K. K., Goyal, R. N. & Shim, Y.-B. Conducting polymer-based electrochemical biosensors for neurotransmitters: A review. Biosens. Bioelectron. 102, 540–552 (2018). 28. Rahman, M., Kumar, P., Park, D.-S. & Shim, Y.-B. Electrochemical sensors based on organic conjugated polymers. Sensors 8, 118–141 (2008). 29. Shrestha, B. K. et al. High-performance glucose biosensor based on chitosan-glucose oxidase immobilized polypyrrole/Nafion/functionalized multi-walled carbon nanotubes bio-nanohybrid film. J. Colloid Interface Sci. 482, 39–47 (2016). 30. Zheng, Q. et al. An electrochemical DNA sensor based on polyaniline/graphene: High sensitivity to DNA sequences in a wide range. Analyst 140, 6660–6670 (2015). 31. White, H. S., Kittlesen, G. P. & Wrighton, M. S. Chemical derivatization of an array of three gold microelectrodes with polypyrrole: Fabrication of a molecule-based transistor. J. Am. Chem. Soc. 106, 5375–5377 (1984). 32. Panzer, M. J. & Frisbie, C. D. Exploiting ionic coupling in electronic devices: Electrolyte-gated organic field-effect transistors. Adv. Mater. 20, 3177–3180 (2008). 33. Casalini, S. et al. Multiscale sensing of antibody–antigen interactions by organic transistors and single-molecule force spectroscopy. ACS Nano 9, 5051–5062 (2015). 34. Berto, M. et al. Biorecognition in organic field effect transistors biosensors: The role of the density of states of the organic semiconductor. Anal. Chem. 88, 12330–12338 (2016). 35. Kim, Z.-S., Lim, S. C., Kim, S. H., Yang, Y. S. & Hwang, D.-H. Biotin-functionalized semiconducting polymer in an organic field effect transistor and application as a biosensor. Sensors 12, 11238–11248 (2012). 36. Rivnay, J. et al. Organic electrochemical transistors. Nat. Rev. Mater. 3, 086 (2018). 37. Strakosas, X., Bongo, M. & Owens, R. M. The organic electrochemical transistor for biological applications. J. Appl. Polym. Sci. 132, 41735 (2015). 38. Hackett, A. J., Malmström, J. & Travas-Sejdic, J. Functionalization of conducting polymers for biointerface applications. Prog. Polym. Sci. 70, 18–33 (2017). 39. Strakosas, X. et al. Catalytically enhanced organic transistors for in vitro toxicology monitoring through hydrogel entrapment of enzymes. J. Appl. Polym. Sci. 134, 44483 (2017). 40. Gerard, M., Chaubey, A. & Malhotra, B. D. Application of conducting polymers to biosensors. Biosens. Bioelectron. 17, 345–359 (2002). 41. Tamiya, E., Karube, I., Hattori, S., Suzuki, M. & Yokoyama, K. Micro glucose using electron mediators immobilized on a polypyrrole-modified electrode. Sensors Actuators 18, 297–307 (1989). 42. Fu, Y., Yuan, R., Chai, Y., Zhou, L. & Zhang, Y. Coupling of a reagentless electrochemical DNA biosensor with conducting polymer film and nanocomposite as matrices for the detection of the HIV DNA sequences. Anal. Lett. 39, 467–482 (2006). 43. Liu, J. et al. Direct imaging of the electrochemical deposition of poly(3,4-ethylenedioxythiophene) by transmission electron microscopy. ACS Macro Lett. 4, 897–900 (2015). 44. Sadik, O. A., Brenda, S., Joasil, P. & Lord, J. Electropolymerized conducting polymers as glucose sensors. J. Chem. Educ. 76, 967 (1999). 45. Sung, W. J. & Bae, Y. H. A glucose oxidase electrode based on electropolymerized conducting polymer with polyanion−enzyme conjugated dopant. Anal. Chem. 72, 2177–2181 (2000). doi:10.1021/AC9908041


Conjugated Polymers

46. Zhao, D., Yuan, D., Sun, D. & Zhou, H.-C. Stabilization of metal–organic frameworks with high surface areas by the incorporation of mesocavities with microwindows. J. Am. Chem. Soc. 131, 9186–9188 (2009). 47. Nien, P.-C., Tung, T.-S. & Ho, K.-C. Amperometric glucose biosensor based on entrapment of glucose oxidase in a poly(3,4-ethylenedioxythiophene) film. Electroanalysis 18, 1408–1415 (2006). 48. Lu, Q. & Li, C. M. One-step co-electropolymerized conducting polymer-protein composite film for direct electrochemistry-based biosensors. Biosens. Bioelectron. 24, 773–778 (2008). 49. Xiao, X., Wang, M., Li, H. & Si, P. One-step fabrication of bio-functionalized nanoporous gold/ poly(3,4-ethylenedioxythiophene) hybrid electrodes for amperometric glucose sensing. Talanta 116, 1054–1059 (2013). 50. Komarova, E., Aldissi, M. & Bogomolova, A. Design of molecularly imprinted conducting polymer protein-sensing films via substrate–dopant binding. Analyst 140, 1099–1106 (2015). 51. Mano, N., Yoo, J. E., Tarver, J., Loo, Y. L. & Heller, A. An electron-conducting cross-linked polyaniline-based redox hydrogel, formed in one step at pH 7.2, wires glucose oxidase. J. Am. Chem. Soc. 129 (22), 7006–7007 (2007). 52. Strakosas, X. et al. A facile biofunctionalisation route for solution processable conducting polymer devices. J. Mater. Chem. B 2, 2537–2545 (2014). 53. Pan, S. & Rothberg, L. Chemical control of electrode functionalization for detection of DNA hybridization by electrochemical impedance spectroscopy. Langmuir 21 (3), 1022–1027 (2005). doi:10.1021/LA048083A 54. Pappa, A.-M. et al. Polyelectrolyte layer-by-layer assembly on organic electrochemical transistors. ACS Appl. Mater. Interfaces 9, 10427–10434 (2017). 55. Li, B. et al. Inkjet printed chemical sensor array based on polythiophene conductive polymers. Sensors Actuators B: Chem. 123, 651–660 (2007). 56. Yemini, M., Reches, M., Gazit, E., & Rishpon, J. Peptide nanotube-modified electrodes for enzyme−biosensor applications. Anal. Chem. 77, 5155–5159 (2005). doi:10.1021/AC050414G 57. Kesik, M. et al. Synthesis and characterization of conducting polymers containing polypeptide and ferrocene side chains as ethanol biosensors. Polym. Chem. 5, 6295–6306 (2014). 58. Akbulut, H. et al. Polythiophene-g-poly(ethylene glycol) with lateral amino groups as a novel matrix for biosensor construction. ACS Appl. Mater. Interfaces 7, 20612–20622 (2015). 59. Yameen, B. et al. A facile avenue to conductive polymer brushes via cyclopentadiene–maleimide Diels–Alder ligation. Chem. Commun. 49, 8623 (2013). 60. Li, L. et al. A nanostructured conductive hydrogels-based biosensor platform for human metabolite detection. Nano Lett. 15, 1146 (2015) doi:10.1021/nl504217p 61. Ohayon, D. et al. Laser patterning of self-assembled monolayers on PEDOT:PSS films for controlled cell adhesion. Adv. Mater. Interfaces 4, 1700191 (2017). 62. Ngoepe, M. et al. Integration of biosensors and drug delivery technologies for early detection and chronic management of illness. Sensors 13, 7680–7713 (2013). 63. Erdmann, V. A. & Barciszewski, J. DNA and RNA Nanobiotechnologies in Medicine: Diagnosis and Treatment of Diseases (Springer, Berlin/Heidelberg, 2013). 64. Lu, C.-H., Willner, B. & Willner, I. DNA anotechnology: From sensing and DNA machines to drug-delivery systems. ACS Nano 7, 8320–8332 (2013). 65. Patolsky, F., Zheng, G. & Lieber, C. M. Nanowire sensors for medicine and the life sciences. Nanomedicine 1, 51–65 (2006). 66. Drummond, T. G., Hill, M. G. & Barton, J. K. Electrochemical DNA sensors. Nat. Biotechnol. 21, 1192–1199 (2003). 67. De Stefano, L. et al. DNA optical detection based on porous silicon technology: From biosensors to biochips. Sensors 7, 214–221 (2007). 68. Chen, J. I. L., Chen, Y. & Ginger, D. S. Plasmonic nanoparticle dimers for optical sensing of DNA in complex media. J. Am. Chem. Soc. 132, 9600–9601 (2010).

Conducting and Conjugated Polymers for Biosensing Applications


69. Zhao, W.-W., Xu, J.-J. & Chen, H.-Y. Photoelectrochemical DNA biosensors. Chem. Rev. 114, 7421– 7441 (2014). 70. Lin, L., Liu, Y., Tang, L. & Li, J. Electrochemical DNA sensor by the assembly of graphene and DNA-conjugated gold nanoparticles with silver enhancement strategy. Analyst 136, 4732 (2011). 71. Cagnin, S. et al. Overview of electrochemical DNA biosensors: New approaches to detect the expression of life. Sensors (Basel) 9, 3122–3148 (2009). 72. Zaffino, R. L., Galan, T., Pardo, W. A., Mir, M. & Samitier, J. Nanoprobes for enhanced electrochemical DNA sensors. Wiley Interdiscip. Rev. Nanomed. Nanobiotechnol. 7, 817–827 (2015). 73. Prabhakar, N., Arora, K., Singh, H. & Malhotra, B. D. Polyaniline-based nucleic acid sensor. J. Phys. Chem. B 112, 4808–4816 (2008). doi:10.1021/JP711853Q 74. Rahman, M., Li, X.-B., Lopa, N., Ahn, S. & Lee, J.-J. Electrochemical DNA hybridization sensors based on conducting polymers. Sensors 15, 3801–3829 (2015). 75. Raymond, F. R. et al. Detection of target DNA using fluorescent cationic polymer and peptide nucleic acid probes on solid support. BMC Biotechnol. 5, 10 (2005). 76. Zhang, L. et al. A conjugated polymer-based electrochemical DNA sensor: Design and application of a multi-functional and water-soluble conjugated polymer. Macromol. Rapid Commun. 29, 1489–1494 (2008). 77. Mangalum, A., Rahman, M. & Norton, M. L. Site-specific immobilization of single-walled carbon nanotubes onto single and one-dimensional DNA origami. J. Am. Chem. Soc. 135, 2451–2454 (2013). 78. Peng, H., Soeller, C. & Travas-Sejdic, J. Novel conducting polymers for DNA sensing. Macromolecules 40, 909–914 (2007). 79. Spain, E., Keyes, T. E. & Forster, R. J. Polypyrrole–gold nanoparticle composites for highly sensitive DNA detection. Electrochim. Acta 109, 102–109 (2013). 80. Shakiba, A. et al. DNA loading and release using custom-tailored poly(l-lysine) surfaces. ACS Appl. Mater. Interfaces 9, 23370–23378 (2017). 81. Wegmann, F. et al. Polyethyleneimine is a potent mucosal adjuvant for viral glycoprotein antigens. Nat. Biotechnol. 30, 883–888 (2012). 82. Velusamy, V. et al. Comparison between DNA immobilization techniques on a redox polymer matrix. Am. J. Anal. Chem. 2, 392–400 (2011). 83. Steel, A. B., Levicky, R. L., Herne, T. M. & Tarlov, M. J. Immobilization of nucleic acids at solid surfaces: Effect of oligonucleotide length on layer assembly. Biophys. J. 79, 975–981 (2000). 84. Mazloum-Ardakani, M., Rajabzadeh, N., Benvidi, A. & Heidari, M. M. Sex determination based on amelogenin DNA by modified electrode with gold nanoparticle. Anal. Biochem. 443, 132–138 (2013). 85. Peng, H. et al. Label-free electrochemical DNA sensor based on functionalised conducting copolymer. Biosens. Bioelectron. 20, 1821–1828 (2005). 86. Galán, T. et al. Label-free electrochemical DNA sensor using ‘click’-functionalized PEDOT electrodes. Biosens. Bioelectron. 74, 751–756 (2015). 87. Wahab, R. et al. Immobilization of DNA on nano-hydroxyapatite and their interaction with carbon nanotubes. Synth. Metals 159, 238–245 (2009). 88. Dupont-Filliard, A., Billon, M., Livache, T. & Guillerez, S. Biotin/avidin system for the generation of fully renewable DNA sensor based on biotinylated polypyrrole film. Anal. Chim. Acta 515, 271–277 (2004). 89. Baur, J. et al. Label-free femtomolar detection of target DNA by impedimetric DNA sensor based on poly(pyrrole-nitrilotriacetic acid) film. Anal. Chem. 82, 1066–1072 (2010). 90. Michalska, A. Optimizing the analytical performance and construction of ion-selective electrodes with conducting polymer-based ion-to-electron transducers. Anal. Bioanal. Chem. 384, 391–406 (2005).


