Non-traditional architectures for semiconducting polymers

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NON-TRADITIONAL ARCHITECTURES FOR SEMICONDUCTING POLYMERS

by

Elizabeth L. (Betsy) Melenbrink

A Dissertation Presented to the
FACULTY OF THE USC GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(CHEMISTRY)

December 2018

Copyright 2018 Elizabeth L. Melenbrink
ii

Epigraph

Never regard study as a duty but as an enviable opportunity to learn…
- Albert Einstein

iii

Dedication

To everyone who believed in me when I could not believe in myself…

iv

Acknowledgements
I never really imagined this journey would come to an end, and yet somehow, here we
are. It takes a village to create a PhD dissertation, and I have a lot of people to thank for
their help along the way, whether professional, personal, or both.
First, I must thank Prof. Barry Thompson for serving as my research advisor during my
5 years at USC. His support enabled me to confront and overcome my fear of organic
synthesis. And though I may have complained bitterly, I am grateful for the challenging
projects he has assigned me during my graduate studies. I know I would have been bored
and unstimulated if I had been handed a successful project from the beginning.
I must also thank the faculty who served on my committee: Prof. Smaranda Marinescu
and Prof. Ralf Haiges for serving on my screening and qualifying exam committees and
Prof. Mark Thompson and Prof. Jongseung Yoon for serving on my screening, qualifying,
and defense committees. Prof. Marinescu has also earned my gratitude for serving as the
faculty mentor for Women in Chemistry for the duration of my time at USC, in the course
of which she was a fount of good ideas and encouragement. Prof. Haiges seems to
understand how everything works, and he taught me not only how to set up, repair, and
maintain a glovebox, but also how to safely work with organic azides. For this I am forever
grateful. Prof. M. Thompson deserves special thanks for letting his lab be my second home
on campus, from my very first lab rotation, to the many OPV subgroup meetings I attended,
to the chemicals, equipment, and expertise I frequently borrowed from his students.
Each chapter of this dissertation was helped by various people to reach its current state.
For Chapters 2, 3, and 4, mechanical measurements were performed by the researchers in
Prof. Lipomi’s group at UC-San Diego. For Chapters 3 and 4, DSC measurements were
v

performed by Negar Kazerouni. Sanket Samal performed SCLC mobility measurements
for Chapters 2, 3, and 4. John Luke McConn and Başak Perçin contributed to synthesis
for several projects. Thomas Saal and Amanda Baxter helped me with reaction monitoring
via IR for Chapter 5. Dr. Jeff Celaje and Antonina Nazirova lent me reagents and shared
their expertise on performing copper-catalyzed “click” reactions for Chapter 5. Abegail
Tadle helped me perform all fluorescence spectroscopy in Chapter 5. Prof. Haiges tested
the explosive nature of the polymer intermediates synthesized in Chapter 5 of this
dissertation.
Next, I would like to acknowledge my lab mates, as one’s coworkers make all the
difference. I am grateful to have learned from the calm demeanors of Dr. Sangtaik Noh and
Alia Latif and the practical approach of Dr. Bing Xu. Dr. Jenna Howard is a pillar of
strength and I appreciate her example of determination and drive. Dr. Şeyma Ekiz holds a
special place in my heart after three years of sharing an office, defeats and triumphs, good
days and bad, and conversations that covered every topic you could imagine. Thank you to
Jenna and Şeyma for your advice, support, and friendship!
I am grateful for everyone I’ve had the privilege to work with through Women in
Chemistry and all the great women I have met in grad school who have inspired me,
mentored me, cheered me, and encouraged me – Dr. Lena Hoober-Burkhart, Dr. Nadia
Korovina, Arunika Ekanayake, and many more. I was lucky to find such great friends in
my cohort: Dr. Courtney Downes, whose strength and determination I admire, soon-to-be
Dr. Caitlin DeAngelo who has an unmatched passion for teaching and an inspiring
confidence, and soon-enough-to-be Dr. Amanda Baxter, whose curiosity and passion for
learning will take her on wonderful adventures.
vi

I am thankful for my hiking buddies, who walked with me on the path to sanity, one
step at a time: Caitlin, Amanda, Eric, Andrew, Sarah, and John.
The past 5 years would have been even more trying had it not been for my friends who
encouraged me when I called with a voice full of tears and frustration – Zuri, Cat, Melissa,
Sarah, and Kelly. Thank you for your constancy and for reaching out when I couldn’t.
I owe a debt of gratitude my undergraduate research advisor, Prof. Wei You, for
introducing me to the joys of research and planting the seed of grad school in my mind.
Gratitude also goes to the incomparable Michele Dea, without whom the chemistry
department would crumble.
Finally, my deepest and most sincere thanks is owed to my family, Eric, Deniese, and
Nathan Melenbrink, who instilled passion and curiosity in me, and supported me
unwaveringly through a rough 5 years.

vii

Table of Contents
Epigraph .............................................................................................................................. ii
Dedication .......................................................................................................................... iii
Acknowledgements ............................................................................................................ iv
Table of Contents .............................................................................................................. vii
List of Schemes ................................................................................................................ xiv
List of Tables ................................................................................................................... xvi
List of Figures .................................................................................................................. xxi
Abstract ........................................................................................................................... xlvi
Chapter 1: Non-Traditional Semiconducting Polymer Architectures ................................. 1
1.1 Background/Motivation ....................................................................................... 1
1.1.1 Benefits of organics vs. inorganics ............................................................... 1
1.1.2 Benefits of fully conjugated polymers vs. small molecules ......................... 1
1.1.3 Drawbacks of fully conjugated polymers ..................................................... 4
1.2 Alternative Architectures ..................................................................................... 6
1.3 Discrete Conjugation ............................................................................................ 9
1.3.1 Motivation/Background ................................................................................ 9
1.3.2 Oligothiophene Systems ............................................................................. 10
1.3.3 Other Studies Using Simple Electroactive Units ........................................ 12
1.3.4 Modern Conjugation-Break Spacers ........................................................... 13
viii

1.4 Pendant Electroactive Polymers ......................................................................... 20
1.4.1 Motivation/Background/History ................................................................. 20
1.4.2 Simple Electroactive Pendants .................................................................... 22
1.4.3 Metathesis Polymerizations and Memory Polymers................................... 23
1.4.4 Polymers with High Chromophore Density ................................................ 26
1.5 Conclusion .......................................................................................................... 32
1.6 References .......................................................................................................... 33
Chapter 2: Influence of Systematic Incorporation of Conjugation-Break Spacers into Semi-
Random Polymers on Mechanical and Electronic Properties ........................................... 47
2.1 Introduction ........................................................................................................ 47
2.2 Results and discussion ........................................................................................ 50
2.2.1 Optical properties ........................................................................................ 52
2.2.2 Structural properties .................................................................................... 54
2.2.3 Electronic properties ................................................................................... 56
2.2.4 Mechanical properties ................................................................................. 57
2.3 Conclusions ........................................................................................................ 65
2.4 References .......................................................................................................... 67
Chapter 3: Influence of Acceptor Side-Chain Length on Mechanical and Electronic
Properties of Semi-Random Polymers with Conjugation-Break Spacers......................... 72
3.1 Introduction ........................................................................................................ 72
ix

3.2 Results and discussion ........................................................................................ 74
3.2.1 Optical properties ........................................................................................ 76
3.2.2 Structural Properties.................................................................................... 77
3.2.3 Electronic Properties ................................................................................... 79
3.2.4 Mechanical properties ................................................................................. 81
3.3 Conclusions ........................................................................................................ 84
3.4 References .......................................................................................................... 85
Chapter 4: Effects of Modulating Content of Conjugation-Break Spacers in Semi-Random
Polymers on Mechanical and Electronic Properties ......................................................... 88
4.1 Introduction ........................................................................................................ 88
4.2 Results and discussion ........................................................................................ 90
4.2.1 Optical Properties........................................................................................ 92
4.2.2 Structural Properties.................................................................................... 94
4.2.3 Electronic Properties ................................................................................... 95
4.2.4 Mechanical Properties ................................................................................. 96
4.3 Conclusions ........................................................................................................ 98
4.4 References ........................................................................................................ 100
Chapter 5: Diketopyrrolopyrrole (DPP) pendant polymers ............................................ 102
5.1 Motivation ........................................................................................................ 102
5.1.1 Motivation for Pendant Polymers ............................................................. 102
x

5.1.2 Rationale for DPP ..................................................................................... 105
5.2 Background ...................................................................................................... 106
5.2.1 Direct polymerization vs. post-polymerization modification ................... 106
5.2.2 Stereoregularity effects ............................................................................. 108
5.3 Thesis ............................................................................................................... 109
5.4 Results and Discussion ..................................................................................... 109
5.4.1 Synthesis ................................................................................................... 109
5.4.2 Optical Properties...................................................................................... 117
5.4.3 Structural Properties.................................................................................. 119
5.4.4 Electronic properties ................................................................................. 120
5.5 Summary and Conclusion ................................................................................ 121
5.6 References ........................................................................................................ 122
Appendix A: Influence of Systematic Incorporation of Conjugation-Break Spacers into
Semi-Random Polymers on Mechanical and Electronic Properties ............................... 126
A.1 Materials and Methods ..................................................................................... 126
A.1.1 SCLC Device Fabrication and Characterization ....................................... 128
A.1.2 Preparation of Substrates for Mechanical Measurements ......................... 128
A.1.3 Preparation of Films for Mechanical Measurements ................................ 129
A.1.4 Preparation of PDMS Elastomers ............................................................. 129
A.1.5 Buckling-Based Metrology for Measuring Elastic Moduli ...................... 130
xi

A.1.6 Tensile Testing of Pseudo–Freestanding Films ........................................ 131
A.2 Synthetic Procedures ........................................................................................ 132
A.3 Synthetic data ................................................................................................... 133
A.4 Polymer Nuclear Magnetic Resonance Spectra ............................................... 135
A.5 Film UV-Vis Absorption Spectra ..................................................................... 140
A.6 Cyclic Voltammograms.................................................................................... 151
A.7 Differential Scanning Calorimetry ................................................................... 156
A.8 GIXRD Patterns ............................................................................................... 162
A.9 Mobility Measurements.................................................................................... 165
A.10 Film-on-Elastomer Mechanical Data............................................................ 169
A.11 References .................................................................................................... 176
Appendix B: Influence of Acceptor Side-Chain Length on Mechanical and Electronic
Properties of Semi-Random Polymers with Conjugation-Break Spacers....................... 177
B.1 Materials and Methods ..................................................................................... 177
B.1.1 SCLC Device Fabrication and Characterization ....................................... 178
B.1.2 Preparation of Substrates for Mechanical Measurements ......................... 179
B.1.3 Preparation of Films for Mechanical Measurements ................................ 180
B.1.4 Tensile Testing of Pseudo–Freestanding Films ........................................ 180
B.2 Synthetic Procedures ........................................................................................ 180
B.3 Synthetic data ................................................................................................... 182
xii

B.4 Polymer Nuclear Magnetic Resonance Spectra ............................................... 184
B.5 Film UV-Vis Absorption Spectra ..................................................................... 188
B.6 Cyclic Voltammograms.................................................................................... 199
B.7 Differential Scanning Calorimetry ................................................................... 204
B.8 GIXRD Patterns ............................................................................................... 212
B.9 Mobility Measurements.................................................................................... 215
B.10 Film-on-Water Mechanical Data .................................................................. 218
B.11 References .................................................................................................... 223
Appendix C: Effects of Modulating Content of Conjugation-Break Spacers in Semi-
Random Polymers on Mechanical and Electronic Properties ......................................... 224
C.1 Materials and Methods ..................................................................................... 224
C.1.1 SCLC Device Fabrication and Characterization ....................................... 225
C.1.2 Preparation of Substrates for Mechanical Measurements ......................... 226
C.1.3 Preparation of Films for Mechanical Measurements ................................ 227
C.1.4 Tensile Testing of Pseudo–Freestanding Films ........................................ 227
C.2 Synthetic Procedures ........................................................................................ 227
C.3 Synthetic data ................................................................................................... 229
C.4 Small Molecule Nuclear Magnetic Resonance Spectra ................................... 231
C.5 Polymer Nuclear Magnetic Resonance Spectra ............................................... 233
C.6 Film UV-Vis Absorption Spectra ..................................................................... 235
xiii

C.7 Cyclic Voltammograms.................................................................................... 239
C.8 Differential Scanning Calorimetry ................................................................... 241
C.9 GIXRD Patterns ............................................................................................... 246
C.10 Mobility Measurements ................................................................................ 249
C.11 Film-on-Water Mechanical Data .................................................................. 252
C.12 References .................................................................................................... 257
Appendix D: Diketopyrrolopyrrole (DPP) pendant polymers ........................................ 258
D.1 Materials and Methods ..................................................................................... 258
D.2 Synthesis........................................................................................................... 259
D.3 Nuclear Magnetic Resonance Spectroscopy .................................................... 269
D.4 Infrared Spectroscopy ...................................................................................... 300
D.5 UV-Visible Spectroscopy ................................................................................. 302
D.6 Fluorescence Spectroscopy .............................................................................. 308
D.7 Differential Scanning Calorimetry (DSC)........................................................ 313
D.8 X-ray Diffraction .............................................................................................. 317
D.9 Cyclic Voltammetry ......................................................................................... 319
D.10 References .................................................................................................... 320
Biographical Sketch ........................................................................................................ 321

xiv

List of Schemes
Scheme 2.1. Semi-random polymers with conjugation-break spacers synthesized for this
study .................................................................................................................................. 50
Scheme 3.1. Semi-random polymers with conjugation-break spacers synthesized for this
study .................................................................................................................................. 73
Scheme 4.1. Semi-random polymers with conjugation-break spacers synthesized for this
study .................................................................................................................................. 90
Scheme 5.1. Polymerization and post-polymerization modification of 4-methylstyrene. i)
AIBN, toluene, 90 °C, 2 hrs (XI-a); MMAO/Zr(IV) catalyst, toluene, RT, 2 days (XI-i);
MMAO/Ti(IV) catalyst, toluene, RT, 2 days (XI-s); ii) AIBN, NBS, CCl4, 85 °C, 1 hr (XI-
a and XI-s); iii) NaN3, DMF, 55 °C, 4 days (XI-a and XI-s) ......................................... 110
Scheme 5.2. Copper-catalyzed azide-alkyne “click” reaction. i) CuSO4·5H2O, Sodium
Ascorbate, DMF/H2O/toluene, RT, 3 days ..................................................................... 114
Scheme 5.3. Functionalization, polymerization, and deprotection of ethynyl styrene. i)
PdCl2(PPh3)2, TIPS acetylene, triethylamine, 50 °C, 16 hrs; ii) AIBN, toluene, 70 °C, 10
hrs; iii) TBAF, THF, 0 °C to RT, overnight ................................................................... 115
Scheme 5.4. Copper-catalyzed azide-alkyne “click” reaction. i) CuSO4·5H2O, Sodium
Ascorbate, DMF/H2O, RT, 24 hrs .................................................................................. 116
Scheme D1. Synthesis of protected TIPS ethynylstyrene monomer (II), polymerization of
TIPS ethynylstyrene (III), and deprotection to atactic poly(ethynylstyrene) (IV) ......... 260
Scheme D2. Synthesis of monoalkylated dtdDPP (VI), synthesis of bromobutyl/dtdDPP
(VII), conversion of bromine to azide (VIII) ................................................................. 261
xv

Scheme D3. Copper-catalyzed azide-alkyne “click” reaction to join azidobutyl/dtdDPP
(VIII) with atactic poly(ethynylstyrene) (IV) to yield aPS-alkyne-click-DPP (IX) ...... 263
Scheme D4. Synthesis of atactic, syndiotactic, and isotactic poly(4-methylstyrene) (XI-a,
XI-s, and XI-i), bromination of atactic and syndiotactic poly(4-methylstyrene) (XII-a and
XII-s), and azide conversion to atactic and syndiotactic poly(4-azidomethylstyrene) (XIII-
a and XIII-s) ................................................................................................................... 264
Scheme D5. Synthesis of unsymmetrical propyne/dtdDPP (XIV) ................................ 265
Scheme D6. Copper-catalyzed azide-alkyne “click” reaction to join propyne/dtdDPP
(XIV) with atactic or syndiotactic poly(4-azidomethylstyrene) (XIII-a and XIII-s) to yield
aPS-azide-click-DPP and sPS-azide-click-DPP (XV-a and XV-s) ................................ 266
Scheme D7. Synthesis of dtdDPP reference compound (XVI). ..................................... 267
Scheme D8. Synthesis of "pre-clicked" styrene-based DPP pendant monomer (XVIII).
......................................................................................................................................... 267
Scheme D9. Synthesis of acrylate-based DPP pendant monomer (XX). ....................... 268

xvi

List of Tables
Table 2.1. SEC, thermal, and electronic data for semi-random polymer family ............. 52
Table 2.2. Mechanical properties obtained from film-on-elastomer measurements for
semi-random polymer family ............................................................................................ 59
Table 2.3. Tabulated values of mechanical properties using the film-on-water technique
........................................................................................................................................... 62
Table 3.1. SEC, thermal, and electronic data for semi-random polymer family ............. 75
Table 3.2. Tabulated values of mechanical properties measured using the film-on-water
technique ........................................................................................................................... 82
Table 4.1. SEC, thermal, and electronic data for semi-random polymer family ............. 91
Table 4.2. Tabulated values of mechanical properties measured using the film-on-water
technique ........................................................................................................................... 97
Table 5.1. Synthetic, optical, and electronic data for clicked polymers and small molecule
reference. ......................................................................................................................... 112
Table A1. Number-averaged polymer molecular weights in kDa as measured by SEC 133
Table A2. Polymerization yields from chloroform Soxhlet fraction ............................. 134
Table A3. Optical properties of neat polymers in thin films spin-cast from o-DCB and
placed in a N2 cabinet for 30 minutes. ........................................................................... 148
Table A4. Optical bandgaps of polymer family calculated from absorption band edge.149
Table A5. Solvent effects on optical bandgap for T-6-T subfamily of polymers. ......... 150
Table A6. Highest occupied molecular orbitals of polymer family calculated from
oxidation onset versus ferrocene. .................................................................................... 155
Table A7. Melting points for polymer family obtained from DSC curves. ................... 160
xvii

Table A8. Crystallization points for polymer family obtained from DSC curves. ........ 161
Table A9. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation.
7,8
....... 164
Table A10. Hole mobilities of eight polymers in thin films spin-cast from chloroform.
Results averaged over a minimum of 8 pixels. ............................................................... 166
Table A11. Hole mobilities of eight polymers in thin films spin-cast from o-DCB. Results
averaged over a minimum of 8 pixels. ............................................................................ 166
Table A12. Average hole mobilities for as-cast films. ................................................... 167
Table A13. Average hole mobilities for annealed films. ............................................... 168
Table A14. Elastic moduli (film-on-elastomer) for polymer family, averaged over 3 trials.
......................................................................................................................................... 169
Table A15. Crack onset strain (COS) percentages for polymer family. ........................ 170
Table A16. Film thicknesses for film-on-elastomer measurements (500 rpm | 1000 rpm |
1500 rpm). ....................................................................................................................... 171
Table A17. Mode of failure obtained from micrographs of polymer films, where cracks
perpendicular to strain direction indicate brittle failure and pinholes that do not propagate
indicate ductile failure..................................................................................................... 174
Table A18. Film thicknesses for film-on-water measurements (1000 rpm). ................. 175
Table B1. Number-averaged polymer molecular weights in kDa as measured by SEC.182
Table B2. Polymerization yields from chloroform Soxhlet fraction. ............................. 183
Table B3. Optical properties of neat polymers in thin films spin-cast from chloroform and
placed in a N2 cabinet for 30 minutes. ............................................................................ 196
xviii

Table B4. Optical properties of neat polymers in thin films spin-cast from chloroform and
annealed under N2 at 150 °C for 30 minutes. ................................................................. 196
Table B5. Optical bandgaps of as-cast polymer family calculated from absorption band
edge. ................................................................................................................................ 197
Table B6. Optical bandgaps of annealed polymer family calculated from absorption band
edge. ................................................................................................................................ 198
Table B7. Highest occupied molecular orbitals of polymer family calculated from
oxidation onset versus ferrocene. .................................................................................... 203
Table B8. Melting points for polymer family obtained from DSC curves. ................... 210
Table B9. Crystallization points for polymer family obtained from DSC curves.......... 211
Table B10. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
7,8
for as-cast
films. ............................................................................................................................... 213
Table B11. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
7,8
for
annealed films. ................................................................................................................ 214
Table B12. Hole mobilities of eight CBS polymers and fully conjugated reference polymer
in thin films spin-cast from chloroform and left in a N2 box for 30 minutes. Results
averaged over at least 4 pixels. ....................................................................................... 216
Table B13. Hole mobilities of eight CBS polymers and fully conjugated reference polymer
in thin films spin-cast from chloroform and annealed under N2 at 150 °C for 30 minutes.
Results averaged over at least 4 pixels. .......................................................................... 217
Table B14. Elastic moduli for polymer family. ............................................................. 218
xix

Table B15. Toughness values for polymer family. ........................................................ 219
Table B16. UTS for polymer family. ............................................................................. 220
Table B17. Fracture strength for polymer family. ......................................................... 221
Table B18. Fracture strain for polymer family. ............................................................. 222
Table C1. Number-averaged polymer molecular weights in kDa as measured by SEC.
......................................................................................................................................... 229
Table C2. Polymerization yields from chloroform Soxhlet fraction. ............................ 230
Table C3. Optical properties of neat polymers in thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. ........................... 235
Table C4. Optical properties of neat polymers in thin films spin-cast from o-
dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for 30 minutes. ................ 236
Table C5. Optical bandgaps of unannealed polymer family calculated from absorption
band edge. ....................................................................................................................... 237
Table C6. Optical bandgaps of annealed polymer family calculated from absorption band
edge. ................................................................................................................................ 238
Table C7. Highest occupied molecular orbitals of polymer family calculated from
oxidation onset versus ferrocene. .................................................................................... 240
Table C8. Melting points for polymer family obtained from DSC curves. ................... 245
Table C9. Crystallization points for polymer family obtained from DSC curves. ........ 245
Table C10. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
6,7
for as-cast
films. ............................................................................................................................... 247
xx

Table C11. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
6,7
for
annealed films. ................................................................................................................ 248
Table C12. Hole mobilities of polymers in thin films spin-cast from hot chloroform and
left in a N2 box for 30 minutes. Results averaged over at least 4 pixels. ........................ 250
Table C13. Hole mobilities of polymers in thin films spin-cast from hot chloroform and
annealed under N2 at 150 °C for 30 minutes. Results averaged over at least 4 pixels. .. 251
Table C14. Elastic moduli for polymer family. ............................................................. 252
Table C15. Toughness for polymer family. ................................................................... 253
Table C16. Ultimate tensile strain (UTS) for polymer family. ...................................... 254
Table C17. Fracture strength for polymer family. ......................................................... 255
Table C18. Fracture strain for polymer family. ............................................................. 256
Table D1. Absorption properties of neat polymers in thin films spin-cast from chloroform
and placed in a N2 cabinet for 30 minutes. ..................................................................... 307
Table D2. Emission properties of neat polymers in thin films spin-cast from chloroform
and placed in a N2 cabinet for 30 minutes. ..................................................................... 312
Table D3. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation.
9,10
...... 318

xxi

List of Figures
Figure 1.1 The alternating single and double bonds of conjugated polymers allows for
electron delocalization and charge transport in a way that is not accessible to organic small
molecules. In this simplistic example, a segment of polyacetylene is considered (1.1A). In
1.1B, the filled π-orbitals of the double bonds are depicted, illustrating the concept that
each filled π-orbital is flanked on each side by another filled π-orbital, allowing the
electrons in those orbitals to be delocalized, or not fixed on any particular atom. In 1.1C,
an electron is removed from one π-orbital (shown in white). Its neighbors can quickly fill
the empty space, or hole, leading not only to charge stabilization within conjugated
polymers, but also rapid charge transport along the polymer backbone............................. 2
Figure 1.2. Cartoon of polymer morphology, with dark red crystallites embedded in light
red amorphous segments. The amorphous segments improve the polymer’s mechanical
performance and film integrity while the crystallites are responsible for rapid charge
transport. ............................................................................................................................. 3
Figure 1.3. Some basic polymer architectures. Gradient and block copolymers are typically
linear, and vary from the linear polymer only in their composition. In contrast to branched
polymers, star polymers have many arms radiating from one center. Dendritic polymers
are fractal structures that are grown in generations, where each generation is represented
by a different shade of blue in this figure. Crosslinked or network polymers have chemical
bonds linking polymer chains together. Comb and brush polymers are sometimes known
as graft polymers, where the pendant chains may or may not be of the same composition
as the backbone. .................................................................................................................. 6
xxii

Figure 1.4. Cartoon illustration of a segment of polymer with discrete conjugation, where
the wavy red lines are meant to indicate a flexible non-conjugated segment linking the
yellow, rod-like conjugated segments................................................................................. 9
Figure 1.5. Polymer structures 1-3 alternating oligothiophene segments with ester-linked
aliphatic segments from studies conducted by the Miller group (Refs. 67,69). ............... 11
Figure 1.6. Polymers incorporating conjugation-break spacers produced by the Sivula
group, where n is 10-12 and m is 4-5 (Refs. 82-83). ........................................................ 13
Figure 1.7. Polymer family synthesized by the Li group (Refs. 62-64) using polyurethane
and polyester condensation polymerizations. ................................................................... 14
Figure 1.8. Polymer structure utilized in CBS studies by the Mei group, where x was varied
from 0 to 1 and n was varied from 0 to 20. (Refs. 84-88) ................................................ 15
Figure 1.9. Polymers synthesized by the Bao lab of Stanford. In both structures 9 and 10,
x was varied between 0-20%. (Refs. 84,60) ..................................................................... 17
Figure 1.10. Schematic of two primary ways to make non-conjugated polymers with
electroactive pendants. A represents post-polymerization modification of a non-conjugated
backbone with an electroactive dye and B represents direct polymerization of a monomer
with an electroactive dye preattached. .............................................................................. 21
Figure 1.11. Simple electroactive pendant polymers with cofacial π-π stacking. A) Vinyl
carbazole and its polymer, poly(vinyl carbazole), or PVK. B) Dibenzofulvene and its
polymer, poly(dibenzofulvene), or pDBF. ........................................................................ 22
Figure 1.12. Schematic of ring opening metathesis polymerization (ROMP, A) and acyclic
diene metathesis polymerization (ADMET, B) with electroactive pendant dyes. ............ 23
xxiii