Conjugated Polymers

91. Bobacka, J. Conducting polymer-based solid-state ion-selective electrodes. Electroanalysis 18, 7–18 (2006). 92. Saberi, R.-S., Shahrokhian, S. & Marrazza, G. Amplified electrochemical DNA sensor based on polyaniline film and gold nanoparticles. Electroanalysis 25, 1373–1380 (2013). 93. Shoaie, N., Forouzandeh, M. & Omidfar, K. Highly sensitive electrochemical biosensor based on polyaniline and gold nanoparticles for DNA detection. IEEE Sens. J. 99, 1–1 (2017). doi:10.1109/ JSEN.2017.2787024 94. Livache, T. et al. Preparation of a DNA matrix via an electrochemically directed copolymerization of pyrrole and oligonucleotides bearing a pyrrole group. Nucl. Acids Res. 22, 2915–2921 (1994). 95. Booth, M. A., Harbison, S. & Travas-Sejdic, J. Development of an electrochemical polypyrrolebased DNA sensor and subsequent studies on the effects of probe and target length on performance. Biosens. Bioelectron. 28, 362–367 (2011). 96. Mai, A. T., Duc, T., Chu, T. & Nguyen, H. Highly sensitive DNA sensor based on polypyrrole nanowire. Appl. Surface Sci. 309, 285–289 (2014). 97. Esmaeili, C. et al. A DNA biosensor based on kappa-carrageenan-polypyrrole-gold nanoparticles composite for gender determination of Arowana fish (Scleropages formosus). Sensors Actuators B Chem. 242, 616–624 (2017). 98. Guler, Z., Erkoc, P. & Sarac, A. S. Electrochemical impedance spectroscopic study of singlestranded DNA-immobilized electroactive polypyrrole-coated electrospun poly(-caprolactone) nanofibers. Mater. Exp. 5, 269–279 (2015). 99. Luo, S.-C., Xie, H., Chen, N. & Yu, H. Trinity DNA detection platform by ultrasmooth and functionalized PEDOT biointerfaces. ACS Appl. Mater. Interfaces 1, 1414–1419 (2009). 100. Radhakrishnan, S. et al. Polyp​yrrol​e–pol​y(3,4​-ethy​lened​ioxyt​hioph​ene)–​Ag (PPy–PEDOT–Ag) nanocomposite films for label-free electrochemical DNA sensing. Biosens. Bioelectron. 47, 133–140 (2013). 101. Tao, W. et al. A sensitive DNA sensor based on an organic electrochemical transistor using a peptide nucleic acid-modified nanoporous gold gate electrode. RSC Adv. 7, 52118–52124 (2017). 102. Appa, A. et al. Polyelectrolyte Layer-by-Layer Assembly on Organic Electrochemical Transistors. (2017). doi:10.1021/acsami.6b15522 103. Lassalle, N., Vieil, E., Correia, J. P. & Abrantes, L. M. Study of DNA hybridization on polypyrrole grafted with oligonucleotides by photocurrent spectroscopy. Biosens. Bioelectron. 16, 295–303 (2001). 104. Huang, Y.-Q. et al. Tuning backbones and side-chains of cationic conjugated polymers for optical signal amplification of fluorescent DNA detection. Biosens. Bioelectron. 24, 2973–2978 (2009). 105. Kergoat, L. et al. DNA detection with a water-gated organic field-effect transistor. Org. Electron. 13, 1–6 (2012). 106. Lin, P., Luo, X., Hsing, I.-M. & Yan, F. Organic electrochemical transistors integrated in flexible microfluidic systems and used for label-free DNA sensing. Adv. Mater. 23, 4035–4040 (2011). 107. Yoo, E.-H. & Lee, S.-Y. Glucose biosensors: An overview of use in clinical practice. Sensors 10, 4558–4576 (2010). 108. Pappa, A.-M. et al. Organic electronics for point-of-care metabolite monitoring. Trends Biotechnol. 36, 45–59 (2018). 109. Inal, S., Malliaras, G. G. & Rivnay, J. Benchmarking organic mixed conductors for transistors. Nat. Commun. 8, 1767 (2017). 110. Foulds, N. C. & Lowe, C. R. Enzyme entrapment in electrically conducting polymers. Immobilisation of glucose oxidase in polypyrrole and its application in amperometric glucose sensors. J. Chem. Soc. Faraday Trans. 1 Phys. Chem. Condens. Phases 82, 1259 (1986). 111. Palomera, N. et al. Zinc oxide nanorods modified indium tin oxide surface for amperometric urea biosensor. J. Nanosci. Nanotechnol. 11, 6683–6689 (2011). 112. Swann, M. J., Bloor, D., Haruyama, T. & Aizawa, M. The role of polypyrrole as charge transfer mediator and immobilization matrix for d-fructose dehydrogenase in a fructose sensor. Biosens. Bioelectron. 12, 1169–1182 (1997).

Conducting and Conjugated Polymers for Biosensing Applications


113. Schuhmann, W., Zimmermann, H., Habermüller, K. & Laurinavicius, V. Electron-transfer pathways between redox enzymes and electrode surfaces: Reagentless biosensors based on thiol-monolayer-bound and polypyrrole-entrapped enzymes. Faraday Discuss. 116, 245–255 (2000). 114. Thompson, B. C., Winther-Jensen, O., Vongsvivut, J., Winther-Jensen, B. & MacFarlane, D. R. Conducting polymer enzyme alloys: Electromaterials exhibiting direct electron transfer. Macromol. Rapid Commun. 31, 1293–1297 (2010). 115. Malhotra, B. et al. Polyaniline-based biosensors. Nanobiosensors Dis. Diagn. 4, 25 (2015). 116. Granot, E., Basnar, B., Cheglakov, Z., Katz, E. & Willner, I. Enhanced bioelectrocatalysis using single-walled carbon nanotubes (SWCNTs)/polyaniline hybrid systems in thin-film and microrod structures associated with electrodes. Electroanalysis 18, 26–34 (2006). 117. Komathi, S., Gopalan, A. I. & Lee, K.-P. Covalently linked silica-multiwall carbon nanotubepolyaniline network: An electroactive matrix for ultrasensitive biosensor. Biosens. Bioelectron. 25, 944–947 (2009). 118. Xian, Y. et al. Glucose biosensor based on Au nanoparticles-conductive polyaniline nanocomposite. Biosens. Bioelectron. 21, 1996–2000 (2006). 119. Wang, J.-Y., Chen, L.-C. & Ho, K.-C. Synthesis of redox polymer nanobeads and nanocomposites for glucose biosensors. ACS Appl. Mater. Interfaces 5, 7852–7861 (2013). 120. Pappa, A.-M. et al. Direct metabolite detection with an n-type accumulation mode organic electrochemical transistor. Sci. Adv., 4, 0911 (2018). 121. Zhang, L. et al. A polyaniline microtube platform for direct electron transfer of glucose oxidase and biosensing applications. J. Mater. Chem. B 3, 1116–1124 (2015). 122. Bartic, C., Campitelli, A., Borghs, S. & Campitelli, A. Field-effect detection of chemical species with hybrid organic/inorganic transistors. Appl. Phys. Lett. 82, 475–477 (2003). 123. Hoa, D. T. et al. A biosensor based on conducting polymers. Anal. Chem. 64, 2645–2646 (1992). 124. Bernards, D. A. et al. Enzymatic sensing with organic electrochemical transistors. J. Mater. Chem. 18, 116–120 (2008). 125. Shim, N. Y. et al. All-plastic electrochemical transistor for glucose sensing using a ferrocene mediator. Sensors (Basel) 9, 9896–9902 (2009). 126. Scheiblin, G. et al. Screen-printed organic electrochemical transistors for metabolite sensing. MRS Commun. 5, 507–511 (2015). 127. Liao, C., Mak, C., Zhang, M., Chan, H. L. W. & Yan, F. Flexible organic electrochemical transistors for highly selective enzyme biosensors and used for saliva testing. Adv. Mater. 27, 676–681 (2015). 128. Bihar, E. et al. A Disposable paper breathalyzer with an alcohol sensing organic electrochemical transistor. Sci. Rep. 6, 27582 (2016). 129. Gualandi, I. et al. Textile organic electrochemical transistors as a platform for wearable biosensors. Sci. Rep. 6, 33637 (2016). 130. Pappa, A.-M. et al. Organic transistor arrays integrated with finger-powered microfluidics for multianalyte saliva testing. Adv. Healthc. Mater. 5, 2295–2302 (2016). 131. Braendlein, M. et al. Lactate detection in tumor cell cultures using organic transistor circuits. Adv. Mater. (2017). doi:10.1002/adma.201605744 132. Jayasena, S. D. Aptamers: An emerging class of molecules that rival antibodies in diagnostics. Clin. Chem. 45, 1628–1650 (1999). 133. Lee, I., Luo, X., Cui, X. T. & Yun, M. Highly sensitive single polyaniline nanowire biosensor for the detection of immunoglobulin G and myoglobin. Biosens. Bioelectron. 26, 3297–302 (2011). 134. Sun, X., Qiao, L. & Wang, X. A novel immunosensor based on Au nanoparticles and polyaniline/ multiwall carbon nanotubes/chitosan nanocomposite film functionalized interface. Nano-Micro Lett. 5, 191–201 (2013). 135. Kanungo, M., Srivastava, D. N., Kumar, A. & Contractor, A. Q. Conductimetric immunosensor based on poly(3,4-ethylenedioxythiophene). Chem. Commun. 7, 680–681 (2002).


Conjugated Polymers

136. Mouffouk, F. & Higgins, S. J. A biotin-functionalised poly(3,4-ethylenedioxythiophene)-coated microelectrode which responds electrochemically to avidin binding. Electrochem. Commun. 8, 15–20 (2006). 137. Suspène, C. et al. Copolythiophene-based water-gated organic field-effect transistors for biosensing. J. Mater. Chem. B 1, 2090 (2013). 138. Palazzo, G. et al. Detection beyond Debye’s length with an electrolyte-gated organic field-effect transistor. Adv. Mater. 27, 911–916 (2015). 139. Casalini, S., Leonardi, F., Cramer, T. & Biscarini, F. Organic field-effect transistor for label-free dopamine sensing. Org. Electron. 14, 156–163 (2013). 140. Mulla, M. Y. et al. Capacitance-modulated transistor detects odorant binding protein chiral interactions. Nat. Commun. 6, 6010 (2015). 141. White, S. P., Dorfman, K. D. & Frisbie, C. D. Operating and sensing mechanism of electrolytegated transistors with floating gates: Building a platform for amplified biodetection. J. Phys. Chem. C 120, 108–117 (2016). 142. White, S. P., Sreevatsan, S., Frisbie, C. D. & Dorfman, K. D. Rapid, selective, label-free aptameric capture and detection of ricin in potable liquids using a printed floating gate transistor. ACS Sensors 1, 1213–1216 (2016). 143. Tyler McQuade, D., Anthony E. Pullen, & Swager, T. M. Conjugated polymer-based chemical sensors. Chem. Rev. 100, 2537–2574 (2000). 144. Fu, Y. et al. Highly sensitive detection of protein biomarkers with organic electrochemical transistors. Adv. Mater. 29, 1703787 (2017). 145. Noh, H.-B., Gurudatt, N. G., Won, M.-S. & Shim, Y.-B. Analysis of phthalate esters in mammalian cell culture using a microfluidic channel coupled with an electrochemical sensor. Anal. Chem. 87, 7069–7077 (2015). 146. Kolusheva, S., Zadmard, R., Schrader, T. & Jelinek, R. Color fingerprinting of proteins by calixarenes embedded in lipid/polydiacetylene vesicles. J. Am. Chem. Soc. 128, 13592–13598 (2006). 147. Bernards, D. A., Malliaras, G. G., Toombes, G. E. S. & Gruner, S. M. Gating of an organic transistor through a bilayer lipid membrane with ion channels. Appl. Phys. Lett. 89, 53505 (2006). 148. Zhang, Y. et al. Supported lipid bilayer assembly on PEDOT:PSS films and transistors. Adv. Funct. Mater. 26, 7304–7313 (2016). 149. Kwon, N.-H., Rahman, M. A., Won, M.-S. & Shim, Y.-B. Lipid-bonded conducting polymer layers for a model biomembrane: Application to superoxide biosensors. Anal. Chem. 78, 52–60 (2006). 150. Gurudatt, N. G., Naveen, M. H., Ban, C. & Shim, Y.-B. Enhanced electrochemical sensing of leukemia cells using drug/lipid co-immobilized on the conducting polymer layer. Biosens. Bioelectron. 86, 33–40 (2016). 151. Simmons, G. et al. Folate receptor alpha and caveolae are not required for Ebola virus glycoprotein-mediated viral infection. J. Virol. 77, 13433–13438 (2003). 152. Reppy, M. A. & Pindzola, B. A. Biosensing with polydiacetylene materials: Structures, optical properties and applications. Chem. Commun. 0, 4317 (2007). 153. Rangin, M. & Basu, A. Lipopolysaccharide identification with functionalized polydiacetylene liposome sensors. J. Am. Chem. Soc. 126, 5038–5039 (2004). 154. Shao, Y. et al. Conducting polymer polypyrrole supported bilayer lipid membranes. Biosens. Bioelectron. 20, 1373–1379 (2005). 155. Tarabella, G. et al. Organic electrochemical transistors monitoring micelle formation. Chem. Sci. 3, 3432 (2012). 156. Magliulo, M. et al. Electrolyte-gated organic field-effect transistor sensors based on supported biotinylated phospholipid bilayer. Adv. Mater. 25, 2090–2094 (2013). 157. Ahmed, A., Rushworth, J. V, Hirst, N. A. & Millner, P. A. Biosensors for whole-cell bacterial detection. Clin. Microbiol. Rev. 27, 631–46 (2014). 158. Miller, M. B. & Bassler, B. L. Quorum sensing in bacteria. Annu. Rev. Microbiol. 55, 165–199 (2001).

Conducting and Conjugated Polymers for Biosensing Applications


159. Zhang, P. et al. Cationic conjugated polymers-induced quorum sensing of bacteria cells. Anal. Chem. 88, 2985–2988 (2016). 160. Tuncagil, S., Odaci, D., Varis, S., Timur, S. & Toppare, L. Electrochemical polymerization of 1-(4-nitrophenyl)-2,5-di(2-thienyl)-1 H-pyrrole as a novel immobilization platform for microbial sensing. Bioelectrochemistry 76, 169–174 (2009). 161. Rivnay, J. et al. Organic electrochemical transistors for cell-based impedance sensing. Appl. Phys. Lett. 106, 43301 (2015). 162. Ramuz, M. et al. Optimization of a planar all-polymer transistor for characterization of barrier tissue. ChemPhysChem 16, 1210–1216 (2015). 163. Yao, C. et al. Organic electrochemical transistor array for recording transepithelial ion transport of human airway epithelial cells. Adv. Mater. 25, 6575–6580 (2013). 164. Ramuz, M. et al. Combined optical and electronic sensing of epithelial cells using planar organic transistors. Adv. Mater. 26, 7083–7090 (2014). 165. Hempel, F. et al. PEDOT:PSS organic electrochemical transistor arrays for extracellular electrophysiological sensing of cardiac cells. Biosens. Bioelectron. 93, 132–138 (2017). 166. Giaever, I. & Keese, C. R. Micromotion of mammalian cells measured electrically. Proc. Natl. Acad. Sci. U. S. A. 88, 7896–900 (1991). 167. Lin, P., Yan, F., Yu, J., Chan, H. L. W. & Yang, M. The application of organic electrochemical transistors in cell-based biosensors. Adv. Mater. 22, 3655–3660 (2010). 168. Jimison, L. H. et al. Measurement of barrier tissue integrity with an organic electrochemical transistor. Adv. Mater. 24, 5919–5923 (2012). 169. Tria, S. A. et al. Dynamic monitoring of Salmonella typhimurium infection of polarized epithelia using organic transistors. Adv. Healthc. Mater. 3, 1053–1060 (2014). 170. Ramuz, M., Hama, A., Rivnay, J., Leleux, P. & Owens, R. M. Monitoring of cell layer coverage and differentiation with the organic electrochemical transistor. J. Mater. Chem. B 3, 5971–5977 (2015). 171. Spanu, A. et al. An organic transistor-based system for reference-less electrophysiological monitoring of excitable cells. Sci. Rep. 5, 8807 (2015). 172. Yao, C., Li, Q., Guo, J., Yan, F. & Hsing, I.-M. Rigid and flexible organic electrochemical transistor arrays for monitoring action potentials from electrogenic cells. Adv. Healthc. Mater. 4, 528–533 (2015). 173. Dervisevic, M., Senel, M., Sagir, T. & Isik, S. Highly sensitive detection of cancer cells with an electrochemical cytosensor based on boronic acid functional polythiophene. Biosens. Bioelectron. 90, 6–12 (2017). 174. Koutsouras, D. A. et al. PEDOT:PSS microelectrode arrays for hippocampal cell culture electrophysiological recordings. MRS Commun. 7, 259–265 (2017). 175. Mohan, K., Donavan, K. C., Arter, J. A., Penner, R. M. & Weiss, G. A. Sub-nanomolar detection of prostate-specific membrane antigen in synthetic urine by synergistic, dual-ligand phage. J. Am. Chem. Soc. 135, 7761–7767 (2013). 176. Donavan, K. C., Arter, J. A., Weiss, G. A. & Penner, R. M. Virus-poly(3,4-ethylenedioxythiophene) biocomposite films. Langmuir 28, 12581–12587 (2012). 177. Arter, J. A., Taggart, D. K., McIntire, T. M., Penner, R. M. & Weiss, G. A. Virus-PEDOT nanowires for biosensing. Nano Lett. 10, 4858–4862 (2010). 178. Donavan, K. C. et al. Virus–poly(3,4-ethylenedioxythiophene) composite films for impedancebased biosensing. Anal. Chem. 83, 2420–2424 (2011). 179. Arter, J. A. et al. Virus-polymer hybrid nanowires tailored to detect prostate-specific membrane antigen. Anal. Chem. 84, 2776–83 (2012). 180. Ogata, A. F. et al. Virus-enabled biosensor for human serum albumin. Anal. Chem. 89, 1373–1381 (2017). 181. Curto, V. F. et al. Organic transistor platform with integrated microfluidics for in-line multi-parametric in vitro cell monitoring. Microsyst. Nanoeng. 3, 17028 (2017).


Conjugated Polymers

182. Pas, J. et al. Neurospheres on patterned PEDOT:PSS microelectrode arrays enhance electrophysiology recordings. Adv. Biosyst. 2, 1700164 (2018). 183. Inal, S. et al. Conducting polymer scaffolds for hosting and monitoring 3D cell culture. Adv. Biosyst. 1, 1700052 (2017). 184. Sengupta, P. P., Gloria, J. N., Parker, M. K. & Flynt, A. S. A polyaniline-based sensor of nucleic acids. J. Vis. Exp. 117 (2016). doi:10.3791/54590 185. Shi, H. et al. Hyper-branched phosphorescent conjugated polyelectrolytes for time-resolved heparin sensing. ACS Appl. Mater. Interfaces 5, 4562–4568 (2013). doi:10.1021/am4000408 186. Miodek, A., Poturnayová, A., Šnejdárková, M., Hianik, T. & Korri-Youssoufi, H. Binding kinetics of human cellular prion detection by DNA aptamers immobilized on a conducting polypyrrole. Anal. Bioanal. Chem. 405, 2505–2514 (2013). 187. Srinivas, A. R. G., Peng, H., Barker, D. & Travas-Sejdic, J. Switch on or switch off: An optical DNA sensor based on poly(p-phenylenevinylene) grafted magnetic beads. Biosens. Bioelectron. 35, 498–502 (2012). 188. Li, J., Zhao, Q. & Tang, Y. Label-free fluorescence assay of S1 nuclease and hydroxyl radicals based on water-soluble conjugated polymers and WS 2 nanosheets. Sensors 16, 865 (2016). doi:10.3390/ s16060865 189. Chen, Y., Hong, P., Xu, B., He, Z. & Zhou, B. Streptavidin sensor and its sensing mechanism based on water-soluble fluorescence conjugated polymer. Spectrochim. Acta Part A Mol. Biomol. Spectrosc. 122, 441–446 (2014). 190. Zhu, H., Lu, F., Wu, X.-C. & Zhu, J.-J. An upconversion fluorescent resonant energy transfer biosensor for hepatitis B virus (HBV) DNA hybridization detection. Analyst 140, 7622–7688 (2015). doi:10.1039/c5an01634g 191. Liu, X. et al. Target-induced conjunction of split aptamer fragments and assembly with a watersoluble conjugated polymer for improved protein detection. ACS Appl. Mater. Interfaces 6, 3406– 3412 (2014). 192. Wang, C., Tang, Y., Liu, Y. & Guo, Y. Water-soluble conjugated polymer as a platform for adenosine deaminase sensing based on fluorescence resonance energy transfer technique. Anal. Chem. 86, 6433–6438 (2014). 193. Huang, H. et al. Determination of catecholamine in human serum by a fluorescent quenching method based on a water-soluble fluorescent conjugated polymer–enzyme hybrid system. Analyst 137, 1481–1486 (2012). doi:10.1039/c2an16143e 194. Park, M.-K., Kim, K.-W., Ahn, D. J. & Oh, M.-K. Label-free detection of bacterial RNA using polydiacetylene-based biochip. Biosens. Bioelectron. 35, 44–49 (2012). 195. Huang, Y. et al. Cationic conjugated polymer/fluoresceinamine-hyaluronan complex for sensitive fluorescence detection of CD44 and tumor-targeted cell imaging. ACS Appl. Mater. Interfaces 6, 19144–19153 (2014). 196. Du, C. et al. Competition-derived FRET-switching cationic conjugated polymer-Ir(III) complex probe for thrombin detection. Biosens. Bioelectron. 100, 132–138 (2018). 197. Lu, X. et al. Label-free detection of histone based on cationic conjugated polymer-mediated fluorescence resonance energy transfer. Talanta 180, 150–155 (2018). 198. Liu, X. et al. Highly sensitive detection of DNA-binding proteins based on a cationic conjugated polymer via a target-mediated fluorescence resonance energy transfer (TMFRET) strategy. Polym. Chem. 3, 703 (2012). 199. Yu, M. et al. Development of near-infrared ratiometric fluorescent probe based on cationic conjugated polymer and CdTe/CdS QDs for label-free determination of glucose in human body fluids. Biosens. Bioelectron. 95, 41–47 (2017). 200. Qian, C.-G. et al. Conjugated polymer nanoparticles for fluorescence imaging and sensing of neurotransmitter dopamine in living cells and the brains of zebrafish larvae. ACS Appl. Mater. Interfaces 7, 18581–18589 (2015).

Conducting and Conjugated Polymers for Biosensing Applications


201. Huang, H., Li, Y., Liu, J., Tong, J. & Su, X. Detection of bisphenol A in food packaging based on fluorescent conjugated polymer PPESO3 and enzyme system. Food Chem. 185, 233–238 (2015). 202. Huang, Y. et al. Cationic conjugated polymer/hyaluronan-doxorubicin complex for sensitive fluorescence detection of hyaluronidase and tumor-targeting drug delivery and imaging. ACS Appl. Mater. Interfaces 7, 21529–21537 (2015). 203. Yuan, H. et al. Graphene-oxide-conjugated polymer hybrid materials for calmodulin sensing by using FRET strategy. Adv. Funct. Mater. 25, 4412–4418 (2015). 204. Wang, Y. et al. Fluorescence ratiometric assay strategy for chemical transmitter of living cells using H2O2-sensitive conjugated polymers. ACS Appl. Mater. Interfaces 7, 24110–24118 (2015). doi:10.1021/acsami.5b07172 205. Zhang, Z., Xia, X., Xiang, X., Huang, F. & Han, L. Conjugated cationic polymer-assisted amplified fluorescent biosensor for protein detection via terminal protection of small molecule-linked DNA and graphene oxide. Sensors Actuators B Chem. 249, 8–13 (2017). 206. Soylemez, S. et al. Electrochemical and optical properties of a conducting polymer and its use in a novel biosensor for the detection of cholesterol. Sensors Actuators B Chem. 212, 425–433 (2015). 207. Bo, Y., Yang, H., Hu, Y., Yao, T. & Huang, S. A novel electrochemical DNA biosensor based on graphene and polyaniline nanowires. Electrochim. Acta 56, 2676–2681 (2011). 208. Al-Sagur, H., Komathi, S., Khan, M. A., Gurek, A. G. & Hassan, A. A novel glucose sensor using lutetium phthalocyanine as redox mediator in reduced graphene oxide conducting polymer multifunctional hydrogel. Biosens. Bioelectron. 92, 638–645 (2017). 209. Kerr-Phillips, T. E. et al. Conducting electrospun fibres with polyanionic grafts as highly selective, label-free, electrochemical biosensor with a low detection limit for non-Hodgkin lymphoma gene. Biosens. Bioelectron. 100, 549–555 (2018). 210. Hu, Y., Zhao, Z. & Wan, Q. Facile preparation of carbon nanotube-conducting polymer network for sensitive electrochemical immunoassay of Hepatitis B surface antigen in serum. Bioelectrochemistry 81, 59–64 (2011). 211. Piro, B. et al. Direct and rapid electrochemical immunosensing system based on a conducting polymer. Talanta 82, 608–612 (2010). 212. Miodek, A., Castillo, G., Hianik, T. & Korri-Youssoufi, H. Electrochemical aptasensor of cellular prion protein based on modified polypyrrole with redox dendrimers. Biosens. Bioelectron. 56, 104–111 (2014). 213. Holford, T. R. J., Holmes, J. L., Collyer, S. D., Davis, F. & Higson, S. P. J. Label-free impedimetric immunosensors for psoriasin—Increased reproducibility and sensitivity using an automated dispensing system. Biosens. Bioelectron. 44, 198–203 (2013). 214. Hussain, K. K., Gurudatt, N. G., Mir, T. A. & Shim, Y.-B. Amperometric sensing of HIF1α expressed in cancer cells and the effect of hypoxic mimicking agents. Biosens. Bioelectron. 83, 312–318 (2016). 215. Xu, F., Ren, S. & Gu, Y. A novel conductive poly(3,4-ethylenedioxythiophene)-BSA film for the construction of a durable HRP biosensor modified with nanoAu particles. Sensors (Basel) 16, 374 (2016). 216. Gao, Y.-S. et al. Overoxidized polypyrrole/graphene nanocomposite with good electrochemical performance as novel electrode material for the detection of adenine and guanine. Biosens. Bioelectron. 62, 261–267 (2014). 217. Lien, T. T. N. et al. Multi-wall carbon nanotubes (MWCNTs)-doped polypyrrole DNA biosensor for label-free detection of genetically modified organisms by QCM and EIS. Talanta 80, 1164–1169 (2010). 218. Wang, P., Ni, Y. & Kokot, S. A novel dsDNA/polydiphenylamine-4-sulfonic acid electrochemical biosensor for selective detection of the toxic catechol and related DNA damage. Analyst 138, 1141–1148 (2013). doi:10.1039/c2an36389e 219. Istamboulie, G. et al. Screen-printed poly(3,4-ethylenedioxythiophene) (PEDOT): A new electrochemical mediator for acetylcholinesterase-based biosensors. Talanta 82, 957–961 (2010).


Conjugated Polymers

220. Mir, T. A. et al. An amperometric nanobiosensor for the selective detection of K+-induced dopamine released from living cells. Biosens. Bioelectron. 68, 421–428 (2015). 221. Sethuraman, V., Muthuraja, P., Anandha Raj, J. & Manisankar, P. A highly sensitive electrochemical biosensor for catechol using conducting polymer reduced graphene oxide–metal oxide enzyme modified electrode. Biosens. Bioelectron. 84, 112–119 (2016). 222. Özer, B. O. & Çete, S. Development of a novel biosensor based on a polypyrrole–dodecylbenzene sulphonate (PPy–DBS) film for the determination of amperometric cholesterol. Artif. Cells Nanomed. Biotechnol. 45, 824–832 (2017). 223. Gokoglan, T. C. et al. A novel approach for the fabrication of a flexible glucose biosensor: The combination of vertically aligned CNTs and a conjugated polymer. Food Chem. 220, 299–305 (2017). 224. Buber, E. et al. Construction and amperometric biosensing performance of a novel platform containing carbon nanotubes-zinc phthalocyanine and a conducting polymer. Int. J. Biol. Macromol. 96, 61–69 (2017). 225. Kiilerich-Pedersen, K., Poulsen, C. R., Jain, T. & Rozlosnik, N. Polymer based biosensor for rapid electrochemical detection of virus infection of human cells. Biosens. Bioelectron. 28, 386–392 (2011). 226. Tarabella, G. et al. Liposome sensing and monitoring by organic electrochemical transistors integrated in microfluidics. Biochim. Biophys. Acta – Gen. Subj. 1830, 4374–4380 (2013). 227. Daprà, J., Lauridsen, L. H., Nielsen, A. T. & Rozlosnik, N. Comparative study on aptamers as recognition elements for antibiotics in a label-free all-polymer biosensor. Biosens. Bioelectron. 43, 315–320 (2013). 228. Goda, T., Toya, M., Matsumoto, A. & Miyahara, Y. Poly(3,4-ethylenedioxythiophene) bearing phosphorylcholine groups for metal-free, antibody-free, and low-impedance biosensors specific for C-reactive protein. ACS Appl. Mater. Interfaces 7, 27440–27448 (2015). 229. Mak, C. H. et al. Highly-sensitive epinephrine sensors based on organic electrochemical transistors with carbon nanomaterial modified gate electrodes. J. Mater. Chem. C 3, 6532–6538 (2015). 230. Kwon, O. S. et al. Ultrasensitive and selective recognition of peptide hormone using close-packed arrays of hPTHR-conjugated polymer nanoparticles. ACS Nano 6, 5549–5558 (2012). 231. Aydın, E. B., Aydın, M. & Sezgintürk, M. K. A highly sensitive immunosensor based on ITO thin films covered by a new semi-conductive conjugated polymer for the determination of TNFα in human saliva and serum samples. Biosens. Bioelectron. 97, 169–176 (2017). 232. Azak, H., Yildiz, H. B. & Bezgin Carbas, B. Synthesis and characterization of a new poly(dithieno (3,2-b:2′, 3′-d) pyrrole) derivative conjugated polymer: Its electrochromic and biosensing applications. Polymer (Guildf) 134, 44–52 (2018). 233. Gualandi, I. et al. Selective detection of dopamine with an all PEDOT:PSS organic electrochemical transistor. Nat. Publ. Gr. 6, 35419 (2016).