Figure 1.13. Comparison of electroactive pendant polymers made by the grafting-on
strategy (post-polymerization modification, 11) against those made by direct
polymerization (12 and 13). Ref. 122. .............................................................................. 27
Figure 1.14. Typical nanomorphologies formed in thermodynamic equilibrium from coil-
coil block copolymers. Going from left to right, the relative fraction of the red monomer is
decreasing as the fraction of the black monomer increases. Reprinted with permission from
Marencic, A. P.; Register, R. A. Controlling Order in Block Copolymer Thin Films for
Nanopatterning Applications. Annual Review of Chemical and Biomolecular Engineering
2010, 1, 277-297. Copyright 2010 Annual Reviews. ....................................................... 30
Figure 1.15. (A) Coil-coil-type block copolymer with hole-transporting triphenylamine
pendant and electron-transporting perylene diimide blocks and (B) a TEM cross section of
a film showing thermodynamic phase segregation of the block copolymer. (B) is reprinted
with permission from Lindner, S. M.; Thelakkat, M. Nanostructures of N-Type Organic
Semiconductor in a P-Type Matrix Via Self-Assembly of Block Copolymers.
Macromolecules 2004, 37, 8832-8835. Copyright 2004 American Chemical Society. ... 31
Figure 2.1. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. The P3HTT-
DPP-10% absorption spectrum (dashed black line) is provided for a fully conjugated
reference. (■) 10% T-8-T; (●) 20% T-8-T; (▲) 30% T-8-T; (▼) 40% T-8-T. ............... 53
Figure 2.2. Pseudo free-standing tensile test of P3HT-based semi-random copolymers.
Schematic representation of experimental setup, which includes a floating film, linear
actuator, load cell, and trough filled with water. Inset photograph demonstrates the real
experimental apparatus ..................................................................................................... 62
xxiv

Figure 2.3. Stress-strain curves and corresponding mechanical properties as functions of
the length (with 10% incorporation of the spacer) and fraction of the aliphatic spacer (with
10 carbon atoms in the spacer). (a,d) Representative stress-strain curves of pristine films
obtained using the film-on-water technique. Correlation of toughness and elastic modulus
with (b) the length, and (e) the fraction of the spacer. Values of toughness are obtained by
integrating the total area under the stress-strain curves. Values of elastic modulus are
calculated as the slope of the linear region of the graph. Relationship between ultimate
tensile strength (UTS) and (c) the length or (f) the fraction of the spacer. UTS values were
obtained from the stress–strain curves from the stress at fracture. Mean values and error
bars (standard deviations based on 95% confidence bounds) are based on data collected
from at least three separate measurements. Dashed lines are to guide the eyes. .............. 64
Figure 3.1. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast
from chloroform and (a) placed in a N2 cabinet for 30 minutes or (b) annealed under N2 at
150 °C for 30 minutes. The P3HTT-dtdDPP-10% absorption spectrum (dashed black line)
is provided for a fully conjugated reference. (■) 10% T-8-T; (●) 20% T-8-T; (▲) 30% T-
8-T; (▼) 40% T-8-T. ........................................................................................................ 76
Figure 4.1. (a) UV-Vis absorption spectra of the T-8-T series thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. (b) UV-Vis absorption
spectra of the T-8-T series thin films spin-cast from o-DCB and annealed under N2 at 150
°C for 30 minutes. The P3HTT-DPP-10% absorption spectrum (dashed black line) is
provided for a fully conjugated reference. (■) 10% T-8-T/10% DPP; (●) 20% T-8-T/10%
DPP; (▲) 30% T-8-T/10% DPP; (▼) 40% T-8-T/10% DPP; (♦) 50% T-8-T/10% DPP. 93
xxv

Figure 5.1. Two general approaches to creating non-conjugated polymers with
electroactive pendants. In A, monomer with a reactive functional group is first
polymerized, then coupled to the electroactive dye as a form of post-polymerization
modification. In B, the monomer is first coupled with the electroactive dye and then
polymerized directly. ...................................................................................................... 104
Figure 5.2. The fundamental diketopyrrolopyrrole (DPP) unit and the form typically used
in organic electronics, with aryl substituents at the 3 and 6 positions and solubilizing alkyl
chains on the amide nitrogens. ........................................................................................ 105
Figure 5.3. Solution (left) and film (right) absorption spectra for three clicked polymers
and small molecule reference. Solution spectra collected in chloroform; films spin-cast
from chloroform solutions and kept in a N2 cabinet for 30 min. .................................... 117
Figure 5.4. Fluorescence spectra of three clicked polymers and reference small molecule
in chloroform solution..................................................................................................... 118
Figure A1. Trends in number-averaged polymer molecular weights as measured by SEC
......................................................................................................................................... 133
Figure A2.
1
H NMR spectrum of 20% T-10-T. The NMR was taken in CDCl3 at 50 °C on
a 600 MHz instrument. This spectrum is being used as an example of how the amount of
monomer in each polymer was calculated. The aromatic peak at 8.88 ppm from the DPP
monomer is highlighted in blue; the aromatic peak at 6.73 ppm from the T-10-T monomer
is in orange; and the aromatic peak at 6.97 ppm from the 3HT monomer is in yellow. The
integrations of the DPP and T-10-T peaks match at 1.00 and 1.01, as they should. Each of
these peaks account for 2 hydrogens on these symmetrical monomers. The 3HT peak in
yellow accounts for only one hydrogen on the asymmetric monomer, therefore its
xxvi

integration is doubled from 1.46 to 2.92. If these 3 integrations are summed, they give a
value of 4.93. Dividing each integration by this sum gives the percentage of each monomer
in the polymer chain. Therefore in this example, there is 20.3% DPP, 20.5% T-10-T, and
59.2% 3HT incorporated. This holds true to the feed ratio of 20% DPP, 20% T-10-T, and
60% 3HT. ........................................................................................................................ 135
Figure A3. Stacked
1
H NMR spectra of 10% T-4-T, 20% T-4-T, 30% T-4-T, and 40% T-
4-T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks
from DPP monomer are highlighted in blue; T-4-T in orange; and 3HT in yellow. ...... 136
Figure A4. Stacked
1
H NMR spectra of 10% T-6-T, 20% T-6-T, 30% T-6-T, and 40% T-
6-T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks
from DPP monomer are highlighted in blue; T-6-T in orange; and 3HT in yellow. ...... 137
Figure A5. Stacked
1
H NMR spectra of 10% T-8-T, 20% T-8-T, 30% T-8-T, and 40% T-
8-T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks
from DPP monomer are highlighted in blue; T-8-T in orange; and 3HT in yellow. ...... 138
Figure A6. Stacked
1
H NMR spectra of 10% T-10-T, 20% T-10-T, 30% T-10-T, and 40%
T-10-T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks
from DPP monomer are highlighted in blue; T-10-T in orange; and 3HT in yellow. .... 139
Figure A7. UV-Vis absorption spectra of 10% CBS subfamily thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 140
Figure A8. UV-Vis absorption spectra of 20% CBS subfamily thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 141
xxvii

Figure A9. UV-Vis absorption spectra of 30% CBS subfamily thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 142
Figure A10. UV-Vis absorption spectra of 40% CBS subfamily thin films spin-cast from
o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 143
Figure A11. UV-Vis absorption spectra of the T-4-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference. ............................... 144
Figure A12. UV-Vis absorption spectra of the T-6-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference. ............................... 145
Figure A13. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference. ............................... 146
Figure A14. UV-Vis absorption spectra of the T-10-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference. ............................... 147
Figure A15. Trends in optical bandgaps of polymer family calculated from absorption
band edge. ....................................................................................................................... 149
Figure A16. Trends in solvent effects on optical bandgap for T-6-T subfamily of polymers.
......................................................................................................................................... 150
xxviii

Figure A17. Cyclic voltammograms of the 10% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference. ........................................................................................ 151
Figure A18. Cyclic voltammograms of the 20% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. .......................................................................... 151
Figure A19. Cyclic voltammograms of the 30% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. .......................................................................... 152
Figure A20. Cyclic voltammograms of the 40% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. .......................................................................... 152
Figure A21. Cyclic voltammograms of the T-4-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference. ........................................................................................ 153
Figure A22. Cyclic voltammograms of the T-6-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference. ........................................................................................ 153
Figure A23. Cyclic voltammograms of the T-8-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference. ........................................................................................ 154
Figure A24. Cyclic voltammograms of the T-10-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference. ........................................................................................ 154
Figure A25. Trends in highest occupied molecular orbitals of polymer family calculated
from oxidation onset versus ferrocene. ........................................................................... 155
xxix

Figure A26. DSC curve for 3.5 mg of P3HTT-DPP-10% using a scan rate of 10 °C/min.
The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions). ................................................................................................. 156
Figure A27. DSC curve for 5.0 mg of 30% T-8-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 157
Figure A28. DSC curve for 4.2 mg of 30% T-10-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 157
Figure A29. DSC curve for 4.3 mg of 40% T-6-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 158
Figure A30. DSC curve for 4.3 mg of 40% T-8-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 158
Figure A31. DSC curve for 3.8 mg of 40% T-10-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 159
Figure A32. Trends in melting points for polymer family obtained from DSC curves. 160
Figure A33. Trends in crystallization points for polymer family obtained from DSC
curves. ............................................................................................................................. 161
xxx

Figure A34. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-coated from o-dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for
30 minutes. ...................................................................................................................... 163
Figure A35. Trends in average hole mobilities for as-cast films. .................................. 167
Figure A36. Trends in average hole mobilities for annealed films. ............................... 168
Figure A37. Trends in elastic moduli for polymer family. ............................................ 169
Figure A38. Trends in COS percentages for polymer family. ....................................... 170
Figure A39. Micrograph of P3HTT-DPP-10% film on PDMS at 24% elongation. ...... 172
Figure A40. Micrograph of 10% T-4-T film on PDMS at 48% elongation................... 173
Figure A41. Micrograph of 20% T-6-T film on PDMS at 52% elongation................... 174
Figure B1. Trends in number-averaged polymer molecular weights as measured by SEC.
......................................................................................................................................... 182
Figure B2. Trends in polymerization yields from chloroform Soxhlet fraction. ........... 183
Figure B3.
1
H NMR spectrum of 40% T-4-T. The NMR was taken in CDCl3 at 50 °C on
a 600 MHz instrument. This spectrum is being used as an example of how the amount of
monomer in each polymer was calculated. The aromatic peak at 8.86 ppm from the DPP
monomer is highlighted in blue; the aromatic peak at 6.75 ppm from the T-4-T monomer
is in orange; and the aromatic peak at 6.98 ppm from the 3HT monomer is in yellow. The
integrations of the DPP and T-4-T peaks match at 1.00 and 1.01, as they should. Each of
these peaks account for 2 hydrogens on these symmetrical monomers. The 3HT peak in
yellow accounts for only one hydrogen on the asymmetric monomer, therefore its
integration is doubled from 0.27 to 0.54. If these 3 integrations are summed, they give a
value of 2.55. Dividing each integration by this sum gives the percentage of each monomer
xxxi

in the polymer chain. Therefore in this example, there is 39.2% DPP, 39.6% T-8-T, and
21.2% 3HT incorporated. This holds true to the feed ratio of 40% DPP, 40% T-8-T, and
20% 3HT. ........................................................................................................................ 184
Figure B4. Stacked
1
H NMR spectra of 10% T-4-T and 40% T-4-T. NMRs were taken in
CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from DPP monomer are
highlighted in blue; T-4-T in orange; and 3HT in yellow. ............................................. 185
Figure B5. Stacked
1
H NMR spectra of 10% T-8-T, 20% T-8-T, 30% T-8-T, and 40% T-
8-T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks
from DPP monomer are highlighted in blue; T-8-T in orange; and 3HT in yellow. ...... 186
Figure B6. Stacked
1
H NMR spectra of 10% T-10-T and 20% T-10-T. NMRs were taken
in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from DPP monomer are
highlighted in blue; T-10-T in orange; and 3HT in yellow. ........................................... 187
Figure B7. UV-Vis absorption spectra of 10% CBS subfamily thin films spin-cast from
chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10% absorption
spectrum provided for fully conjugated reference. ......................................................... 188
Figure B8. UV-Vis absorption spectra of the T-4-T CBS subfamily thin films spin-cast
from chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 189
Figure B9. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast
from chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 190
xxxii

Figure B10. UV-Vis absorption spectra of the T-10-T CBS subfamily thin films spin-cast
from chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 191
Figure B11. UV-Vis absorption spectra of 10% CBS subfamily thin films spin-cast from
chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 192
Figure B12. UV-Vis absorption spectra of the T-4-T CBS subfamily thin films spin-cast
from chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 193
Figure B13. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast
from chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 194
Figure B14. UV-Vis absorption spectra of the T-10-T CBS subfamily thin films spin-cast
from chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 195
Figure B15. Trends in optical bandgaps of as-cast polymer family calculated from
absorption band edge. ..................................................................................................... 197
Figure B16. Trends in optical bandgaps of annealed polymer family calculated from
absorption band edge. ..................................................................................................... 198
Figure B17. Cyclic voltammograms of the 10% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference. ......................................................................... 199
xxxiii

Figure B18. Cyclic voltammograms of the T-4-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference. ......................................................................... 200
Figure B19. Cyclic voltammograms of the T-8-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference. ......................................................................... 201
Figure B20. Cyclic voltammograms of the T-10-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference. ......................................................................... 202
Figure B21. Trends in highest occupied molecular orbitals of polymer family calculated
from oxidation onset versus ferrocene. ........................................................................... 203
Figure B22. DSC curve for P3HTT-dtdDPP-10% using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 204
Figure B23. DSC curve for 10% T-4-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 205
Figure B24. DSC curve for 40% T-4-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 206
Figure B25. DSC curve for 10% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 206
xxxiv

Figure B26. DSC curve for 20% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 207
Figure B27. DSC curve for 30% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 207
Figure B28. DSC curve for 40% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 208
Figure B29. DSC curve for 10% T-10-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 208
Figure B30. DSC curve for 20% T-10-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).
......................................................................................................................................... 209
Figure B31. Trends in melting points for polymer family obtained from DSC curves. 210
Figure B32. Trends in crystallization points for polymer family obtained from DSC curves.
......................................................................................................................................... 211
Figure B33. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-cast from chloroform and placed in a N2 cabinet for 30 minutes. ................. 213
Figure B34. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-cast from chloroform and annealed under N2 at 150 °C for 30 minutes........ 214
Figure B35. Trends in as-cast hole mobilities of eight CBS polymers. ......................... 216
xxxv

Figure B36. Trends in annealed hole mobilities of eight CBS polymers. ..................... 217
Figure B37. Trends in elastic moduli for polymer family. ............................................ 218
Figure B38. Trends in toughness for polymer family. ................................................... 219
Figure B39. Trends in UTS for polymer family. ........................................................... 220
Figure B40. Trends in fracture strength for polymer family.......................................... 221
Figure B41. Trends in fracture strain for polymer family. ............................................. 222
Figure C1. Trends in number-averaged polymer molecular weights as measured by SEC.
......................................................................................................................................... 229
Figure C2. Trends in polymerization yields from chloroform Soxhlet fraction. ........... 230
Figure C3.
1
H NMR spectrum for 1,8-bis((5-bromo)thiophen-2-yl)octane taken in CDCl3
at 25 °C on a 600 MHz instrument. ................................................................................ 231
Figure C4.
13
C NMR spectrum for 1,8-bis((5-bromo)thiophen-2-yl)octane taken in CDCl3
at 25 °C on a 600 MHz instrument. ................................................................................ 232
Figure C5.
1
H NMR spectrum of 20% T-8-T/10% DPP. The NMR was taken in CDCl3 at
50 °C on a 600 MHz instrument. This spectrum is being used as an example of how the
amount of monomer in each polymer was calculated. The aromatic peak at 8.88 ppm from
the DPP monomer is highlighted in blue; the aromatic peaks from 6.63-6.73 ppm from the
T-8-T monomer are in orange; and the aromatic peak at 6.96 ppm from the 3HT monomer
is in yellow. The integration of the DPP peak is 1.00 and that of the T-8-T peak is 2.10.
Each of these peaks account for 2 hydrogens on these symmetrical monomers. The 3HT
peak in yellow accounts for only one hydrogen on the asymmetric monomer, therefore its
integration is doubled from 3.53 to 7.06. If these 3 integrations are summed, they give a
value of 10.16. Dividing each integration by this sum gives the percentage of each
xxxvi

monomer in the polymer chain. Therefore in this example, there is 9.8% DPP, 20.7% T-8-
T, and 69.5% 3HT incorporated. This holds true to the feed ratio of 10% DPP, 20% T-8-
T, and 70% 3HT. ............................................................................................................. 233
Figure C6. Stacked
1
H NMR spectra of 10% T-8-T/10% DPP, 20% T-8-T/10% DPP, 30%
T-8-T/10% DPP, 40% T-8-T/10% DPP, and 50% T-8-T/10% DPP. NMRs were taken in
CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from DPP monomer are
highlighted in blue; T-8-T in orange; and 3HT in yellow. ............................................. 234
Figure C7. UV-Vis absorption spectra of the T-8-T CBS family thin films spin-cast from
o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 235
Figure C8. UV-Vis absorption spectra of the T-8-T CBS family thin films spin-cast from
o-dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference. ............................... 236
Figure C9. Trends in optical bandgaps of unannealed polymer family calculated from
absorption band edge. ..................................................................................................... 237
Figure C10. Trends in optical bandgaps of annealed polymer family calculated from
absorption band edge. ..................................................................................................... 238
Figure C11. Cyclic voltammograms of the T-8-T CBS family of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference. ........................................................................................ 239
Figure C12. Trends in highest occupied molecular orbitals of polymer family calculated
from oxidation onset versus ferrocene. ........................................................................... 240
xxxvii

Figure C13. DSC curve for P3HTT-DPP-10% using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 241
Figure C14. DSC curve for 5.6 mg of 10% T-8-T/10% DPP using a scan rate of 10 °C/min.
The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions). ................................................................................................. 242
Figure C15. DSC curve for 20% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 242
Figure C16. DSC curve for 30% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 243
Figure C17. DSC curve for 40% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 243
Figure C18. DSC curve for 50% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 244
Figure C19. Grazing-incidence X-ray diffraction patterns of thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference. ....................................... 247
Figure C20. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-cast from o-dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for 30
xxxviii

minutes. P3HTT-DPP-10% absorption spectrum provided for fully conjugated reference.
......................................................................................................................................... 248
Figure C21. Trends in hole mobility of polymer family in as-cast films. ..................... 250
Figure C22. Trends in hole mobility of polymer family in annealed films. .................. 251
Figure C23. Trends in elastic moduli for polymer family. ............................................ 252
Figure C24. Trends in toughness for polymer family. ................................................... 253
Figure C25. Trends in UTS for polymer family. ........................................................... 254
Figure C26. Trends in fracture strength for polymer family. ........................................ 255
Figure C27. Trends in fracture strain for polymer family. ............................................ 256
Figure D1.
1
H NMR spectrum of poly(TIPS ethynylstyrene) (III) taken at 50 °C in CDCl3
on a 600 MHz instrument. .............................................................................................. 269
Figure D2.
13
C NMR spectrum of poly(TIPS ethynylstyrene) (III) taken at 50 °C in CDCl3
on a 600 MHz instrument. .............................................................................................. 270
Figure D3.
1
H NMR spectrum of poly(ethynylstyrene) (IV) taken at 50 °C in CDCl3 on a
600 MHz instrument. ...................................................................................................... 271
Figure D4.
13
C NMR spectrum of poly(ethynylstyrene) (IV) taken at 50 °C in CDCl3 on a
600 MHz instrument. ...................................................................................................... 272
Figure D5.
1
H NMR spectrum of monoalkylated dtdDPP (VI) taken at 25 °C in CDCl3 on
a 400 MHz instrument. ................................................................................................... 273
Figure D6.
1
H NMR spectrum of azidobutyl dtdDPP (VIII) taken at 25 °C in CDCl3 on a
400 MHz instrument. ...................................................................................................... 274
Figure D7.
1
H NMR spectrum of clicked polymer (IX) taken at 50 °C in CDCl3 on a 600
MHz instrument. ............................................................................................................. 275
xxxix

Figure D8.
13
C NMR spectrum of clicked polymer (IX) taken at 50 °C in CDCl3 on a 600
MHz instrument. ............................................................................................................. 276
Figure D9. Stacked
1
H NMR spectra of (III, VI, IX) taken in CDCl3 at 50 °C on a 600
MHz instrument. Distinctive peaks are highlighted from each successive reaction. The peak
highlighted in red in the alkyl region of the bottom spectrum is from the triisopropyl
protecting group and it disappears entirely in the next spectrum, being replaced by the
deprotected alkyne peak highlighted in green in the middle spectrum. That peak then
disappears entirely and new DPP resonances appear in the top spectrum, highlighted in
blue, indicating that the reaction proceeded to completion. ........................................... 277
Figure D10. Stacked
13
C NMR spectra of (III, VI, IX) taken in CDCl3 at 50 °C on a 600
MHz instrument. Distinctive peaks are highlighted from each successive reaction. The
peaks highlighted in red on the bottom spectrum are from the triisopropyl protecting group
and it disappears entirely in the next spectrum, being replaced by the deprotected alkyne
peak highlighted in green in the middle spectrum. That peak then disappears entirely and
new DPP resonances appear in the top spectrum, highlighted in blue, indicating that the
reaction proceeded to completion. .................................................................................. 278
Figure D11.
1
H NMR spectrum of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in CDCl3
on a 600 MHz instrument. .............................................................................................. 279
Figure D12.
13
C NMR spectrum of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in
CDCl3 on a 600 MHz instrument. ................................................................................... 280
Figure D13. Deconvolution of the aromatic peaks at ~143 ppm on the
13
C NMR spectrum
of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals 26.4% mm, 39.0% mr/rm, and 34.6% rr triads.
5-7
..................... 280
xl

Figure D14. Deconvolution of the methine peaks at ~42-46 ppm on the
13
C NMR spectrum
of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals pentad peaks, but we were unable to fully assign the pentads. .. 281
Figure D15.
1
H NMR spectrum of s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3
on a 600 MHz instrument. .............................................................................................. 281
Figure D16.
13
C NMR spectrum of s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3
on a 600 MHz instrument. .............................................................................................. 282
Figure D17. Deconvolution of the aromatic peak at ~143 ppm on the
13
C NMR spectrum
of s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals only one peak, indicating rr triads > 95%.
8
............................... 282
Figure D18. Deconvolution of the methine peak at ~44 ppm on the
13
C NMR spectrum of
s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals only one peak, indicating rr triads > 95%.
8
............................... 283
Figure D19.
1
H NMR spectrum of a-poly(4-bromomethylstyrene) (XII-a) taken at 50 °C
in CDCl3 on a 600 MHz instrument. ............................................................................... 283
Figure D20.
13
C NMR spectrum of a-poly(4-bromomethylstyrene) (XII-a) taken at 50 °C
in CDCl3 on a 600 MHz instrument. ............................................................................... 284
Figure D21.
1
H NMR spectrum of s-poly(4-bromomethylstyrene) (XII-s) taken at 50 °C
in CDCl3 on a 600 MHz instrument. ............................................................................... 285
Figure D22.
13
C NMR spectrum of s-poly(4-bromomethylstyrene) (XII-s) taken at 50 °C
in CDCl3 on a 600 MHz instrument. ............................................................................... 286
Figure D23.
1
H NMR spectrum of a-poly(4-azidomethylstyrene) (XIII-a) taken at 55 °C
in DMSO-d6 on a 600 MHz instrument. ......................................................................... 287
xli

Figure D24.
13
C NMR spectrum of a-poly(4-azidomethylstyrene) (XIII-a) taken at 55 °C
in DMSO-d6 on a 600 MHz instrument. ......................................................................... 288
Figure D25.
1
H NMR spectrum of s-poly(4-azidomethylstyrene) (XIII-s) taken at 55 °C
in toluene on a 600 MHz instrument............................................................................... 289
Figure D26.
13
C NMR spectrum of s-poly(4-azidomethylstyrene) (XIII-s) taken at 25 °C
in tetrahydrofuran on a 600 MHz instrument. ................................................................ 290
Figure D27.
1
H NMR spectrum of unsymmetrical propyne/dtdDPP (XIV) taken at 25 °C
in CDCl3 on a 400 MHz instrument. ............................................................................... 291
Figure D28.
1
H NMR spectrum of a-poly(styrene-click-DPP) (XV-a) taken at 50 °C in
CDCl3 on a 600 MHz instrument. ................................................................................... 292
Figure D29.
13
C NMR spectrum of a-poly(styrene-click-DPP) (XV-a) taken at 50 °C in
CDCl3 on a 600 MHz instrument. ................................................................................... 293
Figure D30.
1
H NMR spectrum of s-poly(styrene-click-DPP) (XV-s) taken at 50 °C in
CDCl3 on a 600 MHz instrument. ................................................................................... 294
Figure D31.
13
C NMR spectrum of s-poly(styrene-click-DPP) (XV-s) taken at 50 °C in
CDCl3 on a 600 MHz instrument. ................................................................................... 295
Figure D32. Stacked
1
H NMR spectra of a-poly(styrene)s (XI-a, XII-a, XIII-a, XV-a)
taken on a 600 MHz instrument. Distinctive peaks are highlighted from each successive
reaction. The peak highlighted in red in the alkyl region of the bottom spectrum is from the
para-methyl group on the styrene and it disappears entirely in the next spectrum, being
replaced by the methylene peak (adjacent to a bromine) highlighted in green in the second
spectrum from the bottom. The peak highlighted in green then shifts to the peak highlighted
in teal as the bromine is converted to an azide in the third spectrum from the bottom. That
xlii

peak then disappears entirely and new DPP resonances appear in the top spectrum,
highlighted in purple, indicating that the reaction proceeded to completion. ................. 296
Figure D33. Stacked
13
C NMR spectra of a-poly(styrene)s (XI-a, XII-a, XIII-a, XV-a)
taken on a 600 MHz instrument. Distinctive peaks are highlighted from each successive
reaction. The peak highlighted in red in the alkyl region of the bottom spectrum is from the
para-methyl group on the styrene and it disappears entirely in the next spectrum, being
replaced by the methylene peak (adjacent to a bromine) highlighted in green in the second
spectrum from the bottom. The peak highlighted in green then shifts to the peak highlighted
in teal as the bromine is converted to an azide in the third spectrum from the bottom. That
peak then disappears entirely and new DPP resonances appear in the top spectrum,
highlighted in purple, indicating that the reaction proceeded to completion. ................. 297
Figure D34. Stacked
1
H NMR spectra of s-poly(styrene)s (XI-s, XII-s, XIII-s, XV-s)
taken on a 600 MHz instrument. Distinctive peaks are highlighted from each successive
reaction. The peak highlighted in red in the alkyl region of the bottom spectrum is from the
para-methyl group on the styrene and it disappears entirely in the next spectrum, being
replaced by the methylene peak (adjacent to a bromine) highlighted in green in the second
spectrum from the bottom. The peak highlighted in green then shifts to the peak highlighted
in teal as the bromine is converted to an azide in the third spectrum from the bottom. That
peak then disappears entirely and new DPP resonances appear in the top spectrum,
highlighted in purple, indicating that the reaction proceeded to completion. ................. 298
Figure D35. Stacked
13
C NMR spectra of s-poly(styrene)s (XI-s, XII-s, XIII-s, XV-s)
taken on a 600 MHz instrument. Distinctive peaks are highlighted from each successive
reaction. The peak highlighted in red in the alkyl region of the bottom spectrum is from the
xliii

para-methyl group on the styrene and it disappears entirely in the next spectrum, being
replaced by the methylene peak (adjacent to a bromine) highlighted in green in the second
spectrum from the bottom. The peak highlighted in green then shifts to the peak highlighted
in teal as the bromine is converted to an azide in the third spectrum from the bottom. That
peak then disappears entirely and new DPP resonances appear in the top spectrum,
highlighted in purple, indicating that the reaction proceeded to completion. ................. 299
Figure D36. Click reaction monitored by IR spectroscopy for a-poly(styrene-azide) (XV-
a). The disappearance of the azide stretching peak at 2100 cm
-1
is indicative of the reaction
going to completion. ....................................................................................................... 300
Figure D37. Click reaction monitored by IR spectroscopy for s-poly(styrene-azide) (XV-
s). The disappearance of the azide stretching peak at 2100 cm
-1
is indicative of the reaction
going to completion. ....................................................................................................... 301
Figure D38. UV-Vis absorption spectra of clicked polymers in solution (chloroform).
Small molecule dtdDPP (XVI) absorption spectrum provided for reference. ................ 302
Figure D39. UV-Vis absorption spectra of clicked polymer thin films spin-cast from
chloroform and placed in a N2 cabinet for 30 minutes. Small molecule dtdDPP (XVI)
absorption spectrum provided for reference. .................................................................. 303
Figure D40. Normalized dtdDPP (XVI) absorption spectra in solution (chloroform) and
film .................................................................................................................................. 304
Figure D41. Normalized aPS-alkyne-click-DPP (IX) absorption spectra in solution
(chloroform) and film ..................................................................................................... 305
Figure D42. Normalized aPS-azide-click-DPP (XV-a) absorption spectra in solution
(chloroform) and film ..................................................................................................... 306
xliv

Figure D43. Normalized sPS-azide-click-DPP (XV-s) absorption spectra in solution
(chloroform) and film ..................................................................................................... 307
Figure D44. Normalized solution emission spectra in CHCl3. Small molecule dtdDPP
(XVI) emission spectrum provided for reference. .......................................................... 308
Figure D45. Normalized dtdDPP (XVI) absorption and emission spectra in CHCl3 .... 309
Figure D46. Normalized aPS-alkyne-click-DPP (IX) absorption and emission spectra in
CHCl3 .............................................................................................................................. 310
Figure D47. Normalized aPS-azide-click-DPP (XV-a) absorption and emission spectra in
CHCl3 .............................................................................................................................. 311
Figure D48. Normalized sPS-azide-click-DPP (XV-s) absorption and emission spectra in
CHCl3 .............................................................................................................................. 312
Figure D49. DSC curve for 7.1 mg of dtdDPP (XVI) using a scan rate of 10 °C/min. The
top curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions)....................................................................................................................... 313
Figure D50. DSC curve for 5.3 mg of aPS-alkyne-click-DPP (IX) using a scan rate of 10
°C/min. The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions). ................................................................................................. 314
Figure D51. DSC curve for 7.5 mg of aPS-azide-click-DPP (XV-a) using a scan rate of 10
°C/min. The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions). ................................................................................................. 315
Figure D52. DSC curve for 7.2 mg of sPS-azide-click-DPP (XV-s) using a scan rate of 10
°C/min. The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions). ................................................................................................. 316
xlv

Figure D53. Grazing-incidence X-ray diffraction pattern of thin film of dtdDPP (XVI)
spin-cast from chloroform and placed in a N2 cabinet for 30 minutes. .......................... 317
Figure D54. Cyclic voltammograms of the clicked polymer films taken in dry acetonitrile
under nitrogen atmosphere. The small molecule dtdDPP (XVI) cyclic voltammogram is
provided for reference. .................................................................................................... 319

xlvi

Abstract
Electronic devices are becoming progressively smaller and more integrated into daily
life. Organic semiconductors offer the possibility of making these devices lightweight,
flexible, and easy to manufacture.
Conjugated polymers are widely explored for a variety of organic electronic devices.
However, there are some drawbacks to using the traditional linear, fully conjugated
polymer. This dissertation explores some alternative architectures for semiconducting
polymers in an attempt to overcome some of the challenges facing traditional architectures.
An overview of alternative architectures for semiconducting polymers is presented in
Chapter 1. Special attention is paid to linear conjugated polymers with non-conjugated
segments interspersed, so-called conjugation-break spacers (CBSs), as well as non-
conjugated polymers with electroactive pendants. The history, advantages, disadvantages,
and recent trends for these systems are discussed.
In Chapter 2 conjugation-break spacers are applied to the semi-random polymer
design developed previously in the Barry C. Thompson research group. The effects of
spacer length and content as it varies concurrently with the other monomers in the polymer
system are studied. Charge mobility is observed to decrease with increasing CBS
incorporation, but increase with increasing CBS length. This is accompanied by an
improvement in mechanical properties as measured by film-on-elastomer (FOE) and film-
on-water (FOW) techniques. This study results in CBS polymers with comparable
electronic properties but greatly improved mechanical properties relative to their fully
conjugated analogs.
xlvii

Some polymers within the family synthesized for Chapter 2 could not be analyzed due
to poor solubility. Consequently, a subset of these polymers with longer solubilizing side-
chains are synthesized and analyzed in Chapter 3. Solubility and mechanical properties
are seen to improve greatly, with some polymers exhibiting properties approaching
elastomeric limits. Free-standing films measured by the FOW technique could be stretched
to many times their original length before breaking. However, this improvement in
mechanical properties is accompanied by a diminishment in electronic properties.
Within the series of CBS polymers synthesized for Chapter 2, the content of all three
monomers was changed simultaneously, making it difficult to isolate the effects of the CBS
content on electronic and mechanical properties. Therefore, in Chapter 4 a small set of
CBS polymers are synthesized and evaluated in which the content of one monomer is held
constant and the other two are varied. Similar to Chapter 3, an improvement in the
mechanical performance of these polymers is observed along with a decrease in desirable
electronic properties.
Finally, in Chapter 5 a set of novel non-conjugated polystyrene-based polymers with
electroactive diketopyrrolopyrrole (DPP) pendants is synthesized and characterized. The
three polymers differ in tacticity and spacer length linking the DPP pendant to the
polystyrene backbone. There is no observed effect of tacticity on optical properties, but a
possible influence of spacer length on optical properties.

1

Chapter 1: Non-Traditional Semiconducting Polymer Architectures

1.1 Background/Motivation
1.1.1 Benefits of organics vs. inorganics
In recent decades there has been increasing interest in using organic semiconductors to
replace or supplement inorganic semiconductors. These organic semiconductors can be
used in devices ranging from photovoltaics (PVs) and light-emitting devices (LEDs) to
field-effect transistors (FETs) and electrochromics (ECs). The use of organic materials,
those primarily composed of carbon and hydrogen atoms, allows these electronics to be
made thinner, flexible, and more lightweight than their inorganic counterparts. This is
attributable in part to a higher absorption coefficient in organic materials,
1,2
and the fact
that organic materials tend to be less dense than inorganic materials.
3,4
Organic electronics
are often considered to be more aesthetically pleasing than inorganic devices due to the
many handles for tuning color available to synthetic organic chemists.
5-7
Organic devices
are also predicted to be cheaper and less energy-intensive to make than their inorganic
counterparts because organic molecules can be made into inks that can be printed,
8,9
and
there is frequently not such a rigorous need for purity as in inorganic wafer technology.
10,11

Though small molecule organic semiconductors are also subject to intense research,
semiconducting polymers will be the focus of this dissertation.
1.1.2 Benefits of fully conjugated polymers vs. small molecules
There are several advantages to using conjugated polymers over small molecules. The
extended conjugation of the polymer backbone allows for electron delocalization (Figure
2

1.1), and charges can be transported in three dimensions: along a polymer backbone,
between π-stacked polymer chains, and across polymer chains.
12

Figure 1.1 The alternating single and double bonds of conjugated polymers allow for
electron delocalization and charge transport in a way that is not accessible to organic
small molecules. In this simplistic example, a segment of polyacetylene is considered
(1.1A). In 1.1B, the filled π-orbitals of the double bonds are depicted, illustrating the
concept that each filled π-orbital is flanked on each side by another filled π-orbital,
allowing the electrons in those orbitals to be delocalized, or not fixed on any particular
atom. In 1.1C, an electron is removed from one π-orbital (shown in white). Its neighbors
can quickly fill the empty space, or hole, leading not only to charge stabilization within
conjugated polymers, but also rapid charge transport along the polymer backbone.
Most polymers are characterized as being semicrystalline, that is they are made up of
segments of amorphous polymer chains as well as segments with three-dimensional
periodicity known as crystallites (Figure 1.2). The amorphous segments of a polymer chain
serve as a bridge between crystallites, allowing charge transport between them and making
the material more mechanically robust.
13
In contrast, small molecules and inorganic
semiconductors form grain boundaries at the edge of their crystallites, across which it is
difficult to achieve charge transport and which serve as fracture points if the material comes
under mechanical strain.
13

3

Figure 1.2. Cartoon of polymer morphology, with dark red crystallites embedded in light
red amorphous segments. The amorphous segments improve the polymer’s mechanical
performance and film integrity while the crystallites are responsible for rapid charge
transport.
These semicrystalline attributes of polymers not only contribute to charge transport
within organic electronics, but also the physical and mechanical properties of the
semiconducting film. Conjugated polymers are often designed to be solution processible
and they can be spin cast, blade coated, slot dye coated, and even spray coated onto a
surface as a semiconducting ink. As the ink dries, it leaves behind a very uniform film that
conforms to the substrate. In contrast, small molecules and inorganic materials have a
tendency to crystallize, providing excellent channels for charge transport,
2,14
but becoming
brittle and leading to uneven films, which in turn lead to device failure.
15

Additionally, the extended conjugation of polymers and possibility of combining
various chromophores in the conjugated backbone allows for the tuning of multiple
properties, including the frontier orbital energy levels, bandgap, absorption spectrum, and
color of the material.
16-18
This is particularly relevant for organic photovoltaics, as
polymers are capable of absorbing broad swaths of the solar spectrum but small molecules
tend have very sharp, narrow absorption bands, underutilizing the solar spectrum.

4

1.1.3 Drawbacks of fully conjugated polymers
Despite the myriad advantages of using conjugated polymers in organic electronics,
they have some inherent weaknesses. In comparison to the non-conjugated polymers
frequently seen in household plastics, fully conjugated polymers tend to be stiff and brittle
with high tensile moduli (meaning a great deal of force is necessary to deform a film – an
undesirable quality) and low ductility (the amount by which a film can be stretched before
cracks appear – also undesirable).
19
For example, the tensile, or elastic (Young’s) modulus,
E, for conventional elastomers such as polydimethylsilane (PDMS) or polyisoprene
(rubber) ranges from 0.3-2.5 MPa,
20,21
and popular thermoplastics such as low-density
polyethylene (LDPE) and high-density polyethylene (HDPE) have E = 0.3-0.7 GPa,
22

whereas conjugated polymers tend to exhibit tensile moduli in the range of 1-4 GPa.
23

Another weakness that conjugated polymers do not share with their non-conjugated
cousins is that the double bonds in the backbone of the polymer are susceptible to
photooxidation and other degradation pathways which render the polymer incapable of
both absorbing light and transporting charge.
24-26
This degradation can even break the
polymer down into small oligomers which do not retain the mechanical, optical, or
electronic properties of the parent chain. Additionally, the twists and kinks that occur along
the conjugated polymer backbone at distorted sp
2
hybridized carbons create electronic
defects that are more susceptible to photodegradation.
27

Conjugated polymers have relatively high melting points and low solubilities, which
inhibits their incorporation into industrial-scale fabrication processes. To address this
weakness, bulky or branched side chains are often added to the conjugated backbone to aid
in solubility or lower the melting point. These side chains, however, can act as charge
5

insulators, preventing the conjugated backbones from drawing close enough to each other
for efficient charge transfer.
28
If these chains are eliminated or shortened, charge transport
improves but solubility and solution processability consequently decrease.
28-31

The polymerization methods most commonly used to create conjugated polymers also
have some undesirable attributes. These polymerization techniques include Stille, Suzuki,
and other transition metal-catalyzed cross-couplings. Unlike the methods available for
creating non-conjugated polymers, these cross-coupling reactions allow for very little
control over polymer molecular weight and dispersity (i.e., they are not living
polymerizations). Since polymer semiconductor device performance has been shown to
depend on both molecular weight and dispersity, this leads in turn to batch-to-batch
variation in device performance.
32
What’s more, precious metal catalysts are necessary for
these polymerizations and a stoichiometric amount of often toxic byproducts are produced
from the condensation reactions. In contrast, non-conjugated polymers can be synthesized
through a great variety of methods, many of which do not require metal catalysts, which
are difficult to remove entirely after the reaction has concluded. Residual metals can hinder
device performance
33
and can prevent the materials from being used in biomedical
applications.
It is evident that fully conjugated semiconducting polymers have several areas of
weakness that should be improved. To address some of the challenges inherent to
traditional linear conjugated polymers, a variety of alternative architectures of
semiconducting polymers has been introduced. Some of these are highlighted and
discussed below.

6

1.2 Alternative Architectures

Figure 1.3. Some basic polymer architectures. Gradient and block copolymers are
typically linear, and vary from the linear polymer only in their composition. In contrast to
branched polymers, star polymers have many arms radiating from one center. Dendritic
polymers are fractal structures that are grown in generations, where each generation is
represented by a different shade of blue in this figure. Crosslinked or network polymers
have chemical bonds linking polymer chains together. Comb and brush polymers are
sometimes known as graft polymers, where the pendant chains may or may not be of the
same composition as the backbone.
Polymers consist of many individual units, or monomers, connected to each other in a
repeating pattern. There are as many different architectures available to polymers as there
are ways in which these units can be linked. Some of the more common architectures are
illustrated in Figure 1.3. By far the most common architecture for semiconducting
polymers is a simple linear polymer. Most monomers used in the synthesis of electroactive
polymers are symmetrical, and are therefore easy to difunctionalize with the groups that
are used in cross-coupling polymerizations. This leads to polymers that grow only in two
dimensions. Branching occurs when a monomer can bond in more than two directions and
chain growth propagates in multiple dimensions. Most of the other common polymer
7

architectures are controlled versions of branching. In a star polymer, there is only one
branching point that becomes the center of multiple growing chains that radiate out. In a
comb or brush polymer, many chains are grafted to/from a polymer backbone. The terms
“comb” and “brush” are used interchangeably by some and by others are distinguished by
the grafting density, with brushes tending to have a higher grafting density than combs
(much like the actual objects). Dendritic polymers are fractal-like molecules exemplified
by their uniformity and perfection. They are built up from a core by generations, with each
generation expanding upon the last in precisely the same branching pattern such that the
final dendrimer is often spherical in nature. A crosslinked polymer has groups that can
form intra and intermolecular attachments between polymer chains, fixing the chains in
place relative to the position of the crosslink.
Each of these architectures could contribute distinct properties to organic electronics.
Dendrimers have long been used to attain site isolation of emitters in polymer OLEDs,
34-36

and have been targeted for organic lasers and other applications due to the ability to tune
optical, electronic, and processing properties independently.
37-40
Star polymers have been
used in OPVs, electrical memory devices, OLEDs, and bioimaging studies with mixed
results.
41-45
Branched polymers have been targeted for OPVs and OFETs, in some cases
explicitly using branching to decrease order in the semiconducting material.
46,47

Crosslinked semiconducting polymers have been pursued for their chemical and
morphological stability in OPV devices.
48
Electroactive block copolymers, gradient
copolymers, conjugated polymer brushes, and several other branched architectures are
discussed in great detail in a book chapter by Geng and Sui in 2016.
49

8

A couple of architectures not included in Figure 1.3 nor in the aforementioned book
chapter include conjugated macrocycles and conjugated microporous polymers.
Conjugated macrocycles are extended rings of electroactive monomers with varying
degrees of control over their bonding order. They have been used for OFETs, OPVs,
50
and
organic photodetectors,
51
both for the benefits imparted by their rigid structure and for their
interesting self-assembly tendancies.
52-54
Conjugated microporous polymers are a type of
network polymer that blurs the line between branching, crosslinking, and covalent organic
frameworks. They are attractive for catalytic activity due to their high surface area and
tunable pore size.
55,56

It is also possible to make semiconducting polymers that are not fully conjugated.
Discrete conjugated units or chromophores can be incorporated into a non-conjugated
polymer backbone either as part of the backbone or pendant to it. By balancing the
conjugated and non-conjugated segments within a linear polymer backbone using
conjugation-break spacers (discussed in Section 1.3), mechanical and optoelectronic
properties can be tuned. This strategy can help semiconducting polymers to overcome the
mechanical limitations inherent to fully conjugated linear polymers. For semiconducting
properties, non-conjugated polymers with pendant electroactive chromophores in the form
of brush or comb graft copolymers require that the pendants stack cofacially to allow
charge transport along the π-π direction. Unlike in conjugated polymers, the backbone is
no longer susceptible to oxidative degradation, which could enhance the polymer’s long-
term stability. A polymer backbone that does not require the formation of aryl-aryl bonds
also opens a wealth of polymerization possibilities, as will be discussed in Section 1.4.
Through the examples given, it can be seen that accessing a variety of geometries allows
9

for tuning of many different properties of semiconducting polymers. Semiconducting
polymers with discrete conjugated units interrupted by conjugation-break spacers and non-
conjugated polymers with electroactive pendants will be discussed in more detail in the
following sections.
1.3 Discrete Conjugation
1.3.1 Motivation/Background
Semiconducting polymers with discrete conjugation are those in which sp
2
hybridized
electroactive segments are interspersed with sp
3
hybridized segments that serve to disrupt
the conjugation of the polymer backbone (Figure 1.4). These sp
3
hybridized segments are
commonly known as conjugation-break spacers, or CBSs. The polymer backbone is
therefore divided into discrete conjugated units that are no longer optoelectronically
interacting with each other in a linear fashion. In the cartoon example of Figure 1.4, the
wavy red lines represent flexible, sp
3
hybridized segments, and the yellow rods represent
rigid, sp
2
hybridized electroactive segments.

Figure 1.4. Cartoon illustration of a segment of polymer with discrete conjugation, where
the wavy red lines are meant to indicate a flexible non-conjugated segment linking the
yellow, rod-like conjugated segments.
CBSs can take the form of a single atom disrupting the conjugation of the chain
57
or
any variety of longer, non-conjugated segment, not limited to linear hydrocarbons.
58-64

These polymers are distinguished from block copolymers both by the length of the
conjugated and non-conjugated segments (both tend to be shorter in CBS polymers than in
block copolymers) and by the method used to produce them. In a polymer with CBS
incorporation, a conjugated monomer is copolymerized with a conjugation-break spacer
10

monomer, whereas in a block copolymer each block is homopolymerized sequentially. This
class of polymers has been in existence nearly since the discovery of conjugated polymers,
though early methods used post-polymerization modification to oxidize double bonds
along the polymer backbone, lending little control over the quantity of conjugation breaks
produced nor their location along the backbone.
65,66

1.3.2 Oligothiophene Systems
In the 1990s, oligothiophenes and polythiophenes were discovered to have interesting
optoelectronic properties, and several groups began investigating how to bridge the gap
between the crystalline small molecules and their polymeric counterparts in a controlled
manner to learn more about their fundamental properties. Hong and Miller introduced
CBSs to oligothiophenes via polyester condensation to halt charge transport along the
polymer backbone.
67
They assumed then that the conductivity observed in the resulting
films could be attributed to charge transport along another axis, i.e., through π-π stacks, a
phenomenon that had yet to be observed in polymers.
68
Simultaneously, several research
groups were exploring the combination of conjugation break spacers with oligothiophenes
to improve solubility, film-forming ability, thermal stability, and mechanical properties
over the small molecules or their parent polymers and achieve tuning of optoelectronic
properties.
58,69-71
Using this strategy, oligothiophenes of defined length would be
synthesized and functionalized, usually with an acyl chloride,
70
then polymerized with a
diol to yield the polymer with alternating conjugated and non-conjugated segments, though
the Miller group used the reverse functionalization strategy (Figure 1.5).
11

Figure 1.5. Polymer structures 1-3 alternating oligothiophene segments with ester-linked
aliphatic segments from studies conducted by the Miller group (Refs. 67,69).
Although one of the purposes of CBS incorporation was to improve the solubility of
the resultant polymers, the researchers were not always successful in that aim,
72
despite the
addition of solubilizing side chains on the oligothiophenes themselves. One study even
attempted to make solution-processable oligothiophenes without sidechains disrupting the
backbone planarity, but found that they had to use polyethyleneoxide CBS segments of
2000 g mol
-1
(about 45 repeat units) to solubilize the oligothiophene blocks.
73
Through the
aforementioned studies, great progress was made in understanding the fundamental
properties of semiconducting polythiophenes, particularly their optoelectronic properties
in relation to effective conjugation length. However, the utility of these polymers in organic
electronic devices was rarely probed, though the Miller group tested their polymers in
organic light-emitting diodes (OLEDs) and photodiodes, obtaining photocurrents
comparable to those reported for poly(3-alkylthiophenes) at the time (1 µA cm
-2
in p3OT
vs 20 µA cm
-2
for structure 3).
72,74
The first reports of this class of thiophene-based
polymers appeared in the early 1990s, and by the end of the decade, research in this area
had mostly died out, although occasional studies using oligothiophenes and CBSs surface
to this day.
75,76

12

1.3.3 Other Studies Using Simple Electroactive Units
While much of the initial research into using conjugation break spacers in
semiconducting polymers focused on oligothiophenes, a few other systems were studied as
well. Similar to the rationale listed above, arylenevinylene-type polymers incorporating
CBSs were studied with the goal of tailoring their bandgap and emission for use in OLEDs
while improving solubility and film quality.
77-80
An interesting study from Woodhouse et
al. in 2009 sought to mitigate the electronic defects caused by backbone kinking typically
observed in conjugated polymers by incorporating CBSs into a perylene diimide (PDI)-
based polymer.
27
They had also hoped to improve film-formation and adhesion while
retaining the good electronic properties of small molecules and create photostable polymers
for use in organic photovoltaics (OPVs). Though they achieved a greater photostability in
the PDI-based polymer when compared to P3HT, their electronic characterization suffered
from low molecular weight polymers with poor film quality. Other studies hoped the use
of flexible spacers would allow the secondary ordering of PDI and naphthalene diimide
(NDI) polymers into π-π stacks favorable for charge transport, but showed disappointing
results in OPVs when compared to devices made with similar small molecules.
59,81
Though
a frequently stated aim of research incorporating spacers into conjugated polymers was to
improve mechanical properties, no objective measurements of these properties were ever
reported. However, in the time since this field first developed, much progress has been
made in polymerization techniques, electroactive monomer synthesis, and the ability to
measure mechanical properties of semiconducting polymers.