24 Conjugated Poly/­OligoElectrolytes for Cancer Diagnosis and Therapy 24.1 Introduction.......................................................................................743 24.2 Diagnosis.............................................................................................744 Tumor Marker Tests  •  Genetic Tests

24.3 Therapy................................................................................................756

Lingyun Zhou, Guillermo C. Bazan, and Shu Wang

Drug Delivery System  •  Gene Delivery Systems  •  Photodynamic Therapy • Photothermal Therapy

24.4 Summary and Outlook.....................................................................773 References.......................................................................................................774

24.1 Introduction Although the earliest description of cancer dates back to ancient Egypt and extraordinary efforts have been devoted toward understanding its underlying physiological mechanisms and providing different treatment options,1,2 the word “cancer” remains associated with an incurable disease that causes misery, fear, and long lasting family trauma. Cancer is the second most common cause of death in the United States and is exceeded only by heart disease. According to an estimate by the American Cancer Society, more than 1,680,000 new cancer cases are expected to be diagnosed in 2017 and approximately 601,000 Americans will succumb to the disease.3,4 Fortunately, due to our increased understanding of the disease and improvements in early detection and therapies, a substantial number of cancer cases can be prevented and the overall death rates continue to decline. The five-year relative survival rate for all cancer combined increased more than 20 percentage points over the past three decades.1,2 Owing to the high brightness, large extinction coefficients, excellent photostability, low cytotoxicity, stability in bodily fluids and versatile structural modifications, research on water dispersible conjugated polymers/oligomers (CPs/OPs) has demonstrated that these materials offer powerful alternatives in biological applications.5,6 Indeed, certain CPs/OPs can distinguish cancer cells from normal cells. Their versatility leads to uses in targeting tumor imaging, early detection, drug delivery, and cancer therapy.7 High brightness and stability reduce the influence of bias light from biomolecules and endow the assays with better signal to noise ratios when considering imaging and diagnosis platforms. The antenna effect and multivalent interactions can improve the detection sensitivity. Longer luminescence lifetimes aid surgery with whole-time specific discrimination of tumor and normal tissues, leading to less damage to normal peripheral tissue and the nervous system. Amphiphilic CPs form the basis of highly effective drug delivery systems because of their strong interaction with chemotherapeutics, controllable releasing response to diverse 743


Conjugated Polymers

stimuli and noninvasively and real-time monitor of drug loading, delivery, and release. In phototherapy, CPs can generate singlet oxygen (1O2) under light irradiation without the need to include a photosensitizer (PS). With the discovery of near infrared emissive CPs, thermotherapy can also be achieved which generates heat to kill tumor cells under light exposure.8 The purpose of this chapter is to provide an overview of the current status and development of organic conjugated materials designed for cancer diagnosis and therapy.

24.2 Diagnosis Early diagnosis of cancer remains and is anticipated to be one of the most effective way to improve the possibility of complete cure and to avoid benign tumors becoming malignant.9–11 Certain kinds of tumor can be detected easily by routine physical examination. For example, digital rectal examination is a conventional method for detection early colorectal cancer. Self-examination of breasts can inform about cancer through detection of a lump. Roentgenograms and color duplex ultrasonography can be used to observe solid tumors in the lung, liver, kidney, and so on. Abnormal blood and stool routine results could also provide evidence for hematopoietic and intestinal cancers. However, because of the lack of precision from self-examination and the relative low sensitivity of routine examinations, an early phase diagnose of cancer remains unreliable. New technologies are needed and continue to be developed for the earliest possible detection with the goal of achieving a more prompt and specific diagnosis of cancer, and thereby significantly improve the opportunity of full-recovery. Tumors secrete characteristic proteins and the body may also generate specific molecules when tumors are present. These molecules, known as biomarkers, constitute relevant targets for cancer detection.12–14 DNA mutations are also important for detection. Single-nucleotide polymorphism (SNP) is a single nucleotide variation.15–23 Most of the SNPs happen in non-coding sequences and cause little consequence. However, when present within coding regions or a specific non-coding region SNPs can interrupt protein expression and function. DNA methylation is an epigenetic mechanism and plays an important role in human cancer.24–29 Cancer cells detaching from a primary tumor and circulating in the peripheral blood are known as circulating tumor cells (CTCs).30–38 Sorting and detection of these cells help to diagnose cancer during its metastatic stage, which plays a key role in determining the prognosis of the patients and monitoring therapeutic effects. Conjugated polymers and oligomers can be used to identify tumor biomarker and tumor DNA sequences and detect tumor cells.8,39 Because of their high brightness and the antenna effect afforded by the electronic delocalization, the use of conjugated polymers and oligomers can greatly improve the limit of detection (LOD). Signals can be detected by the turning on and off of fluorescence aided by super quenching and recovery mechanisms. Additionally, fluorescence resonance energy transfer (FRET)-based methods can be used to provide detectable large shifts in emission profiles and can be adapted to reduce the impact of nonspecific interactions through the use of ratiometric measurements. Conjugated polymer nanoparticles (CPNs) can be prepared to achieve modifications that add multiple functions. The introduction of an additional amphiphilic polymer, referred to as a co-encapsulated reagent, not only contributes to the formation of nanoparticles (NPs) and to modulate hydrophilicity, but also can be used to introduce new chemical functionalities.40,41 Commonly used co-encapsulated reagents include poly(styrene-g-ethylene oxide) (PS-PEG-COOH), poly(styrene-co-maleic anhydride) (PSMA), poly(DL-lactide-co-glycolide) (PLGA). Both water dispersible and insoluble conjugated polymer/oligomer can be prepared into nanoparticles; hence, using nanoparticles enlarges the range of applications for a given conjugated polymer structure. For example, carboxyl groups on the surface of CPN can be used to introduce recognition elements, such as an antibody.41–45

24.2.1 Tumor Marker Tests By utilizing electrostatic interactions between a charged polymer and an oppositely charged biomarker, without modification of targeting groups, one can detect a biomarker directly through quenching of the emission. As an illustration, the charged conjugated polyelectrolyte brush ThNI (Scheme 24.1, structure 1)

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


SCHEME 24.1  Chemical structure of conjugated materials applied in cancer diagnosis and therapy. Part 1 of 3.


Conjugated Polymers

FIGURE 24.1  (A) Fluorescence spectra of ThNI under the addition of AFP at different concentrations. (B) Response of quenching efficiency of ThNI upon the addition at different concentration (0–100 nM). The inset figure shows the linear relationship in the range of 0–6 nM. (C) Fluorescence spectra of ThNI with the addition of AFP, BSA, IgG, Lys, PSA, and Thro. (D) Quenching efficiency and fluorescence color (inset figure) of ThNI in the presence of AFP and five negative proteins (20 nM each). (From Liu, X.; Shi, L.; Zhang, Z.; Fan, Q.; Huang, Y.; Su, S.; Fan, C.; Wang, L.; Huang, W. Analyst. 2015, 140 (6), 1842. With permission.)

was used to sense the human α-fetoprotein (AFP), a tumor marker for the diagnosis of hepatocellular carcinoma, by observing selective super-quenching in several minutes (Figure 24.1) Moreover, ThNI can specifically recognize AFP, and other negative proteins are much less capable of quenching the fluorescence of ThNI.46 Without specific interactions or a responsive group, any quenching of emission caused by non-specific interactions with CPs has the potential to give rise to a false positive result. In this respect, the specificity of antibody and antigen interactions can be used to greatly improve detection accuracy. Therefore, the introduction of antibody recognition elements can exclusively recognize an overexpressed biomarker. HER2 (Human epidermal growth factor receptor 2), also called as receptor tyrosine-protein kinase erbB-2, is overexpressed in certain aggressive types of breast cancer. The linkage of HER2 antibody to PFBTTB (Scheme 24.1, structure 2) allows for targeting and imaging of HER2 receptors on the surface of MCF-7/HER18 and 435.eB cells (both over-expressed HER2 receptors) (Figure 24.2).47 The prostate-specific antigen (PSA) has been identified as a valuable biomarker to detect prostate cancer. The complex of PSA and inhibitors α1-antichymotrysin (PSA–ACT, MW 90 kDa) and α2macroglobulin (PSA–AMG) constitute two major forms of PSA found in serum.48–52 The antibody against prostate-specific antigen−α1-antichymotrypsin (PSA−ACT) complex was functionalized on a polydiacetylene (PDA) chip. The PDA chip exhibits a nonfluorescent-to-fluorescent transition upon

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.2  The specific targeting of Herceptin-conjugated PFBTTB to HER2-overexpressed cells (a) MCF-7 and (b) 435.eB assessed by laser confocal microscope. MCF7/neo and 435/neo cells have lower expression of HER2 receptors. (From Way T.; Chang C.; Lin C., J. Fluoresc. 2011, 21 (4), 1669. With permission.)

the addition of PSA−ACT, and the LOD was 10 ng·mL−1. To further improve the detection sensitivity so as to meet the clinical requirements of  G) in exon 21, resulting in a leucine (L) to arginine (R) change of codon 858 in EGFR protein. Cationic poly(fluorenephenylene) (PFP1, Scheme 24.1, structure 9) and fluorescein labeled deoxyguanosine triphosphate (dGTP) can be selected to form a FRET pair, since PFP1 exhibits emission maximums at 425 and 445 nm, which overlap with the absorption spectrum of fluorescein. Irradiation at 380 nm selectively excites PFP1 and FRET from CCP (donor) to fluorescein (acceptor) is favored. Target EGFR gene of sample was firstly amplified by nested PCR. L858R primer was conducted by SBE, and after removing of leftover primers and dNTPs, PFP1 was added to interact with DNA for FRET. Because the designed primer L858R cannot match wild type EGFR sequence, SBE is not favored. FRET could only be measured on mutated DNA. Upon measuring the FRET ratio from PFP1 to fluorescein (fluorescent intensity on maximum emission of fluorescein divided by that of PFP1, I528nm/I440nm), the proportion of mutant DNA can be determined (Figure 24.9).85 As we have mentioned above, CCP-based SNP detection reduces cost and complexity, while the combination of the molecular wire effect and FRET is capable of enhancing the detection signal. Utilizing PFP1 with a similar procedure, more than one site of mutation can be detected at the same time. Furthermore, the results can be visualized by human-unaided eyes. Mutations on PIK3CA gene have been discovered in different human cancers. Specifically, E545K (G1633A) and H1047R (A3140G), two of four hotspot mutations (E542K (G1624A), E545K (G1633A) in exon 9 (helical domain), and H1047R (A3140G), H1047L (A3140T) in exon 20 (kinase domain)), were chosen as detection targets. Two SBE on E545K primer and H1047R primer were conducted utilizing dATP-TR (TexRed label), dUTP-Fl (fluorescein label) and Taq polymerase. The 3′-terminal base of two primers individually complimented the mutation site of E545K and H1047R, but not the wild type. After adding PFP1, different FRET phenomena (PFP1→Fl, PFP1→TR, Fl→TR) occur depending on single E545K mutation, single H1047R mutation, both E545K and H1047 mutation, and wild type. The fluorescent color is also varied from each situation and can be distinguished by observation (Figure 24.10).86 Additionally, multistep FRET strategies can be designed by involving four types of labeled ddNTP, and a fingerprint spectrum forms via the number of FRET pairs. Not SBE, but regular PCR, is conducted with extra dye-labeled ddNTP to terminate the PCR progress. Different mutations lead to different possibilities of termination, further changing the FRET consequences. By the analysis of the resulting fingerprints, not only SNP, but also insertion and deletion can be detected and differentiated (Figure 24.11).87