13

1.3.4 Modern Conjugation-Break Spacers

Figure 1.6. Polymers incorporating conjugation-break spacers produced by the Sivula
group, where n is 10-12 and m is 4-5 (Refs. 82-83).
Recent developments in synthetic techniques and analysis have led to a resurgence in
the study of semiconducting polymers with conjugation-break spacers in the last five years.
A complicated strategy emerged in 2014 from the Sivula group, as they attempted to
independently control the length of conjugated segments and the overall polymer chain
length.
82
By using an intentionally unbalanced feed ratio of dibromo and distannyl
monomers in a perfectly alternating Stille polymerization, they created short conjugated
polymer segments primarily terminated with bromine. These segments, with a M n of 11.7
kDa (DP of 10-12), they then polymerized again via Stille coupling with a CBS monomer,
leading to longer polymers with a secondary degree of polymerization of 4-5 (meaning 4-
5 CBSs incorporated into each polymer backbone) (Structure 4). The polymers with
flexible spacers incorporated showed improvements in the thin film morphology and
stability over the short conjugated segments, as well as a four-fold improvement in charge
transport after annealing. However, when compared with a fully conjugated polymer of
similar molecular weight, the spacer polymers only achieved half of the reported charge
mobility. In a follow-up study the Sivula group synthesized a perfectly alternating donor
polymer with CBSs (Structure 5) and used it as an additive in small molecule OPVs.
83

14

Though device performance suffered at incorporation levels higher than 1.5%, the CBS
polymer was shown to improve long-term device thermal stability by suppressing small
molecule crystallization.

Figure 1.7. Polymer family synthesized by the Li group (Refs. 62-64) using polyurethane
and polyester condensation polymerizations.
Simultaneously, the Li group began publishing in the field (Figure 1.7).
62-64
His group
was specifically trying to combine the favorable properties of small molecules with those
of polymers for use in organic photovoltaics. Namely, polymers have excellent film-
formation ability, in contrast to small molecules, which tend to crystallize, leading to rough
films with pinholes which can short circuit the device. However, the photovoltaic
properties are dependent on conjugated polymer molecular weight and dispersity, which
are difficult to control. Thus by polymerizing a discrete electroactive unit, the group could
eliminate the effects of polymer molecular weight and dispersity while increasing the small
molecule’s ability to form smooth and uniform films. They tested this hypothesis by
comparing OPV device performance across a series of polymers with conjugation-break
spacers against devices made with a comparable small molecule and observed large relative
gains in overall device efficiency, primarily from improving the device fill factor from
15

43.2% in the small molecule device to a maximum of 58.8% in the device made from
polymer structure 6A. This polymer had the highest overall efficiency at 0.95%, almost
double that of the small molecule at 0.55%, despite having lower hole and electron mobility
when compared to the small molecule. Although most of the molecular weights of the
polymers in this and subsequent studies
64
were low (< 10 kDa), a comparison study of the
molecular weight showed that it had no effect on device performance nor on film
formation. In these studies, the Li group effectively showed that connecting discrete
chromophores into long polymer chains (2-6 repeat units) enhanced film-forming ability,
leading to improved device performance, without negatively affecting the desirable
properties of the small molecule.

Figure 1.8. Polymer structure utilized in CBS studies by the Mei group, where x was varied
from 0 to 1 and n was varied from 0 to 20. (Refs. 84-88)
Since 2015, the Mei group has pursued the incorporation of conjugation break spacers
into a conjugated polymer backbone with the aim of improving polymer solubility and ease
of processing.
84
This group’s initial reports held true to their assumptions: increasing the
proportion of alkyl spacer (n = 3 in Figure 1.8) incorporated into the polymer backbone
led to an increase in solubility accompanied by a logarithmically correlated decrease in
charge mobility. However, with the higher percentages of alkyl spacer, melting transitions
were observed at low enough temperatures to consider melt processing. Films formed by
melt processing displayed charge mobilities approximately double that of their analogs
formed through solution processing (0.032 cm
2
V
-1
s
-1
vs 0.016 cm
2
V
-1
s
-1
organic field-
16

effect transistor (OFET) hole mobility for the perfectly alternating polymer with n = 3 in
Figure 1.8).
84
Melt processing is frequently practiced in industry and eliminates the need
for solvents altogether, reducing the environmental impact, toxicity, and cost of making
organic electronic devices. Subsequent studies showed that blending a fully conjugated
polymer (n = 0 in Figure 1.8) with a perfectly alternating CBS polymer (n = 3 in Figure
1.8) could preserve the charge mobility of the fully conjugated polymer while improving
its solution processability.
85
These blends (n = 0 with n = 5 in Figure 1.8) showed
promising OFET performance when melt-processed into relatively thick (1-2 µm) films.
86

Studies were also performed on varying lengths of alkyl spacer (n = 2-12,20 in Figure 1.8),
showing that while the length of the spacer had no effect on frontier orbital energy levels,
longer spacers corresponded to a logarithmic decrease in charge mobility, even when
blended with a fully conjugated polymer.
87
However, mechanical studies performed by the
Lipomi group on a family of these polymers ranging from fully conjugated to perfectly
alternating conjugation-break spacer polymers (n = 3 in Figure 1.8) showed that the
perfectly alternating polymer was surprisingly brittle and quite comparable to the fully
conjugated polymer.
88
The best mechanical properties were obtained from the polymer
with 70% of the conjugation break spacer incorporated (E = 0.103 GPa, crack-onset strain
(COS) = 12%) attributed to more disorder present in this polymer in contrast to the
perfectly alternating bookend polymers.
17

Figure 1.9. Polymers synthesized by the Bao lab of Stanford. In both structures 9 and 10,
x was varied between 0-20%. (Refs. 84,60)
Since 2016 an impressive but concise body of work in the field of conjugation-break
spacers has emerged from the lab of Zhenan Bao. In their initial study, the Bao group
incorporated 0-20% of a simple six-carbon alkyl conjugation-break spacer into an isoindigo
polymer with the aim of introducing disorder to the backbone of the polymer to prevent
aggregation in solution but allow good packing in film (Structure 9). They found that
though the solubility and processability of the polymers were enhanced with higher
incorporation of the conjugation-break spacer, the photovoltaic properties of this family
did not vary greatly.
89
Furthermore, high charge mobilities (between 0.1 and 2 cm
2
V
-1
s
-1
)
were maintained in OFET mobility measurements of the polymers with CBSs, though the
mobility decreased steadily with increasing CBS incorporation, as previously observed.
84

A subsequent study of a diketopyrrolopyrrole polymer (Structure 10) compared the simple
alkyl conjugation break spacer with a flexible pyridine-based linker capable of hydrogen
bonding, and its derivatives.
60
This work showed that not only were the flexible polymers
easier to process than their conjugated analogs, but they also had excellent mechanical and
18

electronic properties, with elastic moduli between 0.8 and 0.2 GPa and COS ≥ 25% for
CBS polymers. Significantly, the polymers with hydrogen bonding capabilities also
showed the ability to self-heal breaks in their films, restoring lost mechanical and electronic
properties. However, the distinction should be made between self-healing (in this case
requiring solvent- and thermal-annealing for optimum recovery) and autonomous healing,
in which no external stimulus is necessary to recover. More recently, the Bao group has
replaced the hydrogen bonding unit with a branched alkyl spacer for a 70-fold increase in
solubility in non-halogenated solvents when compared to its fully conjugated counterpart.
61

The four examples outlined above illustrate how the idea of using conjugation-break
spacers has developed in the past few years. Though in his initial study Sivula’s methods
were impractical and did not lead to well-defined conjugated segments, they opened the
door for future studies with more complex electroactive units than had been used before.
His study drew important relationships between film morphology, conjugation length, and
charge transport and showed the relative ease of using Stille polymerization to incorporate
CBSs. Li’s studies were successful in proving how he could achieve optoelectronic
properties that were independent from molecular weight. However, his claims would have
been aided by performing mechanical studies. It is likely that the molecular weight of his
polymers was too low for robust mechanical properties, a possible downside of using the
polyurethane/polyester condensation polymerization methods. The Mei and Bao studies
illustrate the potential of conjugation-break spacers for enhancing mechanical properties
and processing conditions while avoiding negative effects on electronic properties.
However, this is a delicate balancing act, for though mechanical properties improve when
more CBS is incorporated, electronic properties tend to deteriorate.
60,88

19

These studies lay the groundwork for our research applying conjugation-break spacers
to semi-random polymers. In all prior examples of polymers containing CBS units, the
polymers were synthesized with a perfectly alternating or semi-alternating
90
(random)
architecture, using AA and BB functionalized monomers. The restricted linkage pattern in
these polymers lends order to the polymer chain, which is propagated through the bulk and
manifests itself in the high degree of crystallinity often seen in perfectly alternating
polymers.
91,92
We predicted that the less restrictive linkage pattern available to monomers
in a semi-random polymer with AA, AB, and BB functionalized monomers can increase
disorder along the polymer backbone, a trait associated with decreased stiffness and
brittleness in polymer films.
93,94
The semi-random architecture has also been shown to
broaden the absorption of conjugated polymers over that of perfectly alternating polymers
by creating a broader range of chromophores within the polymer backbone than is
accessible with a perfectly alternating architecture.
17
Additionally, incorporating 3-
hexylthiophene allows for the retention of the favorable properties of P3HT while tuning
such properties as absorption, electronic energy levels, and surface energy.
18,95
Therefore
we believe that the semi-random architecture alone can provide the randomness necessary
to improve conjugated polymer mechanical properties as well as give us a handle for tuning
multiple properties simultaneously. The incorporation of conjugation-break spacers into
semi-random polymers is expected to open up new possibilities for mechanical properties
and optoelectronic tuning. These ideas led us to create several semi-random polymer
families with conjugation break spacers, varying spacer length and content (Chapter 2),
tuning monomer solubility (Chapter 3), and varying spacer content while holding the
electron-poor monomer content constant (Chapter 4).
20

1.4 Pendant Electroactive Polymers
1.4.1 Motivation/Background/History
Another class of polymer architecture with a great deal of potential for improving the
physical and mechanical properties of semiconducting polymers is that of non-conjugated
polymers with electroactive pendants. Comprising the plastics that we use in everyday life,
non-conjugated polymers are made through a wide variety of synthetic techniques, many
of which have been perfected over decades of use in industrial fabrication processes to give
precise end groups, molecular weights, and backbone stereochemistry. Early studies have
proven that charge transport can occur along π-π stacks,
68
and does not necessitate a
conjugated polymer backbone,
67,96
so if the electroactive pendants can arrange themselves
in a manner favorable for π-π stacking, charge transport should occur along the pendants.
Electroactive pendants can be added to a common monomer and polymerized directly
(Figure 1.10B), or can be added onto a non-conjugated backbone after polymerization via
post-polymerization modification (Figure 1.10A). Each method has its drawbacks: post-
polymerization modification rarely reacts to completion, leaving some units along the
backbone unreacted with no way of purifying the imperfections; direct polymerization
requires electroactive monomers and catalysts/initiatiors to be compatible. More optically
interesting monomers tend to be more sensitive to reaction conditions and can be difficult
to polymerize using controlled or living methods.
21

Figure 1.10. Schematic of two primary ways to make non-conjugated polymers with
electroactive pendants. A represents post-polymerization modification of a non-conjugated
backbone with an electroactive dye and B represents direct polymerization of a monomer
with an electroactive dye preattached.
Non-conjugated polymers can be made by living polymerization techniques to obtain
precise molecular weights, dispersity, and end groups, leading to less variation in the
properties that influence performance in organic electronic devices. In contrast to typical
conjugated polymer syntheses, transition metal catalysts are not necessary for many non-
conjugated polymerization methods. This not only lowers the synthetic cost but also allows
for biomedical applications and prevents metal-induced organic electronic device defects.
33

Non-conjugated polymers can be made both through step growth (condensation) and chain
growth (ionic, radical) polymerization techniques. The former produces stoichiometric
byproducts (usually water) and the latter has no byproduct. This is in sharp contrast to the
often toxic stoichiometric byproducts produced during conjugated polymer synthesis. Non-
conjugated polymers can frequently be synthesized in the bulk, eliminating the need for
toxic or wasteful organic solvents. They may also exhibit improved solubility and
mechanical properties over fully conjugated polymers, although these parameters are rarely
discussed in semiconducting polymer literature. Finally, the aliphatic backbone of the non-
22

conjugated polymers should be stable to photooxidation. This field has been explored with
varying frequency over the last several decades. Interestingly, some have even studied the
inverted architecture, with a conjugated backbone and non-conjugated pendants. A
thorough review on conjugated polymers with non-conjugated pendants was published by
Strover et al. in 2016.
97

1.4.2 Simple Electroactive Pendants

Figure 1.11. Simple electroactive pendant polymers with cofacial π-π stacking. A) Vinyl
carbazole and its polymer, poly(vinyl carbazole), or PVK. B) Dibenzofulvene and its
polymer, poly(dibenzofulvene), or pDBF.
Poly(vinyl carbazole) (PVK, Figure 1.11A) is one of the earliest examples of a pendant
electroactive polymer. Discovered to have photoconductive properties in 1957 by Helmut
Hoegl and coworkers,
98
PVK has been used for decades in electrophotography and
photocopiers.
99,100
More recently, the polymer has been investigated in photorefractive
materials and as a hole transport material for OLEDs.
101-103
Vinyl carbazole and its
derivatives are simple monomers, easy to functionalize and purify, and photostable when
polymerized. The monomer can be polymerized via anionic, cationic, and radical methods,
though the electronics of the vinyl bond are not conducive to metallocene or Ziegler-Natta
polymerization techniques.
104
The photoconductive nature of PVK derives from the fact
that it forms relatively stable radical cations upon photoexcitation, and the ability of the
23

pendants to π-π stack in a cofacial manner allows for efficient charge transport along the
pendants. A thorough review on pendant carbazole polymers was written by Grazulevicius
et. al in 2003.
99
Similarly, poly(dibenzofulvene) (pDBF, Figure 1.11B) has also been
targeted as a polymer with pendants that are π-π stacked and is stable to photooxidation
(though not to oxidative polymerization
105
).
106-108
PolyDBF was found to have higher
charge mobility than PVK (10
-4
vs 10
-7
cm
2
V
-1
s
-1
), comparable to that of a polymer with
a conjugated backbone.
109
However, although the polymer can be formed through cationic,
radical, and anionic polymerizations,
107
high molecular weights of poly(dibenzofulvene)
have never been achieved due to steric hindrance and poor solubility,
106,108,110
contrary to
the supposition that pendant polymers have better solubility than conjugated polymers.
111

Though these simple pendant polymers are electroactive and provide a solid foundation for
the study of electroactive pendant polymers, their high bandgaps (> 3.5 eV)
110,112
limit their
potential applications.
111

1.4.3 Metathesis Polymerizations and Memory Polymers

Figure 1.12. Schematic of ring opening metathesis polymerization (ROMP, A) and acyclic
diene metathesis polymerization (ADMET, B) with electroactive pendant dyes.
24

Much work has been done in the past two decades to affix electroactive pendants to
monomers that can be linked using metathesis polymerization methods. In particular, ring
opening metathesis polymerization (ROMP, Figure 1.12A) is known to be a living
polymerization with high functional group tolerance and mild reaction conditions, resulting
in polymers with narrow molecular weight dispersity and the potential to make block
copolymers.
113
ROMP proceeds rapidly with strained monomers such as norbornene,
shown in Figure 1.12A, and the resultant polymers are known to have good thermal
stability, solubility, and film-forming ability.
114,115
For this reason, norbornyl-dye
conjugates have been pursued by several groups to create electroactive pendant polymers
with oxidative stability,
116
cross-linkable pendants,
117
and more. A thorough review of dye-
functionalized polymers prepared via ROMP was published by Hollauf, et. al in 2015.
113

The Reynolds and Wagener groups briefly delved into acyclic diene metathesis
polymerization (ADMET, Figure 1.12B) as an alternative to ROMP for electroactive
pendant polymers in the early 2010s. They argued that ADMET could achieve perfectly
regioregular polymers and tailor the distance between pendants, whereas ROMP of
norbornyl-dye conjugates gave a statistical distribution of dyes spaced 4, 5, or 6 carbons
apart.
118
A later study sought to increase the distance between chromophores by increasing
the methylene units (m) in the structures in Figure 1.12B to prevent the fluorescence
quenching that usually accompanies highly ordered materials in OLEDS.
119
They found
that increasing the spacer length between pendants did indeed increase the fluorescence
quantum yield by suppressing quenching, but that longer spacers also led to a decrease in
current density in OLED devices, attributed to a decrease in charge mobility. A final study
found that they could tune the hole mobility over three orders of magnitude by altering the
25

distance between chromophores (10
-10
cm
2
V
-1
s
-1
for electroactive pendants on every 21
st

carbon to 10
-7
cm
2
V
-1
s
-1
for electroactive pendants on every 9
th
carbon).
120
These
mobilities are quite low, such that it is difficult to truly classify the polymers as
semiconducting. Likely the mobilities would improve if the spacers could be shortened,
but for molecules with fewer than three methylene units (m) on each side of the
electroactive pendant, ring closing metathesis (RCM) was a competing reaction. In addition
to the drawback of having RCM as a competing reaction, ADMET requires polymerization
under vacuum to remove the ethylene gas side product and drive the reaction towards
completion, a challenge for polymerizations that require solvent. And unlike the chain
growth that characterizes ROMP, ADMET is a step growth polymerization, meaning that
very high conversions must be achieved to attain high molecular weight polymers, and
control over molecular weight and dispersity is limited.
Yet another strategy involving precise placement of electroactive pendants along a
thermally and chemically stable non-conjugated polymer backbone is that used in electrical
memory devices.
121
This method combines the physical and mechanical properties of
polyimides, polyamides, polyureas, polysulfones, etc. with the optoelectronic properties of
pendants. A thorough review on the use of electroactive pendants with non-conjugated
polymers for electrical memory devices was published by Ree et. al in 2015.
121
All three
methods outlined above provide excellent control over the polymerization, leading to
polymers of precise molecular weights and low molecular weight dispersities. This precise
control could lead to organic electronic devices that do not suffer from batch-to-batch
variation. However, all three methods suffer from poor charge mobility due to the distance
between the electroactive units; charge is best transported through π-π stacks, which these
26

polymers do not appear capable of achieving. Another downside is that all of the examples
covered thus far are limited to hole donors and transporters, or p-type materials.
121
Many
organic electronics, including OLEDs, OPVs, and organic memory devices rely on the
interplay of both hole and electron transporters (p- and n-type materials) for efficient
function. More research in this field is needed to develop pendant electron
donors/transporters.
1.4.4 Polymers with High Chromophore Density
Several strategies have been pursued to shorten the distance between electroactive units
to promote charge transport in pendant polymers. As illustrated in Figure 1.10, this can
take the form of either post-polymerization modification with pendants (grafting-on
strategy) or the direct polymerization of monomers with electroactive pendants. In 2010,
Lang et al. published a comparative study of polyacrylates with perylene diimide pendants,
synthesized either through a direct controlled radical polymerization of the perylene
acrylate monomer (Structures 12-13) or through a post-polymerization “click” reaction of
a perylene azide onto a poly(propargyl acrylate) (Structure 11).
122
Using nitroxide-
mediated radical polymerization (NMRP) in the direct polymerization of the perylene
acrylate monomer, they showed that the polymerization kinetics were not first order, and
the molecular weight dispersity increased with monomer conversion, neither of which
phenomena were observed in typical acrylate NMRPs. The authors attributed the broad
dispersity and unpredictable molecular weight of these polymers to monomer’s poor
solubility, steric hindrance, high molecular weight, and transfer reactions occurring during
polymerization. Despite this, they were still able to achieve relatively high polymer
molecular weights, with number-average molecular weights above 10 kDa. This signifies
27

a degree of polymerization of 18-19 for structure 11, 16-17 for structure 12, and 21-22 for
structure 13. The degree of polymerization for structure 11 is much lower than the expected
value of 30-36 repeat units based on the Mn of the propargyl acrylate polymer before
“clicking” on the perylene azide. The authors attribute this to an underestimation of the
“clicked” polymer mass due to aggregation, instead of incomplete functionalization
leading to fewer than expected perylene azides “clicking” onto the polymer backbone.
However, assuming that each step reacts to completion, the post-polymerization
modification should lead to more uniform pendant polymers with narrow molecular weight
distributions.

Figure 1.13. Comparison of electroactive pendant polymers made by the grafting-on
strategy (post-polymerization modification, 11) against those made by direct
polymerization (12 and 13). Ref. 122.
The example highlighted above is illustrative of the challenges facing this field.
Electroactive polymers made via the grafting-on technique frequently suffer from low
charge mobilities due to low grafting density. This can be attributed to incomplete
conversion (intentional or otherwise) of functional groups along the backbone to pendant
moieties. Alternatively, the electroactive moiety may be diluted with another monomer to
28

attain some other desirable property such as elasticity, self-assembly,
123
crosslinking, water
solubility,
124
or a desired weight ratio of chromophores.
125
On the other hand, direct
polymerization of an electroactive pendant monomer may lead to a greater density of
chromophores along the polymer backbone, but often a lower molecular weight due to poor
solubility. Additionally, fewer synthetic options may be available in terms of both
functionalizing the chromophore with a polymerizable group as well as the polymerization
itself based upon the complexity, solubility, and reactivity of the electroactive group used.
For example, in the study highlighted above, a pre-clicked acrylate monomer was
synthesized, but could not be polymerized due to its poor solubility.
122
The decision to
pursue an electroactive pendant polymer from the grafting-on approach or direct
polymerization must be made on an individual basis and is informed both by the type of
monomer and pendant chosen and the properties desired from the final polymer product.
Despite the inherent challenges, non-conjugated electroactive pendant polymers have
been pursued for a variety of applications, from memory devices to OLEDs to OPVs. In
2009, the Hirao and Chen groups synthesized polymers with oligofluorene pendants
through living anionic polymerization with the aim of attaining good solubility and film
formation along with precise optoelectronic control.
126
They opted for the direct
polymerization method because the oligofluorene monomers were presumed to be stable
to anionic polymerization, and they attained polymers with very low dispersity (< 1.10)
and molecular weights that closely matched the predicted values. That same year, the team
put the pendant polymers into non-volatile memory devices and showed that the memory
characteristics could be systematically tuned with the conjugation length of the pendant.
127

The same polymerization procedure would likely not have worked for a more complex
29

monomer than fluorene. On the other hand, the Fréchet group has combined direct
synthesis with post-polymerization modification to create a single-component, solution
processable polymer with a photoinert backbone for OLEDs.
33
They performed a random
copolymerization of hole transporting, electron transporting, and metal chelating
monomers using NMRP, then used a post-polymerization modification to add in the metal
emitter complex. By varying the content of each monomer, they were able to tune the
single-component active layer to white emission. Several studies have focused exclusively
on oligothiophene pendants, for much the same reasons pursued in the first generations of
CBS polymers above. The Reichmanis group used post-polymerization Stille
functionalization (~96% yield) and coupling (unquantifiable yield) to create pendant
oligothiophenes on a polystyrene backbone.
128
Others have pursued direct polymerization
of pendant thiophenes, but have never been able to achieve high molecular weights.
129-133

A thorough review of electroactive polymers with thiophene pendants was published by
Qiao et al. in 2015.
134

A great deal of research has been done to develop block copolymers: with a non-
electroactive polymer to gain desired physical or mechanical properties;
135
with a
conjugated electroactive polymer in so-called “rod-coil” or “hard-soft” blocks;
136
or with
another non-conjugated electroactive pendant block, taking the form of “coil-coil” or “soft-
soft” blocks.
137
A thorough review primarily focused on hard-soft electroactive block
copolymers was published by McCullough and Matyjaszewski in 2010.
138

30

Figure 1.14. Typical nanomorphologies formed in thermodynamic equilibrium from coil-
coil block copolymers. Going from left to right, the relative fraction of the red monomer is
decreasing as the fraction of the black monomer increases. Reprinted with permission from
Marencic, A. P.; Register, R. A. Controlling Order in Block Copolymer Thin Films for
Nanopatterning Applications. Annual Review of Chemical and Biomolecular Engineering
2010, 1, 277-297. Copyright 2010 Annual Reviews.
Block copolymers are an area of focus because not only can multiple functionalities be
encoded into the two (or more) blocks, but they are also known to self-assemble into
nanomorphologies based upon the relative monomer fraction (Figure 1.14).
139-141
This
feature makes block copolymers an attractive option, and they have been studied
extensively for use in OPV active layers since the Hadziioannou group first proposed in
2000 that block copolymers of electron donating and electron accepting monomers could
be programmed to phase separate on a length scale conducive for exciton diffusion.
142

However, rod-rod and rod-coil block copolymers do not separate into the morphologies
predicted for coil-coil polymers due to the physics of the stiff, conjugated polymer
segment. For this reason, the Thelakkat group has pursued coil-coil donor-acceptor block
copolymers, showing the first evidence of thermodynamic phase segregation in 2004
(Figure 1.15),
137
and demonstrating that the single-component block copolymer
outperforms a comparable 2-component blend in OPV devices.
143
Two thorough reviews
on the subject of donor-acceptor block copolymers to improve morphological control and
enhance long-term device stability in OPVs have been written by the Thelakkat and Hiorns
groups.
144,145

31

Figure 1.15. (A) Coil-coil-type block copolymer with hole-transporting triphenylamine
pendant and electron-transporting perylene diimide blocks and (B) a TEM cross section of
a film showing thermodynamic phase segregation of the block copolymer. (B) is reprinted
with permission from Lindner, S. M.; Thelakkat, M. Nanostructures of N-Type Organic
Semiconductor in a P-Type Matrix Via Self-Assembly of Block Copolymers.
Macromolecules 2004, 37, 8832-8835. Copyright 2004 American Chemical Society.
As phase segregation of electroactive components is not only useful for OPVs, pendant
block copolymers have also been pursued for non-linear optics devices,
146
OLEDs,
36,147

and memory devices,
148
where they were found to exhibit favorable morphologies.
Through all of these studies, it has been shown that electroactive pendants on non-
conjugated polymer backbones can achieve desired morphology and stability. However, it
remains a challenge to synthesize optoelectronically interesting polymers with precise
molecular weights and low dispersities. In the vast majority of pendant polymers produced,
the charge mobility continues to suffer relative to that observed in fully conjugated
polymers. One exception to this general observation is an n-type perylene diimide pendant
polymer which was shown to have electron mobility as high as 1 E-2 cm
2
V
-1
s
-1
when
measured by SCLC.
149
This is four orders of magnitude better than the mobility measured
in OFET devices, and is comparable to or better than many n-type conjugated polymers.
150

The remarkable mobility is attributed to how the pendants arrange themselves in relation
to direction of current travel in the different device architectures.
32

It is clear from the examples illustrated above that charge mobility is related to
structure, and that semiconducting polymer architecture can be leveraged to improve
mobility, processability, and morphology, amongst other properties. The lessons we have
taken from studies on electroactive pendant polymers is that we should aim for high
grafting density (for mobility), high molecular weight (for film properties), and
optoelectronically interesting monomers (for variety of applications). Another avenue that
remains relatively unexplored is the use of backbone stereoregularity to control pendant
alignment and influence mobility.
101,104,151
This topic is explored further in Chapter 5 of
this dissertation.
1.5 Conclusion
In sum, there is a great variety of architectures available for semiconducting polymers,
each with their own advantages and disadvantages. However, when seeking to tune
polymer mechanical properties, charge mobility, and processability simultaneously, the
conjugation-break spacer and electroactive pendant architectures hold the most promise.
This dissertation will explore the synthesis and characterization of novel families of
semiconducting polymers made using either of these strategies and the impact of the
polymer structure on its optoelectronic and mechanical properties.
33

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47

Chapter 2: Influence of Systematic Incorporation of Conjugation-Break Spacers
into Semi-Random Polymers on Mechanical and Electronic Properties

2.1 Introduction
Organic electronic materials have been pursued for several decades with a vision of
creating devices that are lightweight, flexible, and even stretchable to be suitable for a
broad variety of applications. Conjugated polymers have many characteristics that make
them attractive for these applications,
1
however, they tend to be brittle, inflexible, and have
limited solubility and high melting points, making it impractical to incorporate them into
industrial production. The rigid π–conjugated backbone of these materials, combined with
the semicrystalline microstructure and the (often) glassy state of the amorphous domains
at the operating temperature, has the effect of decreasing mechanical deformability and
robustness. Different strategies have been pursued to try to overcome these mechanical and
physical limitations, including side-chain engineering,
2,3
nanoconfinement,
4,5
blends with
non-conjugated polymers,
6,7
and the creation of rod-coil-type block copolymers with
conjugated and non-conjugated segments.
8,9
Many of these approaches dilute the
electroactive component with insulating hydrocarbons, leading to a decrease in charge
mobility. Additionally, several of these strategies require complex synthetic and processing
conditions, increasing their ultimate cost. Thus the desire arises to create simple, single-
component semiconducting polymers with intrinsically robust mechanical properties.
10

Over the last few years, a strategy from the 1990s
11-13
of incorporating non-conjugated
segments into a conjugated polymer backbone has been revived with new versatility lent
by improved polymerization techniques and a large library of electroactive monomers. An
48

increasing number of groups have begun to study semiconducting polymers with a break
in conjugation, seeking properties such as melt processibility,
14
stretchability,
15
and even
healable materials
16
by introducing additional modes for dissipation of mechanical energy
(e.g., stretching of and rotation about aliphatic units).
17-21
The non-conjugated segments
within the polymers are commonly known as conjugation-break spacers, or CBSs. The Bao
group has succeeded in making polymers that could almost fully recover their OFET charge
mobility (1.13 vs. 1.28 cm
2
V
-1
s
-1
) after healing a cut in the film by incorporating
conjugation-break spacers capable of hydrogen bonding. The Mei group has synthesized a
family of polymers with CBSs that so drastically lowered the polymer melting point (from
221 °C to 94 °C as spacer length increased) that it allowed for melt processing of thin films,
eliminating the need for wasteful, toxic organic solvents. Previous work has suggested that
mechanical properties such as the elastic modulus (E) depend on a complex interplay
between molecular structure and packing arrangement in the solid state.
15
A lower elastic
modulus is generally considered to be better for flexible electronics applications, as it
indicates a material which is more easily deformed and which produces decreased
interfacial stresses with other layers in a device. While elastic moduli of conjugated
polymers vary widely (E = 0.1-8 GPa),
3,16,22-24
polymers with CBSs have consistently
shown elastic moduli of less than 1 GPa.
15,16
However, they fall far short of conventional
elastomers such as PDMS (E = 0.6-2.5 MPa)
25
or polyisoprene (E = 0.36 MPa).
26

In previous examples of polymers containing CBS units, polymers were synthesized
with a perfectly alternating or semi-alternating
27
(random) architecture, using AA and BB
functionalized monomers.
21,28,29
The restricted linkage pattern in these polymers lends
order to the polymer chain, which is propagated through the bulk and manifests itself in
49

the high degree of crystallinity often seen in perfectly alternating polymers.
30,31
Herein, we
introduce a semi-random architecture to polymers incorporating conjugation-break
spacers, using AA, AB, and BB monomers. The less restrictive linkage pattern available to
monomers in the semi-random polymer can increase disorder along the polymer backbone,
a trait associated with decreased stiffness and brittleness in polymer films.
32,33
The semi-
random architecture has also been shown to broaden the absorption of conjugated polymers
over that of perfectly alternating polymers by creating a broader range of chromophores
within the polymer backbone than is accessible with a perfectly alternating architecture.
34

Additionally, incorporating 3-hexylthiophene allows for the retention of the favorable
properties of P3HT while tuning such properties as absorption, electronic energy levels,
and surface energy.
35,36

In this study, a family of semi-random polymers with conjugation-break spacers was
synthesized via Stille polymerization (Scheme 2.1). CBSs were incorporated at 10%, 20%,
30%, and 40%, with spacer length varying between 4 and 10 methylene units. Due to the
need to balance the AA/BB functional groups, the amount of the electron poor
diketopyrrolopyrrole (DPP) acceptor monomer was increased at the same rate as the CBS,
while simultaneously decreasing the amount of 3-hexylthiophene (3HT) monomer.
Introducing randomness into the backbone of a conjugated polymer has been shown to
decrease elastic modulus and increase crack-onset strain.
23
It was expected that the semi-
random nature of the polymer backbones in this study would have a similar effect. The
degrees of freedom added by conjugation-break spacers were expected to amplify these
trends. As more methylene units were introduced into the spacer, we anticipated a decrease
in elastic modulus (the ability of the material to resist deformation when a stress is applied
50

to it) and ultimate tensile strength (UTS, the maximum stress that a material can withstand)
due to the increased modes of dissipating mechanical energy excluding fracture.
37,38
In
addition to probing the influence of structure on mechanical properties, we also aimed to
expand our understanding of the scope of structure-function relationships in the versatile
class of semi-random polymers.
2.2 Results and discussion

Scheme 2.1. Semi-random polymers with conjugation-break spacers synthesized for this
study
The family of sixteen polymers with varying contents of CBS (and consequently DPP
and 3HT monomers) and the fully conjugated reference polymer, P3HTT-DPP-10%, were
synthesized via Stille polycondensation using methods previously developed in our
group.
34
Solely even-numbered hydrocarbon chains were studied both for economic
practicality and to eliminate a possible odd-even effect.
39
Polymers were named by the
CBS monomer and the amount by which it was incorporated, e.g., 10% T-4-T indicates a
CBS length of 4 carbons and that monomer makes up 10 mole percent of the polymer
backbone, along with 10 mole percent of DPP and 80 mole percent 3HT. Polymer
51

composition was confirmed by NMR to match the monomer feed ratios (Figure A2,
Figures A3-A6). As increasing amounts of the alkyl spacers were incorporated into the
polymer backbone, solubility in common halogenated solvents decreased, contrary to
trends observed by Zhao et al.
21
Solubility is necessary not only for purifying and analyzing
the polymers, but also for processing them in their eventual use in organic electronics. The
decrease in solubility was evidenced not only by increased difficulty with processing the
polymers, particularly those with shorter spacers, but also the regular decrease in polymer
molecular weight as the content of CBS increased (Table 2.1, Table A1, Figure A1). This
could be due to the fact that DPP content was increased at the same rate as spacer content,
with the poorly soluble acceptor monomer perhaps having a contradictory effect from that
of the CBS. This effect diminished with the 10-carbon alkyl spacer. We suspect that the
solubility-enhancing effects of the ten-carbon spacer outweighed the negative impact DPP
had on solubility. Specifically, when 40% of the T-10-T monomer was incorporated into
the polymer, a higher molecular weight was obtained than when only 30% of the ten-carbon
spacer was used, contrary to the trend seen with every other spacer length. The entire T-
10-T subfamily was easily soluble in chloroform at room temperature in contrast to the T-
4-T subfamily. The polymer incorporating a four-carbon spacer at 40%, and thus DPP at
40%, required a chlorobenzene Soxhlet extraction to obtain a higher molecular weight
(16.9 kDa) polymer fraction, but this fraction was not used for any subsequent analysis due
to its poor solubility. These were generally high-yielding polymerizations, giving
chloroform-soluble fractions at 60%-80% yield (Table A2).

52

Table 2.1. SEC, thermal, and electronic data for semi-random polymer family
Polymer Mn
a
(kDa)
Ð
a
Tm / Tc
b
(°C)
Eg
c
(eV)
HOMO
d
(eV)
µh
e
(cm
2
V
-1
s
-1
)
P3HTT-DPP-10% 9.5 4.15 208 / 182 1.50 5.52 9.29 E-4
10% T-4-T 10.5 4.30 - / -
f
1.54 5.45 -
g

10% T-6-T 29.0 4.79 - / -
f
1.56 5.51 2.90 E-6
10% T-8-T 19.7 6.23 - / -
f
1.55 5.41 2.08 E-5
10% T-10-T 14.0 3.68 - / -
f
1.53 5.43 2.53 E-4
20% T-4-T 8.4 4.90 - / -
f
1.58 5.41 -
g

20% T-6-T 17.6 4.28 - / -
f
1.58 5.41 -
g

20% T-8-T 14.2 5.15 - / -
f
1.61 5.49 6.49 E-6
20% T-10-T 12.4 5.24 - / -
f
1.60 5.44 1.06 E-5
30% T-4-T 6.4 2.87 - / -
f
1.63 5.42 -
g

30% T-6-T 14.4 5.42 - / -
f
1.65 5.40 -
g

30% T-8-T 9.9 5.15 90 / 81 1.66 5.43 -
g

30% T-10-T 9.8 6.47 75 / 73 1.64 5.48 3.55 E-6
40% T-4-T 7.4 2.90 - / -
f
1.70 5.43 -
g

40% T-6-T 10.5 5.09 143 / 112 1.68 5.41 -
g

40% T-8-T 8.8 2.69 125 / 85 1.68 5.45 -
g

40% T-10-T 12.2 2.98 83 / 76 1.68 5.50 -
h

a) Obtained via size-exclusion chromatography (SEC) versus polystyrene standards; b)
Obtained via differential scanning calorimetry (DSC); c) Calculated from the absorption
band edge in thin films, where Eg = 1240/λedge; d) Estimated from cyclic voltammetry (CV)
oxidation onset versus ferrocene; e) Calculated from SCLC mobility measurements in hole-
only devices with the architecture: ITO/PEDOT:PSS/polymer/Al, where the polymer layer
was spin-cast from chloroform and annealed for 30 min at 150 °C prior to aluminum
deposition; f) No apparent thermal transitions; g) Hole mobility not measured due to
difficulty obtaining uniform films; h) Not enough sample for annealed test; unannealed
data in Appendix A.

2.2.1 Optical properties
The UV-vis spectra of the T-8-T subfamily are shown in Figure 2.1 as a representative
example of optical trends. Spectra for all of the subfamilies are provided in Figures A7-
53

A14. All polymers with a break in conjugation had a blue-shifted absorption onset, leading
to a wider optical bandgap, when compared to the fully conjugated P3HTT-DPP-10%
optical bandgap of 1.50 eV (Table 2.1). As the percentage of CBS and DPP incorporated
into the polymer increased, the absorption onset continued to blue-shift to a maximum
bandgap of 1.70 eV for the 40% T-X-T subfamily (Figure A15). Both of these trends are
as expected, for the extended conjugation of a semiconducting polymer allows the highest
occupied and lowest unoccupied molecular orbitals to draw nearer to each other, lowering
the bandgap of the extended solid. By interrupting this extended sp
2
hybridization with the
sp
3
hybridized carbons of alkyl chains, the bandgap increased and therefore the absorption
onset moved to a higher energy. This is in contrast to a previous study from our group in
which the absorption onset was found to red-shift and the bandgap to decrease with
increasing DPP acceptor content in fully conjugated semi-random polymers.
40
It is
interesting to consider that perhaps the DPP and CBS monomers had contradictory effects
on optical properties as well.

Figure 2.1. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast from
o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. The P3HTT-DPP-
10% absorption spectrum (dashed black line) is provided for a fully conjugated reference.
(■) 10% T-8-T; (●) 20% T-8-T; (▲) 30% T-8-T; (▼) 40% T-8-T.

54

Only the percentage of conjugation-break spacer incorporated seemed to affect the
absorption onset; the length of the CBS had no effect on the onset. However, experiments
with different solvent and annealing conditions for the T-6-T subfamily indicated that there
was a strong dependence of optical properties on processing conditions (Table A5, Figure
A16), which has been observed previously in semiconducting polymers.
41
From these
experiments it is seen that both aromatic solvents and thermal annealing slightly increased
the bandgap of the T-6-T polymers. As expected, absorption in the P3HT region of 400-
550 nm decreased as CBS and acceptor content increased due to a consequent decrease in
loading of the 3-hexylthiophene monomer (Figure 2.1).
2.2.2 Structural properties
Conjugated polymers are known for their high melting and crystallization
temperatures, which were apparent in the fully conjugated P3HTT-DPP-10% polymer
when analyzed by differential scanning calorimetry (DSC) (Figure A26). These thermal
transitions can likely be attributed to crystallized P3HT segments present within the
polymer, for P3HT has a melting point of 212 °C,
40
very close to the P3HTT-DPP-10% Tm
of 208 °C. In contrast, no thermal transitions were observed in the semi-random polymers
with lower incorporation of CBS, even though the 10% T-X-T subfamily has the same
P3HT content as the fully conjugated P3HTT-DPP-10% polymer. This lack of thermal
transitions is presumably due to the additional disorder contributed to the semi-random
structure by the addition of the CBS. Melting and crystallization transitions became evident
at higher percentages of CBS and DPP incorporation, but only for longer spacers (e.g., 30%
T-8-T, 30% T-10-T, 40% T-6-T, 40% T-8-T, and 40% T-10-T; see Tables A7-A8 and
Figures A32-A33). The temperature of these thermal transitions decreased with increasing
55

spacer length, as had been observed previously by Zhao et al.
39
This trend is attributable to
a corresponding increase in polymer backbone flexibility, which lowered the melting and
crystallization temperatures by increasing the entropy associated with those transitions.
An attempt to corroborate the DSC findings was made by observing crystallinity in thin
films via grazing incidence X-ray diffraction (GIXRD). However, contrary to the findings
above, within the CBS polymers, peaks were visible primarily for the 10% T-X-T
subfamily (Figure A34, Table A9). This is attributable to the high P3HT content in that
family, as the 100 diffraction peak in P3HT corresponds to a lamellar packing distance of
16.7 Å,
42
which nearly matches the average distance of 16.3 Å for the 10% T-X-T
subfamily. The fully conjugated, semi-crystalline P3HTT-DPP-10% polymer packed
closer at 15.7 Å, as previously observed.
35
Interestingly, the intensity of the diffraction
peak decreased steadily as the spacer length increased across the 10% T-X-T subfamily,
with the exception of 10% T-10-T, which showed a sharp increase, exhibiting even more
crystallinity than the fully conjugated P3HTT-DPP-10% polymer. It could be that when
the spacer reached ten methylene units in length, it had sufficient flexibility to allow better
intra- and inter-molecular packing, leading to a more intense diffraction peak. This
improved packing could also explain why the 20% T-10-T polymer was the only one in
the 20% T-X-T subfamily to exhibit a diffraction peak. The 20% T-10-T peak was likely
much less intense than its 10% T-10-T cousin due to the 20% decrease in 3HT monomer
content between the two, though it appeared to have approximately the same lamellar
packing distance (16.0 Å). A small diffraction peak was observed in the same region for
the 40% T-6-T polymer, but it was not distinct enough to obtain a Gaussian fit.

56

2.2.3 Electronic properties
The highest occupied molecular orbital (HOMO) energy was calculated from oxidation
onset using cyclic voltammetry (CV). As can be seen in Table 2.1 (and Figure A25), there
were no observable trends in the HOMOs of these polymers. All values were
approximately the same, producing an average HOMO of 5.44 ± 0.04 eV (within the
instrument error of ± 0.05 eV). This was slightly more shallow than the 5.52 eV HOMO of
the fully conjugated P3HTT-DPP-10%. This trend (or lack thereof) closely matches that
seen by Zhao et al., where they observed no effect on the HOMO from increasing the CBS
length.
39
However, in their study, the polymers were perfectly alternating between DPP
and spacer monomers, maintaining the same electroactive moiety regardless of spacer
length. Therefore there was no change in the effective conjugation length as spacer length
was varied, which would lead one to assume that there would be no change in frontier
orbital energy levels. In contrast, in this study, the amount of spacer as well as its potential
distribution within the polymer was varied, which makes the consistent HOMO energy
more surprising.
To determine if the spacer length and content influenced charge mobility, SCLC
measurements were performed on hole-only devices. However, due to the poor solubility
of some of the polymers and the resulting difficulty in obtaining uniform films, only 8 of
the 17 polymers were measured (Table 2.1, Tables A10-A13). Annealed films spin-cast
from chloroform are presented herein; films cast from ortho-dichlorobenzene showed no
trends and are presented in Appendix A (Table A11), along with as-cast films from
chloroform (Tables A10 and A12; Figure A35). From the data points obtained, it can be
seen that mobility decreased as the content of spacer increased in the T-10-T subfamily.
57

This correlates with observations made in the literature for mobility values decreasing as
spacer content increased due to a larger amount of insulating alkyl chains interrupting
charge transport along the polymer backbone.
21
However, a counterintuitive trend emerged
within the 10% T-X-T (and, to a lesser extent, the 20% T-X-T) subfamily. In these
measurements, the hole mobility appeared to increase as the length of the CBS increased,
contrary to trends observed by Zhao et al.
39
This is counterintuitive because the polymers
with longer spacers have more insulating hydrocarbon units per mole than those with
shorter CBSs. However, we reason that the longer spacers possibly gave the polymer
increased flexibility and more easily allowed the electroactive portions of the polymer
chain to rearrange and stack themselves in a manner favorable for charge transport. This
corresponds well with the highest charge mobility coming from the 10% T-10-T polymer
(2.53 E-4 cm
2
V
-1
s
-1
), which also had the most intense diffraction peak, indicating higher
crystallinity due to improved packing. Remarkably, this value for hole mobility is the same
order of magnitude as that obtained for the fully conjugated P3HTT-DPP-10% analog (9.29
E-4 cm
2
V
-1
s
-1
), signifying that a 10% incorporation of conjugation-break spacer did not
greatly impede charge mobility.
2.2.4 Mechanical properties
To test our hypothesis about semi-random CBS incorporation contributing to improved
mechanical properties, film-on-elastomer mechanical tests were performed on the entire
family of 16 polymers, as well as the fully conjugated P3HTT-DPP-10% polymer. Several
interesting trends were observed (Table 2.2, Table A14, Figure A37). The elastic
(Young's) modulus, E, is the ratio of stress σ, or force per cross sectional area of a deformed
body, to strain ε, or fractional amount of deformation in the direction of applied force, in
58

the linear region of a stress-strain curve.
1
The modulus is a quantitative measure of the
ability of a material to store mechanical energy reversibly, or its elasticity. This property
is strongly related to the solid-state morphology, as well as the molecular structure, of the
material.
15
The more closely polymer chains can pack, the greater the density of load-
bearing carbon-carbon bonds along the strained axis and the greater the intermolecular
forces between chains. As aforementioned, a lower elastic modulus is generally considered
to be favorable for applications involving flexible electronics, as it corresponds to a more
compliant material. The elastic modulus in our polymers increased when the spacer content
was increased from 10% to 20% incorporation for all spacer lengths and most of the 40%
T-X-T subfamily were unattainable due to poor film-forming ability. It is noteworthy that
the elastic modulus of 0.32 GPa for the fully conjugated semi-random P3HTT-DPP-10%
polymer is itself in the lower range of moduli observed in perfectly alternating fully
conjugated polymers.
16,22,23,43
It is interesting that the 10% CBS subfamily exhibited
improved elastic moduli compared to the fully conjugated analog, but this trend was
reversed for the 20% CBS subfamily. This could be due to the aforementioned
contradictory effect introduced by increasing the ratio of the stiff DPP monomer along with
the flexible CBS monomer. With the exception of 20% T-4-T, this entire polymer family
fit the trend of CBS polymers consistently displaying elastic moduli < 1 GPa.

59

Table 2.2. Mechanical properties obtained from film-on-elastomer measurements for
semi-random polymer family
Polymer
Elastic Modulus
(GPa)
a

Crack-Onset Strain
(%)
b

Mode of
Failure
b

P3HTT-DPP-10% 0.32 ± 0.20 10 Brittle
10% T-4-T 0.33 ± 0.11 27 Ductile
10% T-6-T 0.15 ± 0.02 > 80
c
Ductile
10% T-8-T 0.14 ± 0.06 > 80
c
Ductile
10% T-10-T 0.15 ± 0.04 > 80
c
Ductile
20% T-4-T 1.30 ± 0.42 2 Brittle
20% T-6-T 0.51 ± 0.16 37 Ductile
20% T-8-T 0.65 ± 0.32 > 80
c
Ductile
20% T-10-T 0.52 ± 0.12 > 80
c
Ductile
30% T-4-T 0.15 ± 0.09 < 1 Brittle
30% T-6-T -
d
-
d
-
d

30% T-8-T 0.59 ± 0.15 < 1 Brittle
30% T-10-T 0.60 ± 0.07 > 80
c
Ductile
40% T-4-T -
d
1 Brittle
40% T-6-T -
d
-
d
-
d

40% T-8-T -
d
1 Brittle
40% T-10-T 0.60 ± 0.30 77 Ductile
a) Calculated using the buckling-based metrology and averaged over three measurements;
b) Obtained from optical micrographs; c) Test terminated at 80% strain due to potential for
PDMS substrate breakage; d) Data unattainable due to inability to form uniform films.

Crack-onset strain (COS), or the strain at failure (when cracks first appear in the
polymer film) was also obtained from film-on-elastomer measurements. This value is a
direct measurement of film ductility and provides important insight about film
stretchability.
10,44
There were clear trends both in the COS and mode of failure as spacer
length and percent incorporation increased (Table 2.2, Table A15, Figure A38, Table
60

A17). Ductile films failed by forming pinholes (Figure A40), while brittle films failed by
forming parallel, slender cracks (Figure A39). All of the 10% T-X-T subfamily and the
entire T-10-T subfamily exhibited ductile behavior, while the higher incorporation of CBS
and shorter spacers exhibited brittle behavior. This correlated with the COS increasing as
the alkyl spacer became longer and decreasing as more CBS was incorporated into the
polymer chain. Again, the general decrease in ductility with increasing CBS incorporation
is attributed not to the alkyl spacer itself, which was expected to improve mechanical
properties, but to the corresponding increase in DPP monomer incorporation. Notably, the
COS of 10% for the fully conjugated P3HTT-DPP-10% polymer is at the higher end of the
range typically observed for fully conjugated perfectly alternating polymers.
15,16,22
This is
attributable to the more randomized monomer distribution along the backbone attainable
with semi-random polymers, leading to decreased order.
45
It should be noted that several
of the COS values obtained for this polymer family are extraordinarily high (> 80%). In
measuring COS, we terminated the film-on-elastomer tensile test beyond 80% applied
strain due to breakage of the PDMS substrate. For this reason, further measurements of
free-standing films were pursued using an alternative technique for the seven polymers
exhibiting a COS near or greater than 80%.
The mechanical properties of the seven polymers were further investigated using a
method originally developed by Kim and co-workers, where a tensile test is performed on
a film supported by water (Figure 2.2).
46
This “film-on-water” technique is a modified
version of a conventional pull test (which involves suspending the specimen in air). The
film is able to float and slide unimpeded on the surface due to the high surface tension and
low viscosity of water. The mechanical properties obtained by this method can be
61

considered a “definitive” measurement of the inherent polymer properties, whereas those
obtained through film-on-elastomer measurements are a more realistic approximation of
mechanical behavior in devices.
47
Stress-strain curves acquired for this study were
produced by transforming the obtained curves of force versus displacement into stress-
strain curves using the dimensions of the corresponding sample. From these curves were
obtained the elastic modulus, toughness, ultimate tensile strength (UTS), and fracture
strength, with the results tabulated in Table 2.3. Toughness is the amount of energy per
unit volume that a material will absorb before fracturing. Since toughness is dependent on
the area under the stress-strain curve, an increase in strength or extensibility of the material
will improve the toughness of the material. There is also a general trend suggesting that
greater intermolecular forces between polymer chains will improve the toughness.
1
UTS
reflects the maximum stress that the material withstands prior to fracture, and fracture
strength is the stress at which the material fractures. The extensibility, or the fracture strain,
is the strain at which the material fractures. Sometimes this quantity is called the
“stretchability,” though we do not generally support the use of this word when referring to
mechanical properties because it has different meanings to different communities.
48

62

Figure 2.2. Pseudo free-standing tensile test of P3HT-based semi-random copolymers.
Schematic representation of experimental setup, which includes a floating film, linear
actuator, load cell, and trough filled with water. Inset photograph demonstrates the real
experimental apparatus

Table 2.3. Tabulated values of mechanical properties using the film-on-water technique
a) Elastic moduli can be derived from stress–strain curves by taking the slope of the linear
regime of the graph; b) Toughness values are obtained by integrating the entirety of the
stress–strain curve; c) UTS is obtained from the maximum stress in the stress–strain curve;
d) Fracture strength is obtained from a stress–strain curve by reading the stress at failure.