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.9  (a) Schematic detection strategy using PFP1-based FRET method. L858R primer is perfectly complementary to mutation EGFR gene rather than wild type DNA. As a result, the primer is extended by dGTP-Fl complementary to the base adjacent to mutation base when the extension primer- L858R primer can completely match the PCR product template. SBE indicates single base extension. FRET indicates fluorescence resonance energy transfer. (b) The FRET ratio (I528nm/I440nm) from solutions containing PFP1 and extension products of various percentages of mutant EGFR gene fragments obtained by PCR amplifying recombinant pGEM-T plasmids. The agarose gels show the PCR products amplified from pGEM-T L858R and pGEM-T L858 plasmids. (c) Standard curve for mutation percentages generated by various ratios of pGEM-L858R to pGEM-L858 plasmids. NC indicates negative control. For the mean, error bars indicate ± s.d. (n = 3). All the experiments were performed in Hepes buffer solution (25 mM, pH 7.5). The excitation wavelength is 380 nm. [dGTP-Fl] = 5 × 10−8 M; and [PFP1] = 1.5 × 10−6 M. (From Yang, Q.; Qiu, T.; Wu, W.; Zhu, C. L.; Liu, L. B.; Ying, J. M.; Wang, S. ACS Appl. Mater. Interfaces 2011, 3 (11), 4539. With permission.) DNA Methylation Detection DNA methylation is an epigenetic regulation method of gene expression. Hypermethylation of the CpG islands which contain a high frequency of 5'-C-phosphate-G-3' (CpG) sites in promoter regions of tumor suppressor genes can downregulate or inhibit its expression, leading to abnormal cell proliferation. With the aid of the interaction between CCP and DNA, together with FRET from CCP to fluorescently labeled dNTP, like the method used in the detection of SNP, CCP can also be applied in the detection of DNA methylation. The procedure uses bisulfite treatment to replace cytosine by uracil. In the wild type DNA, i.e. without methylation, the cytosine of the CpG site was modified into uracil upon bisulfite treatment, and the uracil was substituted by thymine after PCR amplification. For methylated CpG sites, the cytosine remains unchanged. Fluorescein labeled dGTP and a specific primer, whose 3′-terminal base is C, complementary to the target sequence from methylated DNA region was used in the performance of SBE. The dGTP-Fl was incorporated into the probe by extension reaction for methylation DNA but not the wild type. Upon addition of the cationic conjugated polyelectrolyte PFP1, efficient FRET from PFP1 to fluorescein takes place on methylation groups. In contrast, FRET was sharply weakened for the wild type case. Utilizing another primer, whose 3′-terminal base is T, the FRET would change over, which means, FRET will occur on wild type group rather than methylation groups (Figure 24.12).88


Conjugated Polymers

FIGURE 24.10  (A) Schematic representation of visual assay for mutations in the PIK3CA gene. E545K is located at exon 9 and H1047R on exon 20, and blue arrows show the occurrence of distance-dependent energy transfer. (B) Fluorescence spectra of E545K, E545K/H1047R, H1047R, and wild type by mixing SBE products with PFP1. SBE products were diluted by 100 times with HEPES buffer solution (25 mM, pH 8.0) before fluorescence measurement. [PFP1] = 0.2 μM. The excitation wavelength is 380 nm. (C) The corresponding images of E545K, E545K/H1047R, H1047R, and wild type when SBE products mixed with PFP1 (15 μM) in PCR tubes under 365 nm UV light irradiation. High-purity water was used as the blank. (From Song, J.; Yang, Q.; Lv, F.; Liu, L.; Wang, S. ACS Appl. Mater. Interfaces 2012, 4 (6), 2885. With permission.)

A more convenient approach takes advantage of HpaII, a methylation-sensitive restriction endonuclease. Cumbersome bisulfite treatment can be omitted. All unmethylated recognition sites can be cleaved while conserving those that are methylated. The subsequent nested PCR amplification procedure incorporates Fl-dNTPs to uncleaved DNA (methylated DNA), but not for unmethylated one. Upon addition of PFP1, distinct FRET from PFP1 to Fl was observed.88 The methylated states of three cancer suppressor genes (p16, HPP1, and GALR2 promoters) from five cancer cell lines, HT29, HepG2, A498, HL60, and M17 can be detected through E-value. Here, E = (R Hpall-Rcontrol)/(R Mock-Rcontrol), where R HpaII refers to FRET ratio (I530nm/I424nm) for the HpaII-PCR sample, while R Mock denotes the FRET ratio (I530nm/I424nm) for the mock sample and Rcontrol denotes the FRET ratio (I530nm/I424nm) for the negative control. Mock reactions refer to exactly the same treatments from HpaII digestion to FRET measurements, except without HpaII. The results afford excellent correlation information between cancers and susceptibility genes, which is very useful for early cancer diagnosis. Taking numbers of different candidate genes, cumulative detection of methylation alteration accomplishes the differential diagnosis of colon cancer and ovarian cancer from tissue samples (Figure 24.13).89,90 In summary, CCP-based FRET techniques provide rapid, sensitive, and flexible strategies for DNA detection. Expensive instrumentation and technical expertise are not required. With the aid of fingerprint spectrum analysis, CCP has a universal applicability to detect DNA sequence variation, such as nucleotide deletion, insertion, and substitution and suitable for the detection of non-adjacent and adjacent mutations. The limitation of CCP-based SNP detection is that only known SNP sites can be detected and PCR amplification for subsequent SNP genotyping is required, thereby restricting high-throughput analysis.

24.3 Therapy Conjugated polyelectrolytes and oligoelectrolytes are promising candidates for the development of new cancer therapies. There are several approaches for CPs and COs to achieve tumor-targeted function.

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.11  Presentation of the CCP-based FRET fingerprint technique for DNA mutation detection and standard fingerprint spectra for different mutations. ΔFRET ratio is the difference value of the FRET ratio in the various DNA mutations: (A) PIK3CA mutations stand for non-adjacent mutations; (B) KRAS mutations stand for adjacent mutations; (C) deletion mutations; (D) insertion mutations. Multiplex FRET can be triggered at a single excitation wavelength of 380 nm. (From Song, J.; Zhang, J.; Lv, F.; Cheng, Y.; Wang, B.; Feng, L.; Liu, L.; Wang, S. Angew. Chemie - Int. Ed. 2013, 52 (49), 13020. With permission.)

The water dispersibility of cationic CP and CO provides a basic precondition for application in biosystems. Cationic systems provide the possibility of interactions with cells and negatively charged biomolecules. To increase the specificity toward cancer cells, antibodies and tumor–microenvironment response moieties can be selected and linked onto CP/CO molecular frameworks. Furthermore, non-water-dispersible CP/CO systems can be processed to yield nanoparticles. Along with increased dispersibility, NPs can load drugs and serve as drug delivery systems (DDS). Drugs loaded onto CP and CO carriers can be delivered and monitored in real-time because of the fluorescent property of CP/ CO and tumor targeting moiety modified on CP/CO molecules. Upon light irradiation, excitations on CP and CO can transfer to photosensitizers to produce reactive oxygen species that kill cancer cells. Moreover, CP and CO can serve alone as a photosensitizer and produce sufficient reactive oxygen species (ROS) for photodynamic therapy (PDT) of tumor. With the development of long wavelength excitation polymers, like red light and IR excitation, CPs can also be applied in photothermal therapy (PTT). In this part, we will describe some of the applications of CPs and COs in drug delivery, PDT, and PTT.


Conjugated Polymers

FIGURE 24.12  (A) Schematic representation of the assay for methylation status of specific CpG sites. (B) Fluorescence spectra (a) and FRET ratio (I530nm/I422nm) (b) of the extension production with varying proportions of methylated DNA in the presence of PFP1. The amounts used were 0.67 pmol of probe, 1.67 pmol of dGTP-Fl, and [PFP1]=0.25 μM. The excitation wavelength is 380 nm. (From Feng, F.; Wang, H.; Han, L.; Wang, S. J. Am. Chem. Soc. 2008, No. 13, 11338. With permission.)

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.13  (A) Schematic representation of CCP-based DNA methylation detection. The methylation-sensitive restriction endonuclease (HpaII) digests unmethylated recognition sites of unmethylated DNA (seen on the left-hand side) while conserving those of methylated DNA (seen on the right-hand side). The Taq polymerase incorporates the fluorescein-labeled dNTP (Fl-dNTP) into the PCR product of target methylated DNA sequence, but not unmethylated DNA by extension reaction. By triggering the fluorescence resonance energy transfer (FRET) signal from PFP1 to fluorescein, DNA methylation is detected. (B) CCP-based quantitative analysis of DNA methylation. E-value is introduced to determine the methylation status of promoter regions of three genes (p16, HPP1, and GALR2) in different cancer cell lines (HT29, HepG2, A498, HL60 and M17). A high E-value that approaches 1.0 represents a high methylation level, a low E-value that approaches zero represents unmethylation and midvalue of E represents partial methylation. (From Feng, F.; Liu, L.; Wang, S. Nat. Protoc. 2010, 5 (7), 1255. With permission.)

24.3.1 Drug Delivery System Chemotherapy is one of the most effective and most common treatments for cancer. However, most of the anticancer drugs are associated with systemic toxicity to normal cells.91–96 Introducing a target specific anticancer DDS can avoid interference with normal cell function and deliver chemotherapeutics specifically to cancer sites.97–100 Polymer-based delivery systems have been comprehensively studied.101–109 Conjugated polyelectrolytes also exhibit their advantages in DDS.7,110,111 The loading and delivery process can be monitored in real time without extra linkage of fluorescent molecules. Drugs can be covalently or non-covalently attached to CP/CO molecules. Side chain modifications of conjugated molecules are convenient and well established, as well as the hydrophobic interaction between the conjugated backbones and mostly hydrophobic chemotherapeutics provides a means of attachment through non-covalent bonding. CPs in aqueous media are mostly present as micelles with diameters ranging from tens to hundreds of nanometers. After uptake by cancer cells, more likely through endocytosis and pinocytosis, the complexes are transferred by endosomes towards lysosomes. Hence, drugs need to be detached from DDS to have their desired therapeutic effect on their targets. Many responses to tumor microenvironment and adscititious stimuli have been designed and introduced to the DDS for effectively release the drug from CPs. The commonly referred responses are the higher expression of GSH by tumor cells, varied redox potentials, lower pH of extracellular matrix in tumor environment, and abnormal enzyme and glucose levels. Heat, light, ultrasound, magnetism, and electric fields can be used in bond breaking methods as external stimuli strategy for drug release.101,112–125 When incorporated into CPNs, the loading capacity of chemotherapeutics can be increased. After delivery into cancer cells, the CPNs disassemble and release the loaded chemotherapeutics. The enhanced permeability and retention effect of CPNs with dimensions of approximately 200 nm provides a unique targeting property, in that CPNs tend to accumulate with more propensity in tumor


Conjugated Polymers

tissue relative to normal tissues. COs have their own advantages. With smaller molecular weights, they migrate with a higher speed, are more effective in penetrating the cell membrane, and accumulate in the cytosol rather than the lysosomes. Cationic conjugated polyelectrolyte PFO (Scheme 24.2, structure 10) was applied for the first time in simultaneous imaging and DDS due to its good quantum yield (QY) and little cytotoxicity. Electrostatic assembly of PFO and anionic poly(L-glutamic acid) conjugated with anticancer drug Dox (PFO/ PG-Dox) forms 50 nm particles, which can be internalized by tumor cells. When present within PFO/ PG-Dox complexes, the fluorescence of PFO is quenched by Dox (fluorescence “turn-off”). After the carrier is transported into cancer cells, the PG is hydrolysed by intracellular carboxypeptidase and Dox is released, leading to the fluorescence recovery (fluorescence “turn-on”). The resulting fluorescent intensity can be correlated to the loading and detachment of Dox, and the release of the drug can be visualized in the cells through fluorescence microscopy (Figure 24.14).41 Conjugated oligomer pentathiophene (5T, Scheme 24.2, structure 11) can be used in both anticancer activity design and molecular imaging techniques. 5T selectively accumulates in mitochondria to provide specific organelle imaging capabilities. Many positively charged oligomers tend to gather in mitochondria because of the electrostatic potential difference between the inner and outer membrane. Through electrostatic interaction with sodium chlorambucil, an anti-cancer drug that damages DNA function, 5T forms nanoparticles with a diameter of ca. 50 nm and delivers sodium chlorambucil to the mitochondria. Such nanoparticles exhibit 2–9 fold more cytotoxicity than free sodium chlorambucil, along with the ability to image mitochondria.126 A CPN-based light-triggered releasing vector has also been developed with functionalities intended to enable specific targeting. Dox was conjugated to a PEGylated PFVBT (Scheme 24.2, structure 12) through a ROS cleavable linker (PFVBT-g-PEG-Dox, Scheme 24.2, structure 13). PFVBT-g-PEG-Dox can self-assemble into nanoparticles. For a better selectivity toward tumor cells, a cyclic arginine–glycine–aspartic acid tripeptide (cRGD) targeting αvβ3 integrin overexpressing cancer cells was modified on the surface of the nanoparticles. The target PFVBT-g-PEG-Dox NPs are denoted as TCP-Dox NPs. Because PFVBT can produce ROS under irradiation, the produced ROS will further cleave the ROS sensitive linker and release Dox from the DDS. Furthermore, excessive ROS can further injure cancer cells thereby providing an additional photodynamic therapy effect (Figure 24.15).127 For in vivo applications, excitation and emission using short wavelengths faces the problem of limited tissue penetration. An innovative CP-based nanocarrier fitted with near infrared (NIR) imaging and adenosine-5’-triphosphate (ATP)-responsive anticancer drug release was developed. Specifically, PFFP (Scheme 24.2, structure 14) was modified through an amidation reaction using 3-fluoro-4-carboxyphenylboronic acid (FPBA) to introduce phenylboronic aid (PBA) tags as binding sites for ATP and sialic acid. PEG can be introduced via reaction with the remaining amino groups for better biocompatibility and reduced nonspecific interactions. The drug loading capacity was attributed to the hydrophobic interaction between Dox and polymer backbones. Dox and PFFP can form NPs with a diameter of approximately 100 nm. Most tumor cells overexpress sialic acid. Thus, before entering cells PFFP anchors on the tumor cell through recognition and interaction between phenylboronic acid on polymer and sialic acid on glycoprotein outside of membrane. After endocytosis, because of the distinctly higher ATP levels of the intracellular milieu compared to the extracellular matrix, ATP binds on the phenylboronic acid and promotes the disassembly of the drug carrier, due to the transition from hydrophobic to hydrophilic side chain. PFP without the modification of PBA serves as control polymers, which is used for explanation of the targeting function. The dependence of release on ATP was also studied to ensure the validity of the overall strategy. It is worth pointing out that the overall drug release can be monitored in real time monitored in HepG2 tumor-grafted mice (Figure 24.16).128 In most cases, NPs serve as drug delivery vector and can, to some extent, overcome drug resistance of tumor cells. Application of nanocarriers provides the means to deliver larger amounts of chemotherapeutics within a short period of time and tumor cells have more difficulty eliminating drugs from the cytosol. In a recent study, a cationic oligo(p-phenylenevinylene) (OPV, Scheme 24.2, structure 15)

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


SCHEME 24.2  Chemical structure of conjugated materials applied in cancer diagnosis and therapy. Part 2 of 3.