Trends within the film-on-water data are plotted in Figure 2.3, where 2.3a-2.3c show
the effects of spacer length and 2.3d-2.3f demonstrate the effects of spacer content on
mechanical properties. Increasing the number of sp
3
bonds in the backbone of the polymer
Polymer
Modulus
(GPa)
a

Toughness
(MPa)
b

UTS
(MPa)
c

Fracture Strength
(MPa)
d

10% T-6-T 0.23 ± 0.03 1.4 ± 0.3 7.8 ± 0.7 7 ± 1
10% T-8-T 0.14 ± 0.03 1.5 ± 0.5 6.9 ± 0.9 6.6 ± 0.5
10% T-10-T 0.13 ± 0.01 1.0 ± 0.5 5 ± 1 5 ± 1
20% T-8-T 0.42 ± 0.03 2.6 ± 0.4 13 ± 1 13 ± 1
20% T-10-T 0.32 ± 0.03 1.8 ± 0.2 8.6 ± 0.4 7 ± 2
30% T-10-T 0.33 ± 0.04 2 ± 1 11 ± 1 10 ± 1
40% T-10-T 0.75 ± 0.08 1.9 ± 0.4 19.3 ± 0.4 16 ± 3
63

increases the ability of the bonds along the polymer chains to rotate; we therefore predicted
that incorporation of conjugation-break spacers within the backbone would increase the
ability of a solid material to be deformed. Indeed, we observed a decrease in elastic
modulus and UTS as spacer length increased (Figure 2.3b and 2.3c), and a relatively small
effect on toughness (2.3b). These measurements allowed us to distinguish the mechanical
properties between 10% T-6-T, 10% T-8-T, and 10% T-10-T, which had nearly identical
behavior when measured using film-on-elastomer techniques. Figure 2.3d shows the effect
of the fraction of the spacer and the DPP unit on the mechanical properties. The most
striking feature of the stress-strain behavior is the increase in stress and decrease in
extensibility with increasing fraction of the spacer and the DPP unit, again perhaps
attributable to the contradictory effects of simultaneously increasing CBS and DPP content.
The modulus also increased, as shown in Figure 2.3e. Toughness, as plotted in Figure
2.3e, is a function of both the stress and the extensibility, so though it appears that there
was little change in toughness among the polymers with varied T-10-T content, their
underlying stress and extensibility properties were quite different. It is worth noting that
the fraction of the spacer and DPP segments seemed to have more significant effects on
the mechanical properties of the materials than the spacer length.
64

Figure 2.3. Stress-strain curves and corresponding mechanical properties as functions of
the length (with 10% incorporation of the spacer) and fraction of the aliphatic spacer (with
10 carbon atoms in the spacer). (a,d) Representative stress-strain curves of pristine films
obtained using the film-on-water technique. Correlation of toughness and elastic modulus
with (b) the length, and (e) the fraction of the spacer. Values of toughness are obtained by
integrating the total area under the stress-strain curves. Values of elastic modulus are
calculated as the slope of the linear region of the graph. Relationship between ultimate
tensile strength (UTS) and (c) the length or (f) the fraction of the spacer. UTS values were
obtained from the stress–strain curves from the stress at fracture. Mean values and error
bars (standard deviations based on 95% confidence bounds) are based on data collected
from at least three separate measurements. Dashed lines are to guide the eyes.

In short, modulating the length and fraction of the spacer in these low-bandgap
conjugated polymers had a strong effect on the deformability. Introducing more carbon
atoms into the spacer generally had the effect of decreasing the elastic modulus, toughness,
and UTS. Increasing the fraction of the conjugation-break spacer and the DPP unit
increased the values of these properties, making the material stronger, but also less
extensible. Overall, the effects from modulating the fraction of the spacer and DPP
65

monomers were more significant than the modulation of the length of the spacer. Although
the samples tested varied in molecular weight, trends were still observed despite the fact
that mechanical properties have been shown to depend strongly on molecular weight in
homopolymers.
49

2.3 Conclusions
In summary, a family of semi-random polymers with conjugation-break spacers was
synthesized exhibiting notable mechanical properties, attributed both to the break in
conjugation as well as the semi-random structure of the polymer backbone. Film-on-
elastomer measurements revealed low elastic moduli between 0.14-1.3 GPa and several
polymers could be stretched beyond 80% strain before film failure. Further testing of these
polymers as free-standing films using film-on-water methods confirmed low elastic moduli
between 0.13-0.75 GPa. While the electronic properties are generally diminished relative
to similar semiconducting polymers, the trends observed in optical bandgap and hole
mobility suggest that these properties could be tuned and optimized in the future. In fact,
the 10% T-10-T polymer seems to be a promising starting point for future optimization,
with its high mobility of 2.53 E-4 cm
2
V
-1
s
-1
, which is comparable to the fully conjugated
P3HTT-DPP-10% mobility of 9.29 E-4 cm
2
V
-1
s
-1
. This ductile polymer had the lowest
elastic modulus as measured by both film-on-water and film-on-elastomer techniques, and
film-on-elastomer measurements revealed a crack-onset strain of greater than 80%. These
results indicate that semi-random polymers with conjugation-break spacers are promising
candidates for further study in flexible electronics. In Chapter 4, we elucidate the effects
of increasing the CBS without increasing the DPP content, as there were several
measurements wherein the two monomers appeared to have contradictory effects. We also
66

clarify the effects of limited solubility on the measurements by lengthening the DPP side
chain in Chapter 3.

67

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41. Kline, R. J.; McGehee, M. D.; Kadnikova, E. N.; Liu, J.; Fréchet, J. M. J.; Toney, M.
F. Dependence of Regioregular Poly(3-Hexylthiophene) Film Morphology and
Field-Effect Mobility on Molecular Weight. Macromolecules 2005, 38, 3312-3319.
42. Verploegen, E.; Mondal, R.; Bettinger, C. J.; Sok, S.; Toney, M. F.; Bao, Z. Effects of
Thermal Annealing upon the Morphology of Polymer–Fullerene Blends. Adv. Funct.
Mater. 2010, 20, 3519-3529.
43. Zhao, Y.; Zhao, X.; Roders, M.; Qu, G.; Diao, Y.; Ayzner, A. L.; Mei, J.
Complementary Semiconducting Polymer Blends for Efficient Charge Transport.
Chem. Mater. 2015, 27, 7164-7170.
71

44. Balar, N.; O’Connor, B. T. Correlating Crack Onset Strain and Cohesive Fracture
Energy in Polymer Semiconductor Films. Macromolecules 2017, 50, 8611-8618.
45. Howard, J. B.; Ekiz, S.; Cuellar de Lucio, A J; Thompson, B. C. Investigation of
Random Copolymer Analogues of a Semi-Random Conjugated Polymer
Incorporating Thieno[3,4-B]Pyrazine. Macromolecules 2016, 49, 6360-6367.
46. Kim, J.; Nizami, A.; Hwangbo, Y.; Jang, B.; Lee, H.; Woo, C.; Hyun, S.; Kim, T.
Tensile Testing of Ultra-Thin Films on Water Surface. Nat. Commun. 2013, 4, 2520.
47. Rodriquez, D.; Kim, J.; Root, S. E.; Fei, Z.; Boufflet, P.; Heeney, M.; Kim, T.;
Lipomi, D. J. Comparison of Methods for Determining the Mechanical Properties of
Semiconducting Polymer Films for Stretchable Electronics. ACS Appl. Mater.
Interfaces 2017, 9, 8855-8862.
48. Lipomi, D. J. Stretchable Figures of Merit in Deformable Electronics. Adv. Mater.
2016, 28, 4180-4183.
49. Bruner, C.; Dauskardt, R. Role of Molecular Weight on the Mechanical Device
Properties of Organic Polymer Solar Cells. Macromolecules 2014, 47, 1117-1121.

72

Chapter 3: Influence of Acceptor Side-Chain Length on Mechanical and Electronic
Properties of Semi-Random Polymers with Conjugation-Break Spacers

3.1 Introduction
Recently, conjugation-break spacers (CBSs) have been incorporated into conjugated
polymers to impart such advantages as solubility,
1,2
melt processability,
3,4
and molecular
weight-independent optoelectronic properties.
5-7
Several studies, including our own, have
shown that polymers with CBSs have vastly improved mechanical properties over their
fully conjugated counterparts, and yet they are able to retain most of the electronic
characteristics that make conjugated polymers attractive as semiconductors.
8,9

Previous studies from our group have shown that semi-random polymers have more
disorder than perfectly alternating or semi-alternating (random)
10
polymers along the
backbone, which is manifested in bulk properties such as crystallinity.
11
This increased
disorder is expected to enhance the mechanical properties of conjugated polymers that are
known for being stiff and brittle.
12-14
Our previous work (Chapter 2) showed that
combining conjugation-break spacers and a semi-random polymer architecture resulted in
very good mechanical properties (low elastic moduli, high crack-onset strain (COS)) for
conjugated polymers. Another advantageous aspect of semi-random polymers is that they
can be used to tune various other properties, including absorption, electronic energy levels,
and surface energy.
11,15,16

Despite encouraging mechanical and electronic results, our previous semi-random CBS
polymers were hampered by poor solubility, which affected molecular weight, dispersity,
and film-formation ability. This in turn affected our ability to complete mobility and
73

mechanical measurements. Because the solubility of our polymers seemed to decrease with
increasing fraction of CBS monomer in the backbone, contrary to the findings of Zhao et.
al,
3
we presumed that the decrease in solubility was attributable to the corresponding
increase in diketopyrrolopyrrole (DPP) monomer incorporation in the backbone. In an
attempt to alleviate these problems, we undertook the synthesis of a portion of the first set
of polymers with decyltetradecyl (dtd) side-chains on the DPP monomer in place of the
original ethylhexyl (eh) side-chains (Scheme 3.1).

Scheme 3.1. Semi-random polymers with conjugation-break spacers synthesized for this
study
We hypothesized that the longer side-chains on DPP would allow for higher molecular
weight polymers to be synthesized (and isolated) due to their increased solubility and
processability. We anticipated that the polymers with higher fractions of CBS and DPP
would be better able to form films for mechanical and mobility studies than in our previous
study with ethylhexyl side-chains on the DPP. However, we also expected that longer side-
chains would have detrimental effects on optoelectronic properties by increasing the weight
fraction of insulating hydrocarbons within the polymer and increasing the space between
polymer backbones when packed in the solid state.
74

3.2 Results and discussion
A series of eight polymers with varying contents of CBS (and consequently dtdDPP
and 3-hexylthiophene (3HT) monomers) and the fully conjugated reference polymer,
P3HTT-dtdDPP-10%, were synthesized via Stille polycondensation using methods
previously developed in our group.
15
As with the previous study in Chapter 2, only even-
numbered hydrocarbon chains were used as CBSs both for economic practicality and to
eliminate the possibility of an odd-even effect.
17
Polymers were named by the CBS
monomer and the amount by which it was incorporated, as well as the dtdDPP content,
e.g., 10% T-4-T/10% dtdDPP indicates a CBS length of 4 carbons and that monomer makes
up 10 mole percent of the polymer backbone, along with 10 mole percent of dtdDPP and
80 mole percent 3HT. A smaller selection of polymers was targeted for this follow-up study
than for the previous study: all of the T-8-T family so that trends with increasing DPP and
CBS (and decreasing 3HT) content could be observed; bookends (10% and 40%) of the T-
4-T family because that family exhibited the worst solubility; and the lower CBS fractions
of the T-10-T family because of their favorable electronic performance in our previous
study. Polymer composition was confirmed by NMR to match the monomer feed ratios
(Figure B3, Figures B4-B6). Higher molecular weights were attained throughout this
family (Mn = 20.9 – 47.6 kDa, Table 3.1, Table B1, Figure B1) when compared to the
prior family (Mn = 6.4 – 29.0 kDa, Table 2.1), and the fully conjugated polymer had the
lowest molecular weight in this family (Mn = 14.9 kDa). The lower molecular weight of
the P3HTT-dtdDPP-10% polymer is attributed to poor solubility and indicates that the CBS
monomers are indeed enhancing the solubility of this family. Interestingly, as increasing
amounts of the alkyl spacers and dtdDPP were incorporated into the polymer backbone,
75

the polymerization yields increased (Table B2), a trend also linked to solubility. For those
polymers with low CBS and dtdDPP content, large amounts of the polymers (with
molecular weights as high as 7 kDa, much higher than the 2-3 kDa typically observed in
conjugated polymers) were extracted into the hexane Soxhlet fractions giving hexane
yields of 30-50%. These high hexane yields reduced the yield in the subsequent chloroform
Soxhlet fraction. Lower hexane yields and higher chloroform yields were generally
attained at higher CBS and dtdDPP incorporation (Figure B2), indicating that those
polymers were less soluble in hexane, though they still dissolved readily in halogenated
organic solvents.

Table 3.1. SEC, thermal, and electronic data for semi-random polymer family
Polymer
Mn
a
(kDa) Ð
a

Tm / Tc
b
(°C)
Eg
c
(eV)
HOMO
d
(eV)
µh
e
(cm
2
V
-1
s
-1
)
P3HTT-dtdDPP-10% 14.9 3.24 150 / 143 1.51 5.62 1.45 E-5
10% T-4-T/10% dtdDPP 29.9 2.20 108 / 93 1.54 5.55 1.62 E-6
10% T-8-T/10% dtdDPP 26.1 3.90 - / -
f
1.56 5.59 1.37 E-6
10% T-10-T/10% dtdDPP 33.0 3.27 - / -
f
1.56 5.57 8.93 E-7
20% T-8-T/20% dtdDPP 35.5 2.59 72 / 50 1.62 5.56 4.37 E-7
20% T-10-T/20% dtdDPP 31.1 5.05 56 / 33 1.63 5.59 3.25 E-7
30% T-8-T/30% dtdDPP 47.6 2.37 103 / 64 1.67 5.52 2.06 E-7
40% T-4-T/40% dtdDPP 20.9 2.86 164 / 140 1.68 5.57 4.11 E-7
40% T-8-T/40% dtdDPP 44.1 2.54 120 / 86 1.71 5.54 3.25 E-7
a) Obtained via size-exclusion chromatography (SEC) versus polystyrene standards; b)
Obtained via differential scanning calorimetry (DSC); c) Calculated from the absorption
band edge in annealed thin films, where Eg = 1240/λedge; d) Estimated from cyclic
voltammetry (CV) oxidation onset versus ferrocene; e) Calculated from SCLC mobility
measurements in hole-only devices with the architecture: ITO/PEDOT:PSS/polymer/Al,
where the polymer layer was spin-cast from chloroform and stored in a N2 cabinet for 30
min prior to aluminum deposition; f) No apparent thermal transitions.

76

3.2.1 Optical properties
The as-cast and annealed UV-vis spectra of the T-8-T subfamily are shown in Figure
3.1 as a representative example of optical trends. Spectra for all of the polymers, as-cast
and annealed, are provided in Figures B7-B14. Similar to the previous study, all polymers
with a break in conjugation had a blue-shifted absorption onset and wider optical bandgap
when compared to the fully conjugated P3HTT-dtdDPP-10% optical bandgap of 1.51 eV
(Table 3.1). As the percentage of CBS and dtdDPP incorporated into the polymer
increased, the absorption onset continued to blue-shift to a maximum bandgap of 1.71 eV
for the 40% T-X-T subfamily (Table B6, Figure B16). As expected, the optical bandgaps
are not dependent on the length of the DPP side-chain and the values presented herein are
virtually the same as those reported for the polymers with ethylhexyl side-chains on the
DPP monomer.

Figure 3.1. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast from
chloroform and (a) placed in a N2 cabinet for 30 minutes or (b) annealed under N2 at 150
°C for 30 minutes. The P3HTT-dtdDPP-10% absorption spectrum (dashed black line) is
provided for a fully conjugated reference. (■) 10% T-8-T; (●) 20% T-8-T; (▲) 30% T-8-
T; (▼) 40% T-8-T.
Similar to what we reported previously, the length of the CBS had no effect on the
onset. However, thermal annealing slightly increased the bandgap, leading to a blue-shift
77

in absorption (Table B5, Figure B15). Thermal annealing also changed the shape of the
absorption spectra, leading to a relative increase in intensity in the P3HT region of the
spectra (400-550 nm) and corresponding decrease in intensity in the DPP region of the
spectra (550-750 nm) (Figure 3.1). All absorption coefficients for the new family of
polymers were lower than those obtained for the previous family of polymers (Tables B3-
B4). This is as expected, for the longer side-chains on DPP make it so that there is less
electroactive polymer per unit volume than for the polymers with shorter side-chains. This
in turn makes the decyltetradecyl polymers less optically dense than the ethylhexyl
polymers.
3.2.2 Structural Properties
The high melting and crystallization points observed for the fully conjugated P3HTT-
DPP-10% polymer (Tm = 208 °C; Tc = 182 °C) were significantly depressed in the fully
conjugated P3HTT-dtdDPP-10% (Tm = 150 °C; Tc = 143 °C) when analyzed by differential
scanning calorimetry (DSC) (Table 3.1, Figure B22). Similar to our previous study, no
thermal transitions were observed for most of the 10% T-X-T polymers. Many of the
dtdDPP polymers that exhibited thermal transitions in this study did not display such
transitions in the polymers with shorter side-chains from Chapter 2. This makes it difficult
to draw comparisons between the two families. Within the T-8-T/dtdDPP family, however,
a clear trend appeared in which the melting and crystallization temperatures rose with
increasing CBS and DPP incorporation (Tables B8-B9, Figures B31-B32). This mirrors
the trend that was seen in the family of polymers with shorter side-chains. Because the
equilibrium melting point of a polymer is inversely related to the entropy of fusion for
melting,
18
the increasing melt and crystallization temperatures of this polymer family
78

indicated that the entropy of these transitions was decreasing as DPP and CBS content
increased. This decrease in entropy was perhaps due to a decrease in chain flexibility,
17

which is most reasonably attributed to the rise in DPP monomer content, as the
conjugation-break spacer was expected to enhance polymer chain flexibility. This is
contrary to the trend that emerged with increasing spacer length, in which the entropy of
the thermal transitions increased, thus depressing melting and crystallization temperatures,
as we observed in our previous study (Chapter 2).
Crystallinity in thin films was investigated via grazing incidence X-ray diffraction
(GIXRD). While weak diffraction peaks were observed for the majority of the polymers in
as-cast films (Figure B33, Table B10), after annealing the weak peaks disappeared and
intense diffraction peaks were observed only for those polymers incorporating 40% CBS
and dtdDPP (Figure B34, Table B11). The 100 diffraction peaks for as-cast films
corresponded to a lamellar packing distance of 17.3-20.4 Å, larger than the range of 15.6-
16.5 Å observed in our original study. This can be explained by the longer DPP side-chains
generally causing the polymer chains to stack further apart, a phenomenon previously
observed in our group when exchanging ethylhexyl side-chains for decyltetradecyl side-
chains.
19
Within the T-8-T subfamily, a trend appeared in the as-cast films in which the
lamellar packing distance decreased with increasing CBS and dtdDPP incorporation.
Although the fraction of monomer with long side-chains was increasing within this data
set, it seems that increasing the CBS content allowed the polymer chains to pack closer
together, counter to the trend observed by Ekiz et al when increasing DPP content without
the use of a CBS monomer.
19
Interestingly, there is no observed diffraction peak for the
fully conjugated polymer before annealing. These results pointed to the conjugation-break
79

spacer as a key factor in polymer crystallinity, perhaps imbuing the conjugated segments
of the backbone with enough flexibility to arrange themselves in a manner favorable for π-
π stacking. It is curious that this order seemed to disappear in most samples upon annealing,
although the intensity of the 40% CBS/40% dtdDPP samples increased greatly. We could
even observe a 200 reflection in the 40% T-4-T/40% dtdDPP annealed film. Upon
annealing, the 40% T-4-T/40% dtdDPP lamellar packing distance diminished slightly from
20.4 Å to 20.1 Å, whereas the 40% T-8-T/40% dtdDPP lamellar packing distance
decreased from 17.3 Å to 15.9 Å. Presumably the greater flexibility imbued in the polymer
by the longer CBS allowed the polymer backbones to draw closer to each other during
annealing than was possible for the polymer with a shorter CBS.
3.2.3 Electronic Properties
The highest occupied molecular orbital (HOMO) energy was calculated from oxidation
onset using cyclic voltammetry (CV). As can be seen in Table 3.1 (and Figure B21), there
were no observable trends in the HOMOs of these polymers. All values were
approximately the same, producing an average HOMO of 5.56 ± 0.02 eV (well within the
instrument error of ± 0.05 eV). This was slightly more shallow than the 5.62 eV HOMO of
the fully conjugated P3HTT-dtdDPP-10%. This trend exactly mirrors that observed in the
first study, with the exception that all HOMO energies are approximately 0.1 eV deeper
for the polymers with longer DPP side-chains. Although several studies have investigated
the effects of side-chain length on electronic properties of conjugated polymers, most
conclude that the length of the side-chain has no effect on frontier orbital levels. However,
in examining the data from these studies, we find that there is a slight tendency toward
80

polymers with longer side-chains having deeper HOMOs (usually ≤ 0.05 eV),
20,21
a small
but consistent trend that matches our observations from this study.
Space-charge limited current (SCLC) mobility measurements were performed on hole-
only devices for this set of polymers. Films were cast from chloroform solutions, in which
all the polymers readily dissolved, and devices were tested with and without thermal
annealing. Surprisingly, all of the polymers with CBSs exhibited decreased performance
after annealing, contrary to what had been observed in the first family of polymers (Table
3.1, Tables B12-B13). Only the fully conjugated P3HTT-dtdDPP-10% had an improved
performance after annealing, doubling from 1.45 E-5 cm
2
V
-1
s
-1
to 2.97 E-5 cm
2
V
-1
s
-1
.
Similar to our results from Chapter 2 and a study by the Mei group,
3
charge mobilities
were seen to steadily decrease with increasing CBS (and DPP) content, attributed to the
corresponding increase in insulating alkyl chains (Figure B35-B36). However, contrary to
our first study, this family of polymers shows a decrease in charge mobility with increasing
CBS length. This trend is intuitive and matches the trend observed by Zhao et al
17
that an
increase in insulating hydrocarbon per mole of polymer leads to a decrease in hole mobility.
Unlike the phenomenon observed for the 10% T-10-T polymer in our first study, the
apparent increase in crystalline order observed in the 40% T-4-T/40% dtdDPP and 40% T-
8-T/40% dtdDPP GIXRD patterns after annealing did not seem to aid charge transport in
these polymers, as both polymers exhibited about a three-fold decrease in mobility after
thermal annealing. In this polymer family, the fully conjugated P3HTT-dtdDPP-10%
outperformed the CBS polymers by an order of magnitude or more, contrary to our
previous study which found at least one CBS polymer to perform comparable to the fully
conjugated polymer.
81

3.2.4 Mechanical properties
While the SCLC charge mobilities for this family of polymers were diminished
compared to the first family studied, the mechanical properties were enhanced. Because
the entire family was easily soluble in chloroform and formed uniform films, the intrinsic
mechanical properties of the freestanding films could be measured using the film-on-water
technique (Table 3.2). The fully conjugated P3HTT-dtdDPP-10% had an elastic modulus
(0.167 GPa) approximately half that of the fully conjugated analog from the previous study,
P3HTT-DPP-10% (0.32 GPa), a sign that the longer side-chains imbued the polymer with
more ductility, a phenomenon that has previously been noted in the literature.
13,20,22
As
another indication of ductility, the strain necessary to fracture the film was greatly
enhanced in the polymer with longer side-chains on the DPP (63% vs. 10%).

82

Table 3.2. Tabulated values of mechanical properties measured using the film-on-water
technique
Polymer
Modulus

(GPa)
a

Toughness

(MPa)
b

UTS

(MPa)
c

Fracture
Strength

(MPa)
d

Fracture
Strain
(%)
e

P3HTT-dtdDPP-10% 0.167 ± 0.016 4.95 ± 1.00 8.68 ± 0.02 8.20 ± 0.55 63 ± 11
10% T-4-T/10% dtdDPP -
f
-
f
-
f
-
f
-
f

10% T-8-T/10% dtdDPP 0.023 ± 0.003 4.99 ± 0.26 3.61 ± 0.27 3.57 ± 0.32 185 ± 33
10% T-10-T/10% dtdDPP 0.016 ± 0.002 4.13 ± 1.06 2.15 ± 0.50 1.90 ± 0.80 217 ± 6
20% T-8-T/20% dtdDPP 0.026 ± 0.002 3.98 ± 1.20 2.63 ± 0.52 2.51 ± 0.55 190 ± 35
20% T-10-T/20% dtdDPP 0.006 ± 0.001 4.15 ± 0.20 1.75 ± 0.33 1.74 ± 0.32 373 ± 67
30% T-8-T/30% dtdDPP 0.023 ± 0.004 9.05 ± 2.04 4.49 ± 0.55 4.38 ± 0.67 323 ± 46
40% T-4-T/40% dtdDPP 0.038 ± 0.001 0.79 ± 0.06 1.77 ± 0.00 1.58 ± 0.05 53 ± 4
40% T-8-T/40% dtdDPP 0.036 ± 0.004 20.59 ± 1.33 8.45 ± 0.11 8.19 ± 0.04 398 ± 32
a) Elastic moduli were derived from stress–strain curves by taking the slope of the linear
regime of the graph; b) Toughness values were obtained by integrating the entirety of the
stress–strain curve; c) UTS was obtained from the maximum stress in the stress–strain
curve; d) Fracture strength was obtained from a stress–strain curve by reading the stress at
failure; e) Fracture strain was obtained from a stress-strain curve by reading the strain at
failure; f) Due to poor film integrity, no stress-strain curves could be obtained.

On the whole, this family of polymers had lower elastic moduli than the polymers
analyzed in the previous study by at least an order of magnitude, with the 20% T-10-T/20%
dtdDPP polymer yielding a modulus as low as 6 ± 1 MPa! This is nearing the modulus
range expected from conventional elastomers such as PDMS (E = 0.6-2.5 MPa)
23
or
polyisoprene (E = 0.36 MPa).
24
Unfortunately, this polymer had one of the lowest charge
mobilities when measured by SCLC, typifying the tradeoff between mechanical and
electronic properties so often witnessed in semiconducting polymers.
25
Surprisingly, the
dtdDPP family of polymers was also slightly tougher than the original family. Since
toughness is dependent on the area under the stress-strain curve, an increase in strength or
extensibility of the material will improve the toughness of the material. There is also a
83

general trend suggesting that greater intermolecular forces between polymer chains will
improve the toughness.
25
In this case, we attribute the increased toughness of the longer
side-chain polymers to the increase in extensibility these polymers exhibit, with fracture
strains as high as 398%. These polymers had slightly lower or comparable UTS and
fracture strength values than the previous family.
Within the family of CBS polymers with decyltetradecyl side-chains on the DPP
monomer, several trends emerged. All of the CBS polymers had a lower elastic modulus
than the fully conjugated P3HTT-dtdDPP-10% by an order of magnitude. The modulus
appeared to decrease with longer conjugation-break spacers, a trend mirrored by the
previous family with ethylhexyl side-chains. Within polymers of the same CBS length,
moduli seemed to be consistent across spacer content until 40% CBS incorporation, at
which point it increased slightly, perhaps due to an increase in the stiff DPP monomer
content. The 30% and 40% T-8-T polymers had much higher toughness than the other
polymers, including even the fully conjugated polymer, again perhaps attributable to the
increase in DPP monomer content. The UTS and fracture strength were observed to
generally decrease with longer spacers, indicating that though these materials were more
extensible than polymers with shorter spacers, they were also slightly weaker. Within the
T-8-T subfamily, UTS and fracture strength increased with 30% and 40% incorporation of
the CBS and DPP monomer, until the 40% T-8-T/40% dtdDPP polymer had comparable
UTS and fracture strength to the fully conjugated P3HTT-dtdDPP-10% polymer. With the
exception of the 40% T-4-T/40% dtdDPP polymer, the fracture strain increased with longer
conjugation-break spacers and higher percent incorporation of CBS, resulting in films able
84

to withstand strains much higher than fully conjugated P3HTT-dtdDPP-10% polymer and
higher even than the CBS polymers tested within the first study.
3.3 Conclusions
In this study a set of semi-random polymers with varying lengths and compositions of
conjugation-break spacers and decyltetradecyl-substituted DPP monomers were
synthesized and analyzed for optoelectronic and mechanical properties. This family was a
derivative of the polymer family synthesized in Chapter 2 with ethylhexyl-substituted
DPP and was designed with the aim of improving solubility, and thus film-forming ability.
These polymers had much higher solubility than the family from Chapter 2, and attained
higher molecular weights, formed films with high integrity, and displayed extraordinary
mechanical properties, with elastic moduli as low as 6 MPa and fracture strains as high as
398%. Unfortunately, these remarkable mechanical properties were paired with a
diminishment of electronic properties, with much lower hole mobilities than the previous
family. The highest hole mobility measured by SCLC for CBS polymers in this family was
1.62 E-6 cm
2
V
-1
s
-1
, as compared to 2.53 E-4 cm
2
V
-1
s
-1
for the polymer family from
Chapter 2 with shorter side-chains. This study confirmed the oft-reported mechanical-
electronic properties trade-off, but indicates that perhaps a favorable balance can be struck
with side-chains of a middling length, between ethylhexyl and decyltetradecyl.

85

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14. Roth, B.; Savagatrup, S.; V. de los Santos, Nathaniel; Hagemann, O.; Carlé, J. E.;
Helgesen, M.; Livi, F.; Bundgaard, E.; Søndergaard, R. R.; Krebs, F. C.; Lipomi, D.
J. Mechanical Properties of a Library of Low-Band-Gap Polymers. Chem. Mater.
2016, 28, 2363-2373.
15. Burkhart, B.; Khlyabich, P. P.; Cakir Canak, T.; LaJoie, T. W.; Thompson, B. C.
“Semi-Random” Multichromophoric rr-P3HT Analogues for Solar Photon
Harvesting. Macromolecules 2011, 44, 1242-1246.
16. Khlyabich, P. P.; Burkhart, B.; Ng, C. F.; Thompson, B. C. Efficient Solar Cells from
Semi-Random P3HT Analogues Incorporating Diketopyrrolopyrrole.
Macromolecules 2011, 44, 5079-5084.
17. Zhao, X.; Zhao, Y.; Ge, Q.; Butrouna, K.; Diao, Y.; Graham, K. R.; Mei, J.
Complementary Semiconducting Polymer Blends: The Influence of Conjugation-
Break Spacer Length in Matrix Polymers. Macromolecules 2016, 49, 2601-2608.
18. Afifi-Effat, A. M.; Hay, J. N. Enthalpy and Entropy of Fusion and the Equilibrium
Melting Point of Polyethylene Oxide. J. Chem. Soc., Faraday Trans. 2 1972, 68,
656-661.
19. Ekiz, S.; Gobalasingham, N. S.; Thompson, B. C. Exploring the Influence of
Acceptor Content on Semi‐random Conjugated Polymers. J. Polym. Sci., Part A:
Polym. Chem. 2017, 55, 3884-3892.
20. Heintges, G. H. L.; Leenaers, P. J.; Janssen, R. A. J. The Effect of Side-Chain
Substitution and Hot Processing on Diketopyrrolopyrrole-Based Polymers for
Organic Solar Cells. J. Mater. Chem. A 2017, 5, 13748-13756.
21. Duan, C.; Willems, R. E. M.; van Franeker, J. J.; Bruijnaers, B. J.; Wienk, M. M.;
Janssen, R. A. Effect of Side Chain Length on the Charge Transport, Morphology,
and Photovoltaic Performance of Conjugated Polymers in Bulk Heterojunction Solar
Cells. J. Mater. Chem. A 2016, 4, 1855-1866.
87

22. Rodriquez, D.; Kim, J.; Root, S. E.; Fei, Z.; Boufflet, P.; Heeney, M.; Kim, T.;
Lipomi, D. J. Comparison of Methods for Determining the Mechanical Properties of
Semiconducting Polymer Films for Stretchable Electronics. ACS Appl. Mater.
Interfaces 2017, 9, 8855-8862.
23. Khanafer, K.; Duprey, A.; Schlicht, M.; Berguer, R. Effects of Strain Rate, Mixing
Ratio, and Stress–strain Definition on the Mechanical Behavior of the
Polydimethylsiloxane (PDMS) Material as Related to its Biological Applications.
Biomed. Microdevices 2009, 11, 503-508.
24. Baboo, M.; Dixit, M.; Sharma, K.; Saxena, N. Mechanical and Thermal
Characterization of Cis-Polyisoprene and Trans-Polyisoprene Blends. Polym. Bull.
2011, 66, 661-672.
25. Savagatrup, S.; Printz, A. D.; Rodriquez, D.; Lipomi, D. J. Mechanical Properties of
Organic Semiconductors for Stretchable, Highly Flexible, and Mechanically Robust
Electronics. Chem. Rev. 2017, 117, 6467-6499.

88

Chapter 4: Effects of Modulating Content of Conjugation-Break Spacers in Semi-
Random Polymers on Mechanical and Electronic Properties

4.1 Introduction
In several recent studies, semiconducting polymers incorporating a small amount of
conjugation-break spacer (CBS) monomer have been shown to greatly improve the
mechanical properties over the parent fully conjugated polymer while still retaining the
parent polymer’s electronic properties.
1,2
This was an effect we also observed in our first
study (Chapter 2), in which polymers with a small amount of CBS monomer incorporated
into the backbone generally exhibited increased ductility and decreased elastic moduli
when compared to a fully conjugated semi-random polymer.
Contrary to the AA/BB functionalized monomers used to make perfectly alternating or
semi-alternating
3
(random) conjugated polymers, semi-random polymers use AA, AB, and
BB functionalized monomers. The less restrictive linkage pattern available to monomers
in the semi-random polymer increases disorder along the polymer backbone, leading to
decreased stiffness and brittleness in polymer films.
4-6
The semi-random architecture has
also been shown to broaden the absorption of conjugated polymers over that of perfectly
alternating polymers by creating a broader range of chromophores within the polymer
backbone than is accessible with a perfectly alternating architecture.
7
Additionally,
incorporating 3-hexylthiophene allows for the retention of the favorable properties of
P3HT while tuning such properties as absorption, electronic energy levels, and surface
energy.
8,9

89

In Chapter 2, the effects of conjugation-break spacer content in semi-random polymers
were examined, but due to the AA/AB/BB functionality of the three monomers used, the
diketopyrrolopyrrole (DPP) acceptor content was altered along with CBS content to
maintain stoichiometric balance. However, it appeared that increasing the DPP and CBS
monomers simultaneously had competing effects on solubility, bandgap, and mechanical
properties. In this study, we sought to disentangle the effects of increasing CBS content
from the effects of increasing DPP while maintaining the semi-random polymer
architecture.
To this end, a second conjugation-break spacer monomer with dibromo functionality
was added so that DPP content could remain constant while increasing CBS content
(Scheme 4.1). The DPP monomer content was fixed at 10 mole percent. This was done to
minimize negative effects from DPP on solubility and mechanical properties, and because
the 10% DPP polymers from our previous study had the highest charge mobility. The T-8-
T monomer was chosen as the fixed spacer length because it was presumed to have minimal
impact on solubility. We anticipated that this polymer family would have improved
solubility and mechanical properties when compared to the T-8-T family of polymers that
was synthesized with variable DPP content. However, we expected that electronic
properties between the two families would be comparable, because the conjugation-break
spacer content was varied by the same amount, so the effective conjugation length of the
polymer should not be different.
90

Scheme 4.1. Semi-random polymers with conjugation-break spacers synthesized for this
study

4.2 Results and discussion
A family of five polymers with 10-50% CBS (and consequently 10% DPP and 80-40%
3-hexylthiophene (3HT) incorporation) and the fully conjugated reference polymer,
P3HTT-DPP-10%, were synthesized via Stille polycondensation using methods previously
developed in our group.
7
T-8-T was chosen as the CBS monomer because it was presumed
to have minimal impact on polymer solubility. DPP was kept at 10% to minimize its
negative effects on solubility and because the 10% DPP polymers from previous studies
had the highest charge mobility. The side-chains on the DPP monomer were kept short
(ethylhexyl instead of decyltetradecyl) because of the detrimental effects that long side-
chains were shown to have on charge mobility. Polymers were named by the CBS
monomer and the amount by which it was incorporated, as well as the DPP monomer and
its content, e.g., 30% T-8-T/10% DPP indicates a CBS length of 8 carbons and that
monomer makes up 30 mole percent of the polymer backbone, along with 10 mole percent
of DPP and 60 mole percent 3HT. Polymer composition was confirmed by NMR to match
91

the monomer feed ratios (Figure C4-C5). The average molecular weight for this family
was low at 12.7 kDa, attributed to poor solubility (Table 4.1, Table C1, Figure C1). The
low molecular weight and broad dispersity of the 20% T-8-T/10% DPP polymer is due to
a bimodal weight distribution – the hexane Soxhlet performed to remove oligomers did not
succeed in extracting the low molecular weight fraction for this polymer. Yields were
generally greater than 60% (Table C2, Figure C2) with the exception of 50% T-8-T/10%
DPP in which a great deal of solid remained in the Soxhlet thimble after extraction with
chloroform. These solubility obstacles were not anticipated, because it was presumed that
solubility would be enhanced by both keeping DPP content low and increasing CBS
content. However, it is apparent that solubility was not enhanced, the reason for which is
not known at this time.

Table 4.1. SEC, thermal, and electronic data for semi-random polymer family
Polymer
Mn
a
(kDa) Ð
a

Tm / Tc
b
(°C)
Eg
c
(eV)
HOMO
d
(eV)
µh
e
(cm
2
V
-1
s
-1
)
P3HTT-DPP-10% 10.5 4.31 211 / 207 1.49 5.39 1.02 E-4
10% T-8-T/10% DPP 19.7 6.23 - / -
f
1.56 5.52 3.02 E-6
20% T-8-T/10% DPP 8.5 8.18 - / -
f
1.59 5.48 2.20 E-6
30% T-8-T/10% DPP 12.8 3.75 - / -
f
1.65 5.50 2.02 E-7
40% T-8-T/10% DPP 12.4 5.24 - / -
f
1.68 5.63 3.63 E-8
50% T-8-T/10% DPP 10.2 4.97 78 / 42 1.70 5.48 6.82 E-9
a) Obtained via size-exclusion chromatography (SEC) versus polystyrene standards; b)
Obtained via differential scanning calorimetry (DSC); c) Calculated from the absorption
band edge in thin films, where Eg = 1240/λedge; d) Estimated from cyclic voltammetry (CV)
oxidation onset versus ferrocene; e) Calculated from SCLC mobility measurements in hole-
only devices with the architecture: ITO/PEDOT:PSS/polymer/Al, where the polymer layer
was spin-cast from hot chloroform and placed in a N2 box for 30 min prior to aluminum
deposition; f) No apparent thermal transitions.

92

4.2.1 Optical Properties
The UV-vis spectra of the polymers are shown in Figure 4.1. All of the CBS polymers
had a blue-shifted absorption onset, leading to a wider optical bandgap, when compared to
the fully conjugated P3HTT-DPP-10% optical bandgap of 1.49 eV (Table 4.1). As the
percentage of CBS incorporated into the polymer increased, the absorption onset continued
to blue-shift to a maximum bandgap of 1.70 eV for the 50% T-8-T/10% DPP polymer
(Figure C9). Both of these trends are as expected, for the extended conjugation of a
semiconducting polymer allows the highest occupied and lowest unoccupied molecular
orbitals to draw nearer to each other, lowering the bandgap of the extended solid. By
interrupting this extended sp
2
hybridization with the sp
3
hybridized carbons of alkyl chains,
the bandgap increased and therefore the absorption onset moved to a higher energy.
Interestingly, the bandgap did not increase as rapidly in this set of polymers as in the
polymers that increased DPP content simultaneously with CBS content (Chapters 2-3).
This is puzzling because a previous study from our group found that the absorption onset
red-shifted with increasing DPP acceptor content.
10
We expected that by keeping the DPP
content constant at 10 mole percent, the bandgap would increase more rapidly with
increasing CBS content than in the study where DPP content rose simultaneously.

93

Figure 4.1. (a) UV-Vis absorption spectra of the T-8-T series thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. (b) UV-Vis absorption
spectra of the T-8-T series thin films spin-cast from o-DCB and annealed under N2 at 150
°C for 30 minutes. The P3HTT-DPP-10% absorption spectrum (dashed black line) is
provided for a fully conjugated reference. (■) 10% T-8-T/10% DPP; (●) 20% T-8-T/10%
DPP; (▲) 30% T-8-T/10% DPP; (▼) 40% T-8-T/10% DPP; (♦) 50% T-8-T/10% DPP.

As expected, absorption in the P3HT region of 400-550 nm decreased as CBS content
increased due to a consequent decrease in loading of the 3-hexylthiophene monomer
(Figure 4.1). This was accompanied by an absorption peak that grew in at about 300-400
nm. This region is attributed to isolated bithiophene segments, which are likely to form in
this semi-random system in which the thiophene-terminated CBS monomer can bond to
itself (Scheme 4.1). Mysteriously, the DPP region from 600-800 nm also appeared to
decrease as CBS content increased, despite the DPP content being held constant. It could
be that the DPP absorption region was enhanced by extended conjugation, which
diminished as spacer content increased, though this would primarily account for the
narrowing of the peak in that region, not the decrease in intensity. Very little change was
observed in the optical characteristics of the films upon thermal annealing.

94

4.2.2 Structural Properties
Conjugated polymers are known for their high melting and crystallization
temperatures, and it has been shown that incorporating conjugation-break spacers into the
backbone of these polymers can dramatically decrease the temperature of these thermal
transitions, even leading to melt processability.
11-13
However, the polymers synthesized in
this study seemed to be primarily amorphous when analyzed by differential scanning
calorimetry (DSC), with thermal transitions only being observed in the 50% T-8-T/10%
DPP polymer. It could be that the addition of a fourth monomer into the polymerization
system led to increased structural disorder, as has been seen in semi-random systems in the
past.
7,14

Crystallinity in annealed and as-cast polymer films was examined by grazing incidence
X-ray diffraction (GIXRD). The diffraction peak intensity decreased dramatically across
the 10%, 20%, and 30% T-8-T/10% DPP polymers in as-cast films (Figure C19, Table
C10), corresponding to the decrease in 3HT content in those polymers. Surprisingly, after
thermal annealing, only the 10% T-8-T/10% DPP polymer exhibited a 100 diffraction peak,
much reduced in intensity compared to its as-cast film (Figure C20, Table C11), though
the fully conjugated P3HTT-DPP-10% polymer 100 peak increased in intensity after
annealing. All of the polymers in this family had a smaller lamellar packing distance than
P3HT (16.7 Å)
15
, as previously observed in semi-random polymers,
8
but also a smaller
lamellar packing distance than those polymers in which DPP content increased
concurrently with CBS content. It could be that the branched side-chains on the DPP
monomer maintained interchain distance in the original family of polymers, while the
95

polymers in this study with a constant low content of DPP could pack closer because of a
relative decrease in the fraction of branched side-chains.
4.2.3 Electronic Properties
The highest occupied molecular orbital (HOMO) energy was calculated from oxidation
onset using cyclic voltammetry (CV). As can be seen in Table 4.1 (and Figure C12), there
were no observable trends in the HOMOs of these polymers. All values were
approximately the same, producing an average HOMO of 5.52 ± 0.06 eV (nearly within
the instrument error range ± 0.05 eV). This was slightly deeper than the 5.39 eV HOMO
of the fully conjugated P3HTT-DPP-10%. It is interesting to consider that the variation in
frontier orbital levels could be attributed to changing effective conjugation lengths, as the
amount of spacer as well as its potential distribution within the polymer was varied.
Charge mobility was measured by space-charge limited current (SCLC) techniques in
hole-only devices to gauge the polymers’ potential utility in organic electronic devices. As
had been observed by Zhao et al,
13
hole mobility decreased logarithmically with increasing
CBS content. This was a more dramatic trend than the decrease in charge mobility observed
in either Chapter 2 or Chapter 3 in which CBS and DPP content increased concurrently,
leading us to infer that despite its negative effects on solubility and processability, the DPP
monomer contributes favorably to the electronic properties of the polymer. Similar to what
we observed with the long side-chain DPP study (Chapter 3), but contrary to our first
study (Chapter 2), the charge mobilities generally decreased upon thermal annealing.
Though a minor effect, this could be associated with the decrease in order that was observed
in the GIXRD measurements upon thermal annealing.

96

4.2.4 Mechanical Properties
In an effort to understand the mechanical properties that would affect the behavior of
these polymers in a flexible electronic device, film-on-water mechanical testing was
performed, with the data summarized in Table 4.2. Due to poor film integrity,
measurements could not be completed on the 50% T-8-T/10% DPP polymer. From the
film-on-water tests we learned that increasing the CBS content while keeping the DPP
content at 10% led to a logarithmic decrease in elastic modulus, E (Figure C23). This is
contrary to the effect observed in the Chapter 2, in which increasing CBS and DPP content
simultaneously led to a slight increase in elastic modulus. The modulus for the 40% T-8-
T/10% DPP polymer is as low as 2 MPa, within the range of conventional elastomers such
as PDMS (E = 0.6-2.5 MPa)!
16
Unfortunately this same polymer had poor electronic
properties, with an average hole mobility of 3.63 E-8 cm
2
V
-1
s
-1
in as-cast films, another
example of the pervasive trade-off between mechanical and electronic properties in
semiconducting polymers.
17

97

Table 4.2. Tabulated values of mechanical properties measured using the film-on-water
technique
Polymer
Modulus

(GPa)
a

Toughness

(MPa)
b
UTS

(MPa)
c
Fracture
Strength

(MPa)
d

Fracture
Strain
(%)
e

10% T-8-T/10% DPP 0.140 ± 0.030 1.5 ± 0.5 6.9 ± 0.9 6.6 ± 0.5 39 ± 9
20% T-8-T/10% DPP 0.069 ± 0.006 23.65 ± 0.89 18.99 ± 1.26 18.73 ± 1.25 200 ± 5
30% T-8-T/10% DPP 0.009 ± 0.004 32.89 ± 4.43 13.58 ± 1.92 12.74 ± 0.62 448 ± 29
40% T-8-T/10% DPP 0.002 ± 0.000 14.99 ± 3.14 13.07 ± 4.54 9.97 ± 4.78 293 ± 25
50% T-8-T/10% DPP -
f
-
f
-
f
-
f
-
f

a) Elastic moduli were derived from stress–strain curves by taking the slope of the linear
regime of the graph; b) Toughness values were obtained by integrating the entirety of the
stress–strain curve; c) UTS was obtained from the maximum stress in the stress–strain
curve; d) Fracture strength was obtained from a stress–strain curve by reading the stress at
failure; e) Fracture strain was obtained from a stress-strain curve by reading the strain at
failure; f) Due to poor film integrity, no stress-strain curves could be obtained.

In general, this family of polymers was tougher than the original family studied in
Chapter 2. Since toughness is dependent on the area under the stress-strain curve, an
increase in strength or extensibility of the material will improve the toughness of the
material. There is also a general trend suggesting that greater intermolecular forces between
polymer chains will improve the toughness.
17
In this case, we attribute the increased
toughness of the polymers to the increase in extensibility these polymers exhibit, with
fracture strains as high as 448%. The toughness had a slight tendency to increase with
increasing CBS monomer incorporation (Figure C24), mirroring the rise in fracture strain
observed for the 10%, 20%, and 30% T-8-T/10% DPP polymers (Figure C27). Fracture
strain in these samples far exceeded those values obtained in Chapter 2, with films able to
extend several times beyond their original length before fracturing.
The values measured for ultimate tensile strain (UTS) and fracture strength for this set
of polymers were comparable to those obtained from film-on-water tests of the original
98

family of polymers. Neither of these values showed a strong trend with changing
conjugation-break spacer content.
When compared to the film-on-elastomer modulus (0.32 GPa) and crack-onset strain
(10%) measured for the fully conjugated P3HTT-DPP-10% polymer in the original study,
all of the polymers examined in this study have superior mechanical properties. Low elastic
moduli and high fracture (or crack-onset) strains are generally considered to be better for
flexible electronics applications, as they indicate a material which is more easily deformed
and which produces decreased interfacial stresses with other layers in a device.
17

4.3 Conclusions
A series of semi-random polymers with 10% DPP and varying amounts of conjugation-
break spacer was synthesized and analyzed for optoelectronic and mechanical properties.
While the polymers produced in this study did not have exemplary electronic properties,
the data presented herein served to disentangle the effects of increasing CBS and DPP
content simultaneously. Polymer solubility, molecular weight, and processability were not
shown to improve dramatically compared to those polymers in the previous study that
varied the DPP content, indicating that the DPP monomer was not the only factor
hampering solubility. SCLC hole mobilities in this family of polymers dropped more
rapidly as CBS content increased than in the family where CBS and DPP content increased
concurrently, signifying that the DPP monomer contributed beneficially to charge mobility.
Finally, the mechanical properties of this series of polymers were quite notable, with lower
elastic moduli, higher toughness, and higher fracture strains than the previous family.
These observations indicated that the DPP monomer, while beneficial to the electronic
properties of the polymers, indeed was responsible for undesirable mechanical properties.
99

These insights can guide us to more judiciously select monomers to build semiconducting
polymers with both favorable mechanical and electronic properties.