Conjugated Polymers

FIGURE 24.14  (A) Schematic illustration of PFO/PG-Dox complex system, where the fluorescence of PFO is highly quenched by dox resulting in the fluorescence “turn-off’ state. (B) Schematic representation of uptake of the electrostatic complex into cancer cells. Hydrolysis of poly(l-glutamic acid) by hydrolase in lysosome releases the Dox to induce the activation of PFO fluorescence to “turn-on” state, thereby sensing the intracellular delivery of Dox and simultaneously achieving fluorescent imaging. (C) Fluorescence emission spectra of PFO with successive additions of PG-Dox. (D) Ksv plot of PFO in the presence of PG-Dox. [PFO]=1.0 × 10−6 M, [Dox]=0–1.03 × 10−7 M. (E) Drug release profiles at various time intervals after adding carboxypeptidase. [PFO]=1.0 × 10−6 M, [Dox]=1.03 × 10−7 M, 0.5 U of carboxypeptidase. Fluorescence measurements were performed in phosphate buffer (100 mM, pH 5.8) with excitation wavelength of 380 nm. FG Microscopy images of A549 cells after incubation with PFO/PG-Dox complex for 4 h at 37°C, washed two times with PBS buffer, and further incubated at 37°C for (F) 0 and (G) 24 h. PFO and Dox are shown in blue and red, respectively. (H) The fluorescence recovery of PFO in the cells with Dox release. [PFO]=1.0 × 10−6 M, [PG-Dox]=3.35 × 10−6 M. Fluorescence image of PFO was recorded by fluorescence microscopy (Olympus 1 × 71) using a 380/30 nm excitation filter with 200 ms exposure time and that of Dox was recorded using a 455/70 nm excitation filter with 1000 ms exposure time. (From Feng, X.; Lv, F.; Liu, L.; Tang, H.; Xing, C.; Yang, Q.; Wang, S. ACS Appl. Mater. Interfaces 2010, 2 (8), 2429. With permission.)

was applied toward combating the drug resistance of cancer cell toward Dox-based chemotherapeutics through a novel approach. First, OPV molecules were incubated with cancer cells for an exact period of time determined their interaction with the membrane but before they are internalized inside the cytosol. Free OPV molecules were then removed from the medium, followed by incubation of cells with Dox. Without light irradiation, there is no significant difference relative to results when Dox was incubated alone. Surprisingly, after light irradiation, obvious Dox uptake was observed and the cytotoxicity

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.15  (A) Schematic illustration of the light-regulated ROS-activated on-demand drug release and the combined chemo–photodynamic therapy. PEG = poly(ethylene glycol). (B) CLSM images of free DOX, TCP NPs, TCP-DOX NPs, and TCP-DOX NPs with the ROS scavenger VC incubated with MDA-MB-231 cells with and without irradiation with light. Red: fluorescence of TCP-DOX NPs; blue: cell nuclei stained by Hoechst 33342. All images share the same scale bar (20 mm). (C) Flow cytometry analysis of the DOX fluorescence intensity in MDA-MB-231 cell nuclei. (D) Viability of MDA-MB-231 cells incubated with free DOX, TCP NPs, and TCP-DOX NPs with (5 or 15 min) or without irradiation with light. Control samples were cells without any treatment. TCP NPs are control molecules without the linkage of Dox. (From Yuan, Y.; Liu, J.; Liu, B. Angew. Chemie - Int. Ed. 2014, 53 (28), 7163. With permission.)


Conjugated Polymers

FIGURE 24.16  (A) The main components of DOX/PFFP nanoparticles (NPs): near-infrared (NIR) fluorescent core and flexible hydrophilic PEG chains as shell by self-assembly of amphiphilic conjugated polymer (PFFP), and the 3-fluoro-4-carboxyphenylboronic acid (FPBA) tags on the surface designed as binding sites for ATP. (B) ATP-responsive delivery of DOX by DOX/PFFP NPs to nuclei for the targeted cancer therapy. (C) SEM images of DOX/PFFP NPs at different ATP concentration. Scale bars are 200 nm. (D) Normalized UV/Vis absorption and emission spectra of PFFP and DOX/PFFP NPs (excitation at 480 nm). (E) In vitro release of DOX from DOX/PFFP NPs in 0.4 mM, 4 mM ATP, and DOX/PFP NPs (without ATP sensitivity) in 4 mM ATP. Error bars indicate SD (n = 3). (F) Intracellular delivery of (a) DOX/PFFP NPs and (b) DOX/PFP NPs into HepG2 cells at different time observed by CLSM. The endosomes and lysosomes were stained by LysoTracker Green, and the nuclei were stained by Hoechst 33342. Merged (DOX/LysoTracker/Hoechst). Scale bars are 20 μm. (G) Intracellular DOX/PFFP NPs-mediated ATPtriggered DOX release. Intracellular delivery of DOX/PFFP NPs in HepG2 cells treated with different formulations observed by CLSM, including at 37°C, 4°C, and with the ATP inhibitor iodoacetic acid (IAA) at 37°C. ATP generation can be inhibited in low temperature of in the pre presence of IAA.The endosomes and lysosomes were stained by LysoTracker Green, and the nuclei were stained by Hoechst 33342. Merged (DOX/LysoTracker/Hoechst). Scale bar is 20 μm. (H) In vivo fluorescence images of the HepG2 tumor-bearing mice at 4, 12, and 24 h after intravenous injection of DOX/PFFP NPs. Arrow indicates the sites of tumor. (I) Representative images of the HepG2 tumors after treatment with different samples at day 14 (from top to bottom, 1: saline, 2: DOX, 3: DOX/PFP NPs, 4: DOX/PFFP NPs. (J) The HepG2 tumor growth curves after treatment with different samples. (K) The body weight variation of HepG2 tumor-bearing mice during treatment. (From Qian, C.; Chen, Y.; Zhu, S.; Yu, J.; Zhang, L.; Feng, P.; Tang, X.; Hu, Q.; Sun, W.; Lu, Y.; Xiao, X.; Shen, Q. D.; Gu, Z. Theranostics 2016, 6 (7), 1053. With permission.)

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


was enhanced. The reason is that the ROS generated by OPV under irradiation damaged the cell membranes, especially the double bond in the phosphate functional group. Pores formed in the membrane allowing Dox to more readily penetrate into the cytosol. However, the strategy only works well on Dox, but not other drugs, such as paclitaxel (TAX), vincristine (VCR), cisplatin (DDP), and 5-fluorouracil (5-Fu). Such results highlight that the interactions between anticancer drugs and photosensitizers play an important pole in the light-assisted multidrug resistance reversal (Figure 24.17).129

24.3.2 Gene Delivery Systems Gene therapy is under rapid development. A wide variety of diseases can be treated using gene therapy, including cancer, neurodegenerative, antiviral, hematological, and hereditary genetic disorders, with promising success.130–133 Along with viral vectors, nanohydrogels, silica nanoparticles and polyethyleneimine (PEI) or other liposomes, CPs/COs and their nanoparticles are potential candidates for gene delivery.134 Due to the negative charges of DNA and RNA, cationic conjugated molecules can interact and thus be used to load the gene cargo. Additionally, owing to the biocompatibility of CPs/COs, when serve as drug delivery vectors, CPs/COs exhibit little immune response. Early in 2008, the cationic PDA nanovesicle, DAMDPA-bis-PCDA (Scheme 24.2, structure 16) as a monomer, was used for in vitro gene delivery. Although PDA showed little cytotoxicity, it successfully transferred plasmid DNA to human embryonic kidney (HEK) 293 cells with an acceptable level of gene expression, and compared to PEI, PDA maintained higher cell viability.135 Target gene knockdown via degradation of targeted homologous cellular mRNA by forming RNAinduced silencing complexes (RISCs) is an important strategy to regulate intracellular gene expression. Another poly(fluorenephenylene), denoted as PFP2 (Scheme 24.3, structure 17), was also designed and studied in the delivery of MDR1-targeted siRNA. Overexpression of P-glycoprotein, the product of MDR1 causes efflux of intracellular hydrophobic chemotherapeutics, including Dox and paclitaxel, which lower the drug concentration inner cell and provides a mechanism for drug resistance. PFP2 and

FIGURE 24.17  (A) The mechanism of light-triggered drug uptake approach to reverse drug resistance of resistant cells. (B) The cellular location of DOX and OPV in MCF-7/DOX cells under different conditions. [OPV] = 10 μM, [DOX] = 50 μM. (C) Dose-response curves for cell viabilities of MCF-7/DOX cells incubated with DOX for 24 h in the presence or absence of OPV upon white light irradiation for 1 h or under dark. (D) Observed enthalpy changes ΔHobs against the molar ratio of OPV/drug by titrating 0.50 mM OPV into 0.10 mM drug solutions. The solid lines are the corresponding fi tted curves. ΔHobs values are expressed in kJ per mol OPV. The dilution enthalpy of OPV has been deducted. Except DOX, other drugs shows insufficient interaction with OPV (From Wang, B.; Yuan, H.; Liu, Z.; Nie, C.; Liu, L.; Lv, F.; Wang, Y.; Wang, S. Adv. Mater. 2014, 5986. With permission.)


Conjugated Polymers

SCHEME 24.3  Chemical structure of conjugated materials applied in cancer diagnosis and therapy. Part 3 of 3.

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


MDR1–siRNA complexes (PFP/siRNA) efficiently achieved the reversal of drug resistance and enhanced drug sensitivity (Figure 24.18).136 Considering the trapping of gene and vector in the endosome and degradation by lysosomes, the high-effective release of the gene delivery system from the endosome after internalization by cells should significantly increase the delivery efficiency. Hence, an endosomal escaping strategy was designed with the application of PPV (Scheme 24.3, structure 18) as the carrier. Cationic PPV can serve as a siRNA delivery vector due to its cationic and amphipathic features. After self-assembling with siRNA by electrostatic interactions, PPV/siRNA polyplexes can be endocytosed by cells. When exposed to white light with a sufficient intensity, PPV can generate ROS to disrupt the endosomal membrane, leading to the leakage of transported gene (Figure 24.19).137

FIGURE 24.18  (A) The gel result of cDNA from MDR1 mRNA by siRNA interference with PFP2 and Lipo 2000 as transfection agents (CCP indicates the PFP2 group). β-actin was used as an internal control. (B) MCF7/ADR cell viability in the presence of Dox before and after treatment with PFP/siRNA, and that of treatment only with PFP/siRNA as control. (From Feng, X. L.; Lü, F. T.; Liu, L. B.; Wang, S. Chinese Sci. Bull. 2013, 58 (22), 2762. With permission.)

FIGURE 24.19  (A) Schematic illustrations of the mechanism of the white light-enhanced endosomal escape strategy for high siRNA delivery. (B) CLSM images for endosomal escape of PPV/siRNA polyplexes with or without light irradiation. (C) Co-localization of PPV/siRNA polyplexes with the endosome after incubation for 24 h. Blue, green and red represent Lysotracker, siRNA, and PPV, respectively. (D) The transfection efficiency of PPV/siRNA polyplexes was quantified by FACS using Cy5-labeled siRNA. PEI and lipofectamine 2000 were chosen as positive control. (E) Luciferase intensity in HeLa-Luc cells after treatment with PPV/siRNA polyplexes at 1 mg anti-Luc siRNA for 24 h. (From Li, S.; Yuan, H.; Chen, H.; Wang, X.; Zhang, P.; Lv, F.; Liu, L.; Wang, S. Chem. - An Asian J. 2016, 11 (19), 2686. With permission.)


Conjugated Polymers

24.3.3 Photodynamic Therapy Although light was used as a therapeutic agent hundreds of years ago for the treatment of skin diseases like vitiligo and lupus vulgaris, the understanding of how a photosensitizer (PS) works and the demonstration of PDT occurred during the early 20th century. Certain PS agents can convert triplet oxygen in the surroundings to singlet oxygen upon irradiation. As shown in Figure 24.20, the excited state of PS undergoes intersystem crossing (ISC) thus opening an opportunity to transfer energy and for converting triplet state oxygen into singlet oxygen.138 Singlet oxygen is a type of ROS. It can react with organic chemicals and water to produce radicals and other ROS. Excessive ROS induces cell apoptosis and also necrosis. PDT is the application of photochemistry based on modality using a PS, a light source, and tissue oxygen for the killing of diseased cells. The diffusion range of 1O2 is reported to be limited to a range of 45 nm in cellular media, hence with a specific delivery of PS to cancer sites, the controlled killing of diseased cells will not cause toxicity to adjacent normal tissues. For the treatment of cancer, eosin was applied as photosensitizer in the first PDT trial in patients with skin carcinoma. In 1983, Hematoporphyrin was used as a PS to treat tumors.139 Sustained efforts since then have been made toward developing new techniques and novel chromophores for application in PDT. A primary consideration is that PDT is a non-invasive method with less undesirable side effects, when compared with chemotherapy and radiation therapy.140–148 Small PS molecules have been applied in PDT treatment in the past. However, many of them face low absorption cross sections, cytotoxicity, photo bleaching, and they can be easily metabolized and eliminated by cells. Because of their large absorption cross section, wide range of functionalization options, and ways to process so as to increase intake by cells, conjugated polymers and oligomers serve as an advantageous option in PDT. They can not only emit fluorescence and transfer energy to photosensitizers to generate 1O2, but they also can participate directly as a photosensitizer and transfer excitation energy to oxygen. Both approaches have their own advantages. When CPs/COs serve as PSs themselves, the addition of small molecular PSs can be omitted. When CPs/COs transfer energy to added PS, FRET can amplify the sensitization via the molecular wire effect and the dosage of PS can be greatly reduced. As small molecular PSs are always considered to be toxic as well as their leakage and non-specific interaction being non-negligible, this method reduces the side effect encountered when only small molecular PSs are used. Additionally, as we described in previous section, CP/CO are water dispersible and provide versatile modification options for selective recognition of tumor sites. An early approach of CPs in PDT relied on polythiophene–porphyrin conjugates (PTP, Scheme 24.3, structure 19), which transfer energy from the polythiophene backbone to the porphyrin units in order to produce 1O2 (Figure 24.21). Excitation of PTP at 470 nm, where the absorption signal of porphyrin

FIGURE 24.20  Jablonski diagram showing the formation of singlet oxygen. (From Ethirajan, M.; Chen, Y.; Joshi, P.; Pandey, R. K. Chem. Soc. Rev. 2011, 40 (1), 340. With permission.)

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.21  (A) Apoptosis and imaging of A498 and A549 cancer cells by PTP upon irradiation at 470 nm with a 2.5 J cm−2 source. Phase contrast bright field images (upper) of A498 cells upon light irradiation from 0 to 30 min in the presence of PTP and overlapping fluorescence images of A498 cells (bottom). (B) Dose-response curve data for cell viability of A498 and A549 cells treated with PTP and PT by using a typical MTT assay under 470 nm light irradiation for 30 min. PT is control molecule which is absence of porphyrin. (C) Phase contrast bright-field and fluorescence images of KB cancer cells (upper) and NIH-3T3 cells (lower) in the presence of PTPF (20.0 μ g mL−1). Left: phase contrast bright-field images. Right: fluorescence images. The false color of PTPF is yellow. (D) Doseresponse curves for cell viability of KB cells and NIH-3T3 cells treated with PTPF by using a typical MTT assay under 470 nm light irradiation or in the dark. Error bars correspond to standard deviations from three separate measurements. (From Xing, C.; Liu, L.; Tang, H.; Feng, X.; Yang, Q.; Wang, S.; Bazan, G. C. Adv. Funct. Mater. 2011, 21 (21), 4058. With permission.)

units does not appear, leads to emission with peaks at both 578 nm and 658 nm. The peak at 658 nm demonstrates suitability for energy transfer from polythiophene to the porphyrin sites. The anti-tumor ability of PTP under 470 nm light irradiation was tested on pulmonary adenocarcinoma cell (A549) and renal cell carcinoma (A498). Apoptosis was observed after light irradiation through cellular morphology changes that include chromatin compaction, cytoplasm condensation, large amounts of blebbing and further whole-cell shrinkage. It is possible to observe greater than 90% inhibition of cell viability in the PDT system, but without the introduction of porphyrin, the inhibition rate was negligible. Further, with the introduction of folic acid via covalent bonds with PTP to form PTPF (Scheme 24.3, structure 20), selectivity can be improved toward folate receptor overexpressed cancer cells. PTPF exhibits prominent inhibition of the KB cell line (this KB line was originally thought to be an oral epidermal carcinoma but is now known to be a subline of the ubiquitous keratin-forming tumor cell line HeLa, for HeLa marker chromosomes were subsequently found), but not for NIH-3T3, a normal mouse embryo fibroblast cell line.149 Another polythiophene–porphyrin conjugate linked with the tumor targeting drug tamoxifen (PTDP, Scheme 24.3, structure 21), was designed to achieve greater specificity. Tamoxifen is an inhibitor of estrogen receptor α (ERα), which is overexpressed in the breast cancer cell MCF-7. PTDP is endowed with the ability to target and temporarily inhibit ERα under the aid of tamoxifen. When irradiated with light, ROS produced from PTDP can permanently damage the ERα. PTDP shows selective growth inhibition of ERα positive cancer cells, while providing low side effects for our intracellular moleculetargeted therapy system (Figure 24.22).150 In addition to CPs, CPNs can also be used for PDT. As described previously, CPNs can be finetuned to accumulate in tumor sites through the enhanced permeability and retention effect and they can take advantage of surface modification to realize more specific targeting and more efficient tumor cell penetration. Furthermore, if the CPNs successfully target and anchor on the surface of tumor cells,


Conjugated Polymers

PDT is also viable without the internalization, since the lifetime of produced ROS is too short to diffuse and attack adjacent healthy cells, hence only those targeted cells will be damaged. A copolymer of boron-dipyrromethene (BODIPY) and fluorene structural units (PBF, Scheme 24.3, structure 22) can be used to form CPNs with 3,3’-dithiodipropionic acid (SDPA) through electrostatic interaction and these nanoparticles were applied in ROS generation and the killing of tumor cells. PBF served as a PS individually without the addition of other molecules to generate ROS upon exposure to white light.151 Transcriptional activator protein (TAT) and anti-EpCAM were modified on PFT (Scheme 24.3, structure 23)/PS NPs. Both quick internalization and cellular surface targeting on EpCAM overexpress cells can be achieved for further PDT procedures.152 A convenient approach for theranostics of cancer in vivo was also demonstrated using a photosensitizer comprised of tetraphenylporphyrin (TPP) doped PFBT (Scheme 24.3, structure 24) nanoparticles (TPP-doped Pdots). The fluorescence of PFBT was completely quenched by the photosensitizer, yielding an energy transfer efficiency of nearly 100% and singlet oxygen generation quantum yield of around 50%. The in vivo fluorescent imaging and PDT can be simultaneously achieved through this TPP doped PFBT nanoparticle153 (Figure 24.23). A novel and innovative combination of bacteria vector delivery of toxin and PDT was developed with through the use of cationic poly(fluorene-vinylene-benzothiadiazole) (PFVB, Scheme 24.3, structure 25). PE66 toxin-expressing plasmid pET28a-PE66 was installed into the expression vector BL21-PE66. PFVB coated E. coli [BL21-PE66] was obtained by taking advantage of electrostatic and hydrophobic interactions. The loaded toxin PE66 can be released from the capsules by the collaboration of red-emissive PFVB and a membrane disrupting antibiotic polymyxin B (PLB) so to exert their toxic functions on cancer cells and initiate programmed cell apoptosis. Furthermore, tumor cells can uptake PFVB coated BL21-GFP, a green fluorescent protein (GFP) expressive E. coli, to fabricate the bacterial vectors. After endocytosis, cell death was observed under white light with an intensity of 54 J cm–2 (Figure 24.24).154 As mentioned previously, under most situations, a CP/CO exhibits short wavelength excitations, which can less effectively penetrate in tissue, leading to a severe restriction for in vivo applications. Although the application of two-photon absorbing photosensitizers and red to infrared light absorbing materials can reduce this shortcoming, the development of these materials with highly efficient

FIGURE 24.22  Schematic mechanism of PTDP for selective targeting and inactivation of intracellular estrogen signal pathway protein. (From Wang, B.; Yuan, H.; Zhu, C.; Yang, Q.; Lv, F.; Liu, L.; Wang, S. Sci. Rep. 2012, 2, 766. With permission.)

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.23  (A) In vivo fluorescence imaging by TPP-doped Pdots. Bright-field and fluorescence images of whole-animal imaging for tumor-bearing mice with intravenous injection (upper) or intratumoral injection of TPP-Pdots (lower). (B) Ex vivo bright-field and fluorescence images of different mouse organs including tumor, liver, spleen, kidney, lung, heart. The upper panel showed the control group without Pdot administration, the middle panel showed the organs collected from mice with intravenous injection, and the lower panel showed the organs collected from mice with intratumoral injection. I.V. = intravenous injection, I.T. = intratumoral injection. (C) Representative photographs of tumor-bearing mice on the 28th day after various treatments. (upper) The control mice with light irradiation only. (middle) The mice treated with I.V. injection of TPP-Pdot and light irradiation. (lower) The mice treated with I.T. injection of TPP-Pdot and light irradiation. (D) Representative photographs of tumors collected from the mice after various treatments. (E) Tumor growth curves in the tumor-bearing mice after different treatment. Tumor volumes were normalized to their initial sizes. Error bars represent the standard deviations of eight mice per group. (F) Histological haemotoxylin and eosin (H&E) staining of various organs from the mice of control group, I. V. group, and I. T. group. For both the intravenous injection group and intratumoral injection group, no pathological changes were observed in the spleen, kidney, lung, and heart, as compared to control group. Minor liver damage was found in the intravenous injection group, such as hepatocellular necrosis, ballooning degeneration, congestion, and inflammatory cell infiltration. (From Li, S.; Chang, K.; Sun, K.; Tang, Y.; Cui, N.; Wang, Y.; Qin, W.; Xu, H.; Wu, C. ACS Appl. Mater. Interfaces 2016, 8 (6), 3624. With permission.)


Conjugated Polymers

FIGURE 24.24  (A) Schematic diagram of cationic PFVB coated live bacteria for controlled drug (protein) loading and release for multimodal extracellular (upper) and intracellular (lower) killing of cancer cells. (B) The endocytosis of PFVB coated bacteria vectors by A498 cells and their location inside the cells. (a) DIC image, (b) fluorescence image for GFP from [BL21-GFP] that collects the signals from 500 − 550nm, (c) fluorescence image for CP that collects the signals from 580 − 680 nm, (d) merged 3D scanning image of A498 cells. (C) Cell viability analysis of CP coated E. coli [BL21-PE66] against A498 cells upon incubating for and 3 days. (D) Cell viability analysis of CP coated E. coli [BL21] against A498 cells under white light. Standard deviations are shown as error bars from five parallel experiments. (From Zhu, C.; Yang, Q.; Lv, F.; Liu, L.; Wang, S. Adv. Mater. 2013, 25 (8), 1203. With permission.)

absorption has proven to be a challenge.155 Recently, the bioluminescence resonance energy transfer (BRET) process has been developed to circumvent the need of light irradiation from an outside source. In situ bioluminescence of luminol can be absorbed by a cationic oligo(p-phenylene vinylene) (OPV, Scheme 24.2, structure 15) that acts as the photosensitizer for PDT. Luminol emits a striking blue bioluminescence with a maximum of 425 nm in the presence of oxidizing agent (such as hydrogen peroxide) and horseradish peroxidase (HRP). The bioluminescence can excite OPV, which then sensitizes surrounded oxygen to produce 1O2. Using this BRET system as its excitation source, OPV can achieve more than 90% cell inhibition of HeLa cells. Moreover, for the treatment of HeLa cell tumor-bearing nude mice using intratumoral injection, without light irradiation, the synergistic interaction of OPV and the BRET system showed obvious inhibition of tumor growth (Figure 24.25).156

24.3.4 Photothermal Therapy Photothermal therapy (PTT) is yet another approach for cancer treatment. PTT is based on the concept that specific chromophores can effectively convert absorbed light into heat in order to kill cancer cells. Gold and carbon nanomaterials, transition metal dichalcogenides nanomaterials are commonly used for PTT.157–163 Although these materials overcome the weakness of poor penetration of short wavelength light, low efficiency, weak photostability, and non-biodegradation have motivated the development of alternative photothermal agents. Through altering the donor−acceptor (D–A) construction units, conjugated polymer dots (Pdots, CPNs with a relatively small size) can extend absorption toward the NIR region of the spectrum. Moreover, taking advantage of biocompatible, noninvasive, and spatiotemporalcontrolling modes and water dispersible Pdots with NIR-induced photoactivity, CPs exhibit promising potential in the application of PTT. As an example, consider the diketopyrrolopyrrole (DPP)-containing polythiophene PTD (Scheme 24.3, structure 26). PTD displays strong NIR absorbance and high photothermal conversion efficiency. Under

Conjugated poly/oligo-electrolytes for cancer diagnosis and therapy


FIGURE 24.25  A Schematic Illustration of the BRET System for PDT. B Cell viability of HeLa cells after incubation with OPV in the absence and presence of luminol luminescence system (E + S). Values are expressed as means ± SD (n = 3, P PPy– SO4>PEDOT–SO4>PEDOT–pTS with a 45 s deposition of PEDOT–pTS having the lowest impedance. The electrode was placed into a rat inferior colliculus (IC) to record neural activity induced by white noise bursts. Correlations of background root-mean-square (RMS), signal to noise ratio and mean spike count with impedance at 1 kHz were seen across different conducting polymers but not with varying PEDOT-pTS thickness. Local field potentials (LFP) are an electrophysiological signal obtained from a large population of neurons and can be used to control brain machine interfaces [107]. An ionic liquid doped conducting polymer, PEDOT-EMIIM was used to coat a microelectrode array. The electrodes were placed in the primary sensory cortex or barrel cortex of a rat and were able to detect LFPs [108].

Biomedical Applications of Organic Conducting Polymers


FIGURE 25.5  Sample images of PC12 neurite outgrowth on PEDOT at 96 h post-plating with bare polymer (top) and adsorbed whole laminin coated polymer (bottom): (A) PEDOT-pTS; (B) PEDOT/DEDEDYFQRYLI; (C). PEDOT/DCDPGYIGSR. Reproduced with permission from Green et al., Biomaterials 2009; 30(22): 3637–3644.

FIGURE 25.6  Neuronal growth on PEDOT:PSS-co-MA. (A) Contrast phase image of dissociated neural cells plated at 25,000 cells/cm2 and cultured 24 h on the non-functionalised polymer. (B) Cells cultured for 24 h on the polymer with covalently bonded PLL. (C) Fluorescence microscopy image of cells cultured 5 days on the polymer with covalently bonded PLL, and then fixed and processed for MAP2 and Tau immunocytochemistry. The neuronal somas and dendrites appear because they were intensely positive for MAP2, and also showed moderate staining for Tau. However, note that axons were positive only for Tau. (D) Cells cultured in the same conditions as in C, with the difference that the polymer was functionalised by electroadsorption of PLL. Scale bars: 100 μm. Reproduced with permission from Collazos-Caastro et al., Biomaterials 2013; 34(14): 3603–3617.

FIGURE 25.7  (a) Optical micrograph of a conducting polymer modified electrode array. The labels (1–4) represent four different coatings, enabling a statistical analysis of each coating within a single experiment. One uncoated electrode is also labelled. In this example, 1–4 are 15, 30, 45 and 60 s deposition times of PEDOT–pTS. (b) Streaming data measured at each electrode with 70 dB white noise bursts measured in the IC. (c) Signal to noise ratio during insertion and retraction of the electrode array into the IC. 70 dB white noise at representative uncoated (dashed) and conducting polymer coated (solid) electrodes and (d) different sound pressure levels (40–70 dB) on a conducting polymer coated electrode. Reproduced with permission from Harris et al., Advanced Functional Materials 2018, 28(12): 1700587.

794 Conjugated Polymers

Biomedical Applications of Organic Conducting Polymers


Electrodes that penetrate the brain surface can induce tissue damage, reducing the electrode performance. Electrodes that lie on the surface of the brain (termed electrocorticography) are less invasive. Gold electrodes on a flexible parylene C substrate were coated with PEDOT–PSS and were able to detect neural activity from the surface of a rat somatosensory cortex [109]. A flexible parylene NeuroGrid coated with PEDOT–PSS was able to measure local field potentials and neural spiking across the dorsal cortex of a rat [110]. Cells are able to communicate using chemicals as well as electrical cues. Neurotransmitters include dopamine, glutamate, γ-aminobutyric acid (GABA) and serotonin. Neurotransmitters can be detected on an electrode through a redox reaction. Metal electrodes display no selectivity to these neurotransmitters, so accurate concentration measurements cannot be obtained. OCPs are capable of detecting neurotransmitters. Chromaffin cells on PEDOT–PSS modified electrodes underwent exocytotic processes when gently pressed with a glass pipette. Amperometric spikes due to release of catecholamines were detected while holding the electrode potential at 700 mV [111]. Overoxidised PPy-PBS, obtained by applying a potential of 989 mV for over 40 minutes, was coated with Nafion for dopamine detection or Nafion and then glutamate oxidase for glutamate detection [112]. Highly selective measurements of dopamine and glutamate could then be achieved by amperometry. Organic electronic ion pumps (OEIP) use an OCP to electrophoretically pump neurotransmitters through a permselective membrane, enabling high spatiotemporal delivery resolution, without necessitating liquid flow. Delivery of GABA from a reservoir could be controlled by a PEDOT–PSS electrode that also functioned as a recording electrode [113, 114]. The GABA reduced epileptiform activity in mouse hippocampal slices.

25.4 Implantable Electrodes for Electrical Stimulation Charge can be injected from an electrode into tissue, inducing an action potential in neurons or muscle contraction. Stimulating electrodes can provide sensory cues, such as for the cochlear implant and bionic eye, or for controlling tremor in Parkinson’s disease or symptoms of other neurological disorders. Platinum or platinum–iridium is normally used for the stimulating electrodes, but the current passed through the electrode can result in corrosion. Small electrodes are required to target individual cells, resulting in high charge densities at the electrode surface that can enhance electrode damage. OCPs can increase the charge injection capacity of an electrode, allowing smaller electrode sizes. The effective electrode area and charge density of conducting polymer modified electrodes has been measured by cyclic voltammetry and optical microscopy [49, 78, 115]. The charge density was a function of conducting polymer type and electrode area, offering the possibility of increasing charge density of neural electrodes by altering both the electrode material and geometry. Charge passed from an electrode into tissue can also alter cell growth, proliferation and differentiation. PC12 cells grown on PPy were subjected to electrical stimulation. The electrical stimulation induced an increase in fibronectin adsorption [116] and in neurite length compared to non-stimulated cells [117]. Electrical stimulation of cells in vivo has also been undertaken. PEDOT–pTS modified electrodes had a larger charge injection capacity compared to bare platinum. The electrodes were implanted into a cat suprachoroidal space [118]. Neural stimulation at 100 µA resulted in a voltage excursion of 1.5 ± 0.2 V versus 3.3 ± 0.6 V for modified and unmodified electrodes. The electrodes could induce neural activity, which was detected by a recording electrode implanted in the visual cortex. PEDOT–PSS modified electrodes on a polyimide substrate were inserted into a rat cochlear nucleus and able to induce electrically evoked auditory brainstem responses (eABRs) [119]. Photostimulation of neurons was achieved by interfacing poly(3-hexylthiophene-2,5-diyl) with ­phenyl-C61-butyric-acid-methyl ester (rr-P3HT:PCBM). The rr-P3HT acts as an electron donor material while PCBM is an electron acceptor. Light focussed on the polymer could induce activity in nearby neurons [120]. Subsequently, a fully organic retinal prosthesis has been developed (Figure 25.8) [121].


Conjugated Polymers

FIGURE 25.8  (a) Scanning electron microscopy images of the full prosthetic device (top) and of its cross-­section at higher magnification showing the three-layered structure (bottom). (b) Scheme of the subretinal implant strategy. RPE, retinal pigment epithelium; ONL, outer nuclear layer; INL, inner nuclear layer; GCL, ganglion cell layer. (c) Sample confocal scanning laser ophthalmoscopy image of the surgical prosthesis placement in the eye fundus of a dystrophic RCS rat. (d) OCT analysis showing the strict contact between the retina (arrows) and the implant (arrowheads) at 30 and 180 dots per inch (DPI). No retinal detachments or breakages were observed. (e)–(f) Explanted eye fixed (e), stained with bisbenzimide and acquired by confocal microscopy (f) to identify retinal nuclear layers and the position of the device. The high-magnification image Inset shows the integrity and location of the implant in the retina. Reproduced with permission from Maya-Vetencourt et al., Nature Materials 2017; 16: 681–689.

Biomedical Applications of Organic Conducting Polymers


It is composed of three layers: a silk fibroin substrate coated with PEDOT–PSS and poly (3-hexylthiophene) (P3HT). When attached to the retina of dystrophic RSC rats (a model of retinitis pigmentosa), electrophysiological and behavioural analyses indicated recovery of visual acuity for up to ten months post-surgery. Neurotransmitters can be eluted from a conducting polymer to alter cell behaviour. Electrical stimulation of overoxidised PEDOT soaked in GABA, glutamate or aspartate could induce a neural response [122]. 2-amino-5-phosphonopentanoate (AP5) or 6-cyano-7-nitroquinoxaline-2,3-dione (CNQX) have been used as dopants in PPy. These compounds act as inhibitors of NMDA and AMPA-type glutamate receptors. Reduction of PPy ejected the dopant ions and blocked neural activity [123]. An electronic ion pump has also been used to deliver glutamate in the cochlea, leading to a shift in the auditory brainstem response [124].

25.5 Controlled Delivery Systems OCPs are capable of electrically controlled, in situ delivery of bioactive molecules. This, coupled with their excellent electrochemical properties, enables a new paradigm in neural interface. Loading of bioactive molecules can be readily achieved via various mechanisms, leading to a variety of electrically on-demand delivery systems [125]. For example, an anionic molecule can be doped directly into an OCP during polymerisation and released in an electrically tunable manner by controlling the redox state of the polymer [69, 85, 126]. Loading of a cationic molecule is achievable via a reduction process using a pre-formed OCP bearing a large anionic dopant as the drug carrier [127, 128]. Alternatively, other mechanisms can be explored for loading of bioactive molecules, including neutral molecules, such as via hydrophobic interactions [129] and affinity binding [130]. Herein we discuss some examples of OCP systems for controlled delivery of drugs or growth factors. We also highlight the underlying principles for engineering the electrically-controlled release capabilities of these systems.

25.5.1 Controlled Drug Delivery Systems A number of studies undertaken have involved the use of dexamethasone 21-phosphate disodium (Dex-P), an anionic corticosteroid prodrug, as a dopant for preparation of anti-inflammatory OCP coatings for neural electrodes [69, 85, 131, 132]. The underlying biological rationale is to minimise reactive cellular and tissue reactions to the electrode implants, the biological processes that give rise to fibrotic or scar tissue encapsulation of implanted electrode, rendering the electrode-neural communication ineffective. The early work by Wadhwa et al. demonstrated a proof-of-concept in vitro [85]. The PPy coatings showed approximately zero-order drug release of Dex-P as a function of cyclic potential stimulation cycle (up to 30 stimulation cycles), and the dosage released was sufficient for reducing murine glial cells, yet non-toxic to primary murine neurons. To improve drug loading and release efficiency, nanoporous PPy-Dex-P films were developed via a template-assisted approach using self-assembled polystyrene nanobeads [133], to produce a nanoporous structure. Further, expanding the application of this approach, the nanoporous film was loaded with extra drug to maximise the loading capacity, or a different drug for simultaneous release of multiple drugs. In both cases, the drug-loaded structure was sealed with a thin layer of PPy to retain the initial electrically controlled release characteristics [131]. In a separate approach, multi-walled carbon nanotubes (CNTs) were exploited as nanoreservoirs of Dex-P for encapsulation into PPy coatings via electropolymerisation [132]. The resultant PPy coating showed improved drug loading capacity and a more linear and sustainable release profile. This approach also results in improved electrical conductivity of the electrode, attributable to CNT’s inclusion, which makes it very attractive for neural interface applications. The effect of anti-inflammatory OCP coatings on electrode performance has been assessed in vivo [134]. Microelectrodes were coated with PEDOT doped with CNT–Dex-P reservoirs and implanted in adult rats for stimulation in the dorsal root ganglion over a period of two weeks. Compared to the non-coated


Conjugated Polymers

electrodes, the coated electrodes showed lower in vivo impedance characteristics and provoked significantly less neuronal death/damage at the site of implantation. More recently, Boehler et al. examined the chronic effect of local pharmaceutical intervention on electrode performance and electrode-tissue integration [135]. Neural microelectrodes were coated with PEDOT/Dex-P and implanted in the rat hippocampus over a period of 12 weeks, during which drug release was activated on a weekly basis at the OCP-coated electrode sites. The electrodes exposed to Dex-P intervention featured stable recordings and impedance characteristics over the entire course of study. Although no significant difference in foreign body response was observed among the implant electrodes following 12 weeks’ implantation, more close-by neurons were observed at the sites of Dex-P functionalised electrodes, suggesting improved electrode–neural interaction due to on-site pharmaceutical intervention. Electrically on-demand drug delivery systems require an electrical stimulus to activate drug release. This feature, once coupled with an appropriately designed diagnostic system, enables controlled drug release to be triggered by the onset of a specific disease. Our group has been working towards development of implantable seizure-initiated drug delivery systems [136]. Figure 25.9 illustrates the key elements of the implantable system, which encompasses electrodes with OCP coatings that are loaded with anti-epileptic drugs (AEDs) and electronics that are capable of recording human electrocorticographic (ECoG) activity, detecting the features of epileptic seizures, and translating the features into electrical stimuli to modulate AEDs release. The potential of using ECoG recordings and stimulation hardware for seizure-initiated drug release has been demonstrated in vitro (Figure 25.10b). Fosphenytoin sodium (FOS), an AED prodrug, was selected as a dopant for preparation of OCP–AED electrode coatings. AED release was initiated using pre-recorded human ECoG data via a custom ECoG hardware system. A modified HPLC system was developed for in situ quantification of the triggered drug release. Our results showed that the AED delivery can be initiated within 10 s of electrical stimulation, a rapid response of clinical relevance in seizure control, and the drug dose can be readily regulated by controlling the stimulation current (i.e. the amount of charge injected into the OCP coatings).

25.5.2 Controlled Delivery of Growth Factors To achieve a stable and effective neural interface in chronic settings, it is critical to ensure intimate interaction between the implanted electrode and target neurons. One promising approach is to provide controlled and sustained neurotrophic support at the site of implanted electrodes to improve neuron survival and neurite outgrowth [2, 4]. Our group has developed a number of neurotrophin-OCP coatings to improve the nerve–electrode interface of the cochlear implant [66, 67, 137]. These coatings were prepared via a two-layer approach where a commonly used anionic dopant is employed as a co-dopant to provide excellent electrochemical and mechanical properties that facilitate subsequent inclusion of neurotrophic factors. For example, the first layer can be grown from pyrrole and pTS, upon which a second layer is deposited from pyrrole, pTS and neurotrophin-3 (NT3) [66, 138]. The OCP coatings have demonstrated accelerated release of neurotrophins when subjected to a clinically relevant biphasic electrical stimulation [139]. The biological activities of the neurotrophin-OCP coatings were assessed both in vitro [138] and in vivo [139]. SGN explants from rat cochleae were cultured on PPy-pTS-NT3 coated electrodes and showed improved neurite outgrowth, in comparison with the explants cultured on PPy-pTS coated electrode. Electrical stimulation further enhanced neurite outgrowth from the explants cultured on PPy-pTS-NT3 [138]. Furthermore, we have shown a synergistic effect of delivering two neurotrophins simultaneously to encourage neuron survival and neurite elongation [137]. This provides insight into the rational design of OCP coatings for improved efficacy of local neurotrophin intervention. Electrical stimulation of the cochlear neural explants cultured on PPy containing both NT-3 and BDNF produced most pronounced improvement in neurite outgrowth, compared to PPy doped only with NT-3 or BNDF (Figure 25.10). In vivo studies were undertaken in the cochlea of deafened guinea pigs, to assess the effects of electrical stimulation and OCP coating for a period of two weeks [139]. PPy-pTS-NT3 coated electrodes in

Biomedical Applications of Organic Conducting Polymers


FIGURE 25.9  (A) Schematic of the seizure-initiated drug delivery system showing: (1) the onset of pre-seizures; (2) progression into an epileptic seizure event; (3 and 4) surgical intervention where by an AED loaded seizure detection and delivery device is implanted; (5) the implant detects the ECoG per-seizure and subsequently delivers the encapsulated AED (depicted as green molecules); and (6) due to the delivery of the AED at a therapeutic dosage the ECoG activity returns to normal. (B) Schematic of signal flow graph and corresponding interface signals of the ECoG system for three programmable output currents. The charge (Q) injected into the PPy–FOS during the applied current stimulation is shown in brackets. Adapted with permission from Muller et al., Sensors and Actuators B: Chemical 2016; 236: 732–740.


Conjugated Polymers

FIGURE 25.10  (A) Representative images of cochlear neural explants grown on PPy–pTS polymers with/without neurotrophin, following four days of explant culture. Neurites were visualised by immunocytochemistry with a neurofilament-200 primary antibody and a fluorescent secondary antibody (green). Cell nuclei are labelled with DAPI (blue). In the absence of neurotrophin (PPy–pTS), very few neurites were observed from explants, while explants grown on PPy–pTS containing neurotrophin demonstrated increased numbers of sprouting neurites. A greater number of neurites per explant were observed on explants grown on the electrically stimulated PPy films. Scale bar is 200μm for images. (B) Neural response to neurotrophins released from PPy-pTS under stimulated (Stim) and unstimulated (Unstim) conditions. Neurite extension was significantly increased (p