100

4.4 References
1. Oh, J. Y.; Rondeau-Gagné, S.; Chiu, Y.; Chortos, A.; Lissel, F.; Wang, G. N.;
Schroeder, B. C.; Kurosawa, T.; Lopez, J.; Katsumata, T.; Xu, J.; Zhu, C.; Gu, X.;
Bae, W.; Kim, Y.; Jin, L.; Chung, J. W.; Tok, J. B. -.; Bao, Z. Intrinsically
Stretchable and Healable Semiconducting Polymer for Organic Transistors. Nature
2016, 539, 411-415.
2. Savagatrup, S.; Zhao, X.; Chan, E.; Mei, J.; Lipomi, D. J. Effect of Broken
Conjugation on the Stretchability of Semiconducting Polymers. Macromol. Rapid
Commun. 2016, 37, 1623-1628.
3. Howard, J. B.; Thompson, B. C. Design of Random and Semi‐ Random Conjugated
Polymers for Organic Solar Cells. Macromol. Chem. Phys. 2017, 218, 1700255.
4. Printz, A. D.; Savagatrup, S.; Burke, D. J.; Purdy, T. N.; Lipomi, D. J. Increased
Elasticity of a Low-Bandgap Conjugated Copolymer by Random Segmentation for
Mechanically Robust Solar Cells. RSC Adv. 2014, 4, 13635-13643.
5. O'Connor, B.; Chan, E. P.; Chan, C.; Conrad, B. R.; Richter, L. J.; Kline, R. J.;
Heeney, M.; McCulloch, I.; Soles, C. L.; DeLongchamp, D. M. Correlations
between Mechanical and Electrical Properties of Polythiophenes. ACS Nano 2010, 4,
7538-7544.
6. Awartani, O.; Lemanski, B. I.; Ro, H. W.; Richter, L. J.; DeLongchamp, D. M.;
O'Connor, B. T. Correlating Stiffness, Ductility, and Morphology of
Polymer:Fullerene Films for Solar Cell Applications. Adv. Energy Mater. 2013, 3,
399-406.
7. Burkhart, B.; Khlyabich, P. P.; Cakir Canak, T.; LaJoie, T. W.; Thompson, B. C.
“Semi-Random” Multichromophoric Rr-P3HT Analogues for Solar Photon
Harvesting. Macromolecules 2011, 44, 1242-1246.
8. Khlyabich, P. P.; Burkhart, B.; Ng, C. F.; Thompson, B. C. Efficient Solar Cells from
Semi-Random P3HT Analogues Incorporating Diketopyrrolopyrrole.
Macromolecules 2011, 44, 5079-5084.
9. Howard, J. B.; Ekiz, S.; Noh, S.; Thompson, B. C. Surface Energy Modification of
Semi-Random P3HTT-DPP. ACS Macro Lett. 2016, 5, 977-981.
10. Ekiz, S.; Gobalasingham, N. S.; Thompson, B. C. Exploring the Influence of
Acceptor Content on Semi‐ random Conjugated Polymers. J. Polym. Sci., Part A:
Polym. Chem. 2017, 55, 3884-3892.
101

11. Zhao, Y.; Zhao, X.; Roders, M.; Gumyusenge, A.; Ayzner, A. L.; Mei, J. Melt‐
Processing of Complementary Semiconducting Polymer Blends for High
Performance Organic Transistors. Adv. Mater. 2016, 29, 1605056.
12. Zhao, X.; Zhao, Y.; Ge, Q.; Butrouna, K.; Diao, Y.; Graham, K. R.; Mei, J.
Complementary Semiconducting Polymer Blends: The Influence of Conjugation-
Break Spacer Length in Matrix Polymers. Macromolecules 2016, 49, 2601-2608.
13. Zhao, Y.; Zhao, X.; Zang, Y.; Di, C.; Diao, Y.; Mei, J. Conjugation-Break Spacers in
Semiconducting Polymers: Impact on Polymer Processability and Charge Transport
Properties. Macromolecules 2015, 48, 2048-2053.
14. Burkhart, B.; Khlyabich, P. P.; Thompson, B. C. Semi-Random Two-Acceptor
Polymers: Elucidating Electronic Trends through Multiple Acceptor Combinations.
Macromol. Chem. Phys. 2013, 214, 681-690.
15. Verploegen, E.; Mondal, R.; Bettinger, C. J.; Sok, S.; Toney, M. F.; Bao, Z. Effects of
Thermal Annealing upon the Morphology of Polymer–Fullerene Blends. Adv. Funct.
Mater. 2010, 20, 3519-3529.
16. Khanafer, K.; Duprey, A.; Schlicht, M.; Berguer, R. Effects of Strain Rate, Mixing
Ratio, and Stress–strain Definition on the Mechanical Behavior of the
Polydimethylsiloxane (PDMS) Material as Related to its Biological Applications.
Biomed. Microdevices 2009, 11, 503-508.
17. Savagatrup, S.; Printz, A. D.; Rodriquez, D.; Lipomi, D. J. Mechanical Properties of
Organic Semiconductors for Stretchable, Highly Flexible, and Mechanically Robust
Electronics. Chem. Rev. 2017, 117, 6467-6499.

102

Chapter 5: Diketopyrrolopyrrole (DPP) pendant polymers

5.1 Motivation
5.1.1 Motivation for Pendant Polymers
Conjugated polymers (CPs) have many features that make them attractive candidates
for organic electronics, but they have some inherent weaknesses. In comparison to the non-
conjugated polymers frequently seen in household plastics, fully conjugated polymers tend
to be stiff and brittle with high tensile moduli and low ductility.
1
Conjugated polymers have
relatively high melting points and low solubilities, which inhibits their incorporation into
industrial-scale fabrication processes. The backbone of CPs is susceptible to
photooxidation and other degradation pathways which render the polymer incapable of
both absorbing light and transporting charge.
2-4
This degradation breaks the polymer down
into small oligomers which do not retain the mechanical, optical, or electronic properties
of the parent chain. Additionally, the twists and kinks that occur along the conjugated
polymer backbone at distorted sp
2
hybridized carbons create electronic defects that are
more susceptible to photodegradation.
5

The polymerization methods most commonly used to create conjugated polymers also
have some undesirable attributes. These polymerization techniques include Stille, Suzuki,
and other transition metal-catalyzed cross-couplings. Unlike the methods available for
creating non-conjugated polymers, these cross-coupling reactions allow for very little
control over polymer molecular weight and dispersity (i.e., they are not living
polymerizations). Since polymer semiconductor device performance has been shown to
depend on both molecular weight and dispersity, this leads in turn to batch-to-batch
103

variation in device performance.
6
What’s more, precious metal catalysts are necessary for
these polymerizations and a stoichiometric amount of often toxic byproducts are produced
from the condensation reactions. In contrast, non-conjugated polymers can be synthesized
through a great variety of methods, many of which do not require metal catalysts, which
are difficult to remove entirely after the reaction has concluded. Residual metals hinder
device performance
7
and can prevent the materials from being used in biomedical
applications.
Non-conjugated polymers with electroactive pendants possess a great deal of potential
for improving the physical and mechanical properties of semiconducting polymers.
Comprising the plastics that we use in everyday life, non-conjugated polymers are made
through a wide variety of synthetic techniques, many of which have been perfected over
decades of use in industrial fabrication processes to give precise end groups, molecular
weights, and backbone stereochemistry. Early studies have proven that charge transport
can occur along π-π stacks,
8
and does not necessitate a conjugated polymer backbone,
9,10

so if the electroactive pendants can arrange themselves in a manner favorable for π-π
stacking, charge transport should occur along the pendants. Electroactive pendants can be
added to a common monomer and polymerized directly (Figure 5.1B), or can be added
onto a non-conjugated backbone after polymerization via post-polymerization
modification (Figure 5.1A). Each method has its drawbacks: post-polymerization
modification rarely reacts to completion, leaving some units along the backbone unreacted
with no way of purifying the imperfections; direct polymerization requires electroactive
monomers and catalysts/initiatiors to be compatible. More optically interesting monomers
104

tend to be more sensitive to reaction conditions and can be difficult to polymerize using
controlled or living methods.

Figure 5.1. Two general approaches to creating non-conjugated polymers with
electroactive pendants. In A, monomer with a reactive functional group is first
polymerized, then coupled to the electroactive dye as a form of post-polymerization
modification. In B, the monomer is first coupled with the electroactive dye and then
polymerized directly.
In contrast to typical conjugated polymer syntheses, transition metal catalysts are not
necessary for many non-conjugated polymerization methods. Non-conjugated polymers
can be synthesized both through step growth (condensation) and chain growth (ionic,
radical) polymerization techniques. The former produces stoichiometric byproducts
(usually water) and the latter has no byproduct. This is in sharp contrast to the often toxic
stoichiometric byproducts produced during conjugated polymer synthesis. Non-conjugated
polymers can frequently be synthesized in the bulk, eliminating the need for toxic or
wasteful organic solvents. They may also exhibit improved solubility and mechanical
properties over fully conjugated polymers, although these parameters are rarely discussed
in semiconducting polymer literature. Finally, the aliphatic backbone of non-conjugated
polymers is stable to photooxidation, enhancing the lifetime of the polymer.
105

5.1.2 Rationale for DPP
Use of the diketopyrrolopyrrole (DPP) unit and its derivatives has been widely adopted
in the field of organic semiconductors. The dye had previously been known as a high-
performance pigment with excellent thermal and photo stability (Pigment Red 254, or
Ferrari red) but with the discovery of favorable optoelectronic properties it has rapidly
become a staple in the organic semiconductor community.
11-13
It serves as an electron-poor
monomer, broadening absorption and narrowing electronic bandgaps when coupled with
electron donating monomers. It is also quite stable to air and its derivatives can absorb deep
into the red region of the solar spectrum, making it useful for photovoltaic applications.
14,15

Its high fluorescence quantum yields and biological inertness also make it an attractive
biological fluorescent probe.
16
The DPP core has a wide range of versatility for
functionalization. Most frequently, an aryl ring is attached at positions 3 and 6, extending
the conjugation of the compound (Figure 5.2). The resulting compounds can vary widely
in color depending on the nature of the aryl group, and remain fairly planar, allowing the
molecules to stack in the solid state.
17
As the core tends to be insoluble in all but the most
polar aprotic solvents, the amide nitrogens at positions 2 and 5 are also frequently alkylated
to increase the range of solubility.

Figure 5.2. The fundamental diketopyrrolopyrrole (DPP) unit and the form typically used
in organic electronics, with aryl substituents at the 3 and 6 positions and solubilizing alkyl
chains on the amide nitrogens.
106

Although conjugated polymers incorporating derivatives of the DPP monomer have
excellent charge transport values (> 1 cm
2
V
-1
s
-1
),
18-20
the isolated small molecule shows
more moderate values (< 0.01 cm
2
V
-1
s
-1
).
21
Despite their many handles for
functionalization and their attractive planar nature, DPP derivatives have yet to be
incorporated as pendants onto a non-conjugated polymer backbone. This strategy could
allow for the DPP pendants to stack in a manner favorable for charge transport, while the
non-conjugated backbone could imbue the polymer with improved mechanical properties
and processibility. For these reasons, DPP was selected to be used as an electroactive
pendant on a non-conjugated backbone, pursued by the various strategies detailed within.
5.2 Background
5.2.1 Direct polymerization vs. post-polymerization modification
Electroactive pendants can be added to a common monomer and polymerized directly
(Figure 5.1B) or can be added onto a non-conjugated backbone after polymerization via
post-polymerization modification (Figure 5.1A). Each method has its drawbacks: post-
polymerization modification rarely reacts to completion, leaving some units along the
backbone unreacted with no way of purifying the imperfections; direct polymerization
requires electroactive monomers and catalysts/initiatiors to be compatible. More optically
interesting monomers tend to be more sensitive to reaction conditions and can be difficult
to polymerize using controlled or living methods. In 2010, Lang et al. published a
comparative study of polyacrylates with perylene diimide pendants, synthesized either
through a direct controlled radical polymerization of the perylene acrylate monomer or
through a post-polymerization copper-catalyzed azide-alkyne “click” reaction of a
perylene azide onto a poly(propargyl acrylate).
22
Using nitroxide-mediated radical
107

polymerization (NMRP) in the direct polymerization of the perylene acrylate monomer,
they showed that the polymerization kinetics were not first order, and the molecular weight
dispersity increased with monomer conversion, neither of which phenomena were
observed in typical acrylate NMRPs. The authors attributed the broad dispersity and
unpredictable molecular weight of these polymers to monomer’s poor solubility, steric
hindrance, high molecular weight, and transfer reactions that occurred during the direct
polymerization. Despite this, they were still able to achieve relatively high polymer
molecular weights by direct polymerization, with number-average molecular weights
above 10 kDa. These molecular weights were comparable to those obtained by post-
polymerization “clicking” of the perylene unit onto the non-conjugated backbone.
However, the “clicked” polymers had a lower molecular weight than anticipated given the
degree of polymerization of the initial unmodified polymer. The authors attribute this to an
underestimation of the “clicked” polymer mass due to aggregation, instead of incomplete
functionalization leading to fewer than expected perylene azides “clicking” onto the
polymer backbone. However, assuming that each step reacts to completion, the post-
polymerization modification should lead to more uniform pendant polymers with narrow
molecular weight distributions.
The example highlighted above is illustrative of the challenges facing this field.
Electroactive polymers made via the post-polymerization functionalization technique
frequently suffer from low charge mobilities due to low attachment density. This can be
attributed to incomplete conversion (intentional or otherwise) of functional groups along
the backbone to pendant moieties. Alternatively, the electroactive moiety may be diluted
with another monomer to attain some other desirable property such as elasticity, self-
108

assembly,
23
crosslinking, water solubility,
24
or a desired weight ratio of chromophores.
25

On the other hand, direct polymerization of an electroactive pendant monomer may lead to
a greater density of chromophores along the polymer backbone, but often a lower
molecular weight due to poor solubility. Additionally, fewer synthetic options may be
available in terms of both functionalizing the chromophore with a polymerizable group as
well as the polymerization itself based upon the complexity, solubility, and reactivity of
the electroactive group used. For example, in the study highlighted above, a pre-clicked
acrylate monomer was synthesized, but could not be polymerized due to its poor
solubility.
22
The decision to pursue an electroactive pendant polymer from the post-
polymerization modification approach or direct polymerization must be made on an
individual basis and is informed both by the type of monomer and pendant chosen and the
properties desired from the final polymer product.
5.2.2 Stereoregularity effects
While electroactive pendant polymers have been studied for several decades, in general
the charge mobility of the materials is quite low because charge transport is fundamentally
restricted to two dimensions (along π-stacked pendants and across polymer chains) as
opposed to the three-dimensional charge transport available to polymers with conjugated
backbones. Exceptions to the low mobility include instances in which the alkyl chain
attaching the pendant to the backbone is so long that the pendants are able to arrange
themselves in a liquid crystalline phase;
26,27
the alkyl chain attaching the pendant to the
backbone is nonexistent (as in the case of poly(dibenzofulvene), pDBF), forcing the
pendants to arrange themselves cofacially;
28,29
or if the stereochemistry of the backbone is
precisely controlled so that the pendants all hang off of the backbone in the same direction
109

(isotactic).
30,31
The use of backbone stereoregularity to control pendant alignment and
influence mobility has been little explored since the first reports in the 1980s, though recent
reports have begun to investigate the influence of backbone tacticity on optical
properties.
32,33
Several modeling papers have also indicated that stereoregularity is
theoretically expected to contribute to intramolecular excimer formation in electroactive
pendants.
34,35
All of these studies have been limited to simple electroactive monomers such
as carbazole derivatives. No attempt has yet been made to explore the influence of
backbone stereoregularity on charge mobility in more complex electroactive pendants.
5.3 Thesis
In this chapter, the synthesis of novel non-conjugated polystyrene-based polymers with
diketopyrrolopyrrole pendants is reported. Post-polymerization modification and direct
polymerization strategies are compared for this system, and the effect of backbone
stereoregularity on optoelectronic properties for these polymers is explored.
5.4 Results and Discussion
5.4.1 Synthesis
There are several methods that can be used to create stereoregular polymers. Radical
polymerization using bulky Lewis Acid salts that coordinate to an acrylate monomer can
direct subsequent monomer addition.
36,37
This strategy had been pursued by our group
before, but resulted in polymers with only ~70 % stereoregularity – not enough to see its
effects on polymer properties. Anionic polymerization can be used to give highly
stereoregular polymers, but it is generally intolerant of polar functionalities (necessary
either for post-polymerization modification or present on optically interesting monomers)
and requires stringently air- and water-free reaction conditions. Ziegler-Natta or
110

metallocene polymerizations can also produce highly stereoregular polymers, but are also
intolerant of polar functionalities on the monomer. This latter method is more tolerant of a
variety of reaction conditions, so two stereoregular polymers were synthesized using
metallocene/MAO catalyst systems followed by several post-polymerization
modifications, as shown in Scheme 5.1. A half-metallocene titanium catalyst
(Cyclopentadienyltitanium(IV) trichloride) was used to synthesize syndiotactic polymers
and a C2 symmetric zirconium catalyst (rac-Dimethylsilylbis(1-indenyl)zirconium(IV)
dichloride) was used to synthesize isotactic polymers, as detailed in Appendix D. For
comparison, an atactic polymer was also synthesized via free radical polymerization using
AIBN as an initiator. The relatively non-polar 4-methylstyrene (X) was selected as the
monomer because monomers with polar functional groups are known to react with the
catalyst system, deactivating or drastically slowing polymerization;
38
the benzylic methyl
group is easily halogenated; and atactic polystyrenes can be separated from their
stereoregular counterparts via Soxhlet extraction.
39
Due to impurities in the reaction, both
the isotactic and syndiotactic polymerizations produced a mixture of atactic and
stereoregular polymers. Atactic polymer was removed from the stereoregular polymer via
a 24-hour exhaustive Soxhlet extraction in methyl ethyl ketone (MEK).

Scheme 5.1. Polymerization and post-polymerization modification of 4-methylstyrene. i)
AIBN, toluene, 90 °C, 2 hrs (XI-a); MMAO/Zr(IV) catalyst, toluene, RT, 2 days (XI-i);
MMAO/Ti(IV) catalyst, toluene, RT, 2 days (XI-s); ii) AIBN, NBS, CCl4, 85 °C, 1 hr (XI-a
and XI-s); iii) NaN3, DMF, 55 °C, 4 days (XI-a and XI-s)
111

Stereochemistry of the polymers was determined by carbon nuclear magnetic
resonance (NMR) spectroscopy (Figures D12-D14, D16-D18).
40-43
Deconvolution of the
aromatic peak between 141-145 ppm in the (XI-a)
13
C spectrum revealed triads of reveals
26% mm, 39% mr/rm, and 35% rr (Figure D13) while deconvolution of the 142.8 ppm
peak in the (XI-s)
13
C spectrum revealed only one peak, indicating rr triads > 95% (Figure
D17).
40
Representing the methine carbon in the polymer backbone, the atactic polymer
(XI-a) exhibited a broad series of peaks between 42-46 ppm (Figure D14) and the
syndiotactic polymer (XI-s) had a sharp peak at 44.4 ppm (Figure D18). Deconvolution of
this peak again yielded a single peak, confirming the highly syndiotactic nature of the
polymer. A clear NMR of the isotactic polymer was never obtained due to poor solubility;
however, this solubility difference between the isotactic and syndiotactic polymers was
indicative of a difference in stereochemistry. The molecular weights and degree of
polymerization for these polymers are reported in Table 5.1 before any post-
polymerization modification took place. The size-exclusion chromatography instrument is
calibrated using polystyrene standards, and these initial polymers are assumed to more
closely match the hydrodynamic volume of the calibration standards. Additionally, poor
solubility with subsequent post-polymerization modifications prevented further
measurements of polymer molecular weight.

112

Table 5.1. Synthetic, optical, and electronic data for polymers and small molecule
reference.
Material Mn (n)
a

(kDa)
Ð
a

Eg, film
b
/ solution
c

(eV)
HOMO
d

(eV)
iPS
e
(XI-i) 38.8 (328) 2.41 - -
aPS-azide-click-DPP (XV-a) 66.3 (561) 1.89 2.09 / 2.08 5.34
sPS-azide-click-DPP (XV-s) 36.9 (312) 3.44 2.10 / 2.08 5.35
aPS-alkyne-click-DPP (IX) 13.3 (47) 1.81 2.08 / 2.12 5.65
dtdDPP (IX)
f
- - 1.99 / 2.17 5.95
a) Determined by SEC for polystyrenes before any post-polymerization modification,
degree of polymerization in parentheses; b) optical bandgap determined from absorption
onset in films spin-cast from chloroform and placed in N2 cabinet for 30 min; c) optical
bandgap determined from absorption onset in dilute solutions in chloroform; d) HOMO
energy estimated from cyclic voltammogram oxidation onset vs Fc/Fc
+
; e) isotactic
polymer insoluble and could not be further reacted or characterized, SEC measurements
performed on mixed tacticity polymer before isolation of isotactic polymer by Soxhlet
extraction; f) small molecule reference – not measured by SEC

The syndiotactic and atactic polymers were brominated using NBS and AIBN (XII-a
and XII-s). However, because the isotactic polymer would not dissolve in any solvent but
toluene (which would also have been brominated as an undesired side reaction), it could
not be brominated. This also prevented intelligible characterization of the polymer. Due to
its insoluble nature, further experiments with the isotactic polymer were not pursued. After
bromination of the atactic and syndiotactic polymers, the halogen was converted to an azide
via SN2 substitution to yield (XIII-a) and (XIII-s). Extreme caution was used at this stage
due to the potential to create an explosive polymer. Tests on the azide-substituted polymers
revealed that they are impact sensitive, but not thermally or friction sensitive. The risk of
explosion from impact was deemed to be low enough that the polymers could be handled
safely. Both the atactic and syndiotactic polymers exhibited decreased solubility with
progressive post-polymerization modifications. Toluene was the most effective solvent of
113

poly(4-methylstyrene) and its derivatives due to its structural similarity to the 4-
methylstyrene monomer. However, clear spectra could not be obtained when using toluene
as a solvent because its characteristic peaks overlapped with the monomer peaks, making
it the least useful solvent for characterization purposes.
A 3,6-thiophene substituted DPP core was synthesized according to literature
procedures,
44
followed by an unsymmetrical alkylation with the solubilizing 2-
decyltetradecyl chain and propargyl group, as shown in Scheme D5 to yield the
propyne/dtdDPP (XIV) shown in Scheme 5.2. The synthesis was also attempted with the
shorter 2-ethylhexyl chain, but the limited solubility of the product made it difficult to
handle. The electroactive molecule was attached to the azide-substituted polymers via a
copper-catalyzed “click” reaction (Scheme 5.2), which was monitored by infrared (IR)
spectroscopy until completion (Figures D36-D37). Unreacted DPP was washed away with
a methanol Soxhlet, and the polymers were collected in the chloroform fraction. Visibly,
the two polymers (XV-a) and (XV-s) appeared the same: red and non-fluorescent in
solution; dark purple as a solid. They dissolved well in common organic solvents, such as
toluene, chloroform, chlorobenzene, and o-dichlorobenzene, though the syndiotactic
polymer took noticeably longer to dissolve than the atactic polymer.
114

Scheme 5.2. Copper-catalyzed azide-alkyne “click” reaction. i) CuSO4·5H2O, Sodium
Ascorbate, DMF/H2O/toluene, RT, 3 days
For comparison, and for safety reasons, the azide and alkyne functionalities were
reversed in an alternative synthesis. A protected ethynyl styrene monomer was synthesized
(II), followed by polymerization (III), then finally deprotection with tetrabutylammonium
fluoride (TBAF) (IV) (Scheme 5.3). The same catalysts used to produce the syndiotactic
and isotactic polystyrenes above were used on the protected monomer (II), but each
resulted in an insoluble white powder that could not be characterized. A simple free radical
polymerization with AIBN was chosen therefore to make a soluble atactic polymer with
reversed azide/alkyne functionalities. At approximately 47 repeat units, this polymer (III)
had a much lower degree of polymerization than all of the (XI) polymers (all above 300
repeat units) (Table 5.1).
115

Scheme 5.3. Functionalization, polymerization, and deprotection of ethynyl styrene. i)
PdCl2(PPh3)2, TIPS acetylene, triethylamine, 50 °C, 16 hrs; ii) AIBN, toluene, 70 °C, 10
hrs; iii) TBAF, THF, 0 °C to RT, overnight
A 3,6-thiophene substituted DPP core was synthesized according to literature
procedures,
44
followed by a monoalkylation with the solubilizing 2-decyltetradecyl chain,
as shown in Scheme D2 (VI). After purification via column chromatography, the molecule
was again subjected to monoalkylation using a large excess of 1,4-dibromobutane (VII).
The remaining halogen on the DPP molecule was then converted to an azide (VIII) and
the electroactive molecule was attached to the alkyne-substituted polymer via a copper-
catalyzed “click” reaction (Scheme 5.4) (IX). Unreacted DPP was separated from the
polymer using column chromatography and its removal confirmed by the absence of azide
stretching signals in IR spectra. Also red and non-emissive in solution and purple as a solid,
this polymer dissolved rapidly in common organic solvents as well.
116

Scheme 5.4. Copper-catalyzed azide-alkyne “click” reaction. i) CuSO4·5H2O, Sodium
Ascorbate, DMF/H2O, RT, 24 hrs
For comparison with the DPP pendant polymers made through post-polymerization
modification above, several direct polymerization strategies were also pursued. A “pre-
clicked” monomer was made following the synthetic strategy outlined in Scheme 5.4,
except that the alkyne functionality was deprotected and the DPP pendant attached before
polymerization, yielding monomer (XVIII) in Scheme D8. While useful as a model
compound, attempts to polymerize this monomer via AIBN-initiated free radical
polymerization failed, resulting in decomposition of the monomer. Another attempt at
direct polymerization was made with an acrylate-functionalized DPP monomer, the
synthesis of which is outlined in Scheme D9, yielding monomer (XX). This monomer was
subjected to n-butyllithium to initiate an anionic polymerization, but this also failed to
produce any polymer. These two failed attempts at direct polymerization demonstrated that
for this system, post-polymerization modification seems to be a more viable option to
create electroactive pendant polymers.

117

5.4.2 Optical Properties

Figure 5.3. Solution (left) and film (right) absorption spectra for three clicked polymers
and small molecule reference. Solution spectra collected in chloroform; films spin-cast
from chloroform solutions and kept in a N2 cabinet for 30 min.
The three clicked polymers (IX), (XV-a), and (XV-s) were characterized along with a
disubstituted decyltetradecyl (dtd) DPP small molecule (structure (XVI) in Scheme D7)
for comparison (Table 5.1). All four materials were characterized by UV-Visible
spectroscopy in both film and solution, using chloroform as a solvent. The three polymers
and small molecule exhibited nearly identical absorption spectra in solution (Figure 5.3),
though in film the small molecule differed significantly in both shape and intensity from
the three polymers (Figure 5.3). For the polymers, the solution and film spectra were
approximately the same shape, indicating that the arrangement of the electroactive
monomers likely did not differ significantly in the solid state from in dilute solution
(Figures D41-D43). This is in contrast to the small molecule DPP in which the difference
between the film and solution spectra indicated a significant reorganization of the molecule
in the solid state (Figure D40). All of the films had a slightly red-shifted absorption onset
from the solution onset, indicating enhanced intermolecular packing in the solid state.
However, the shift is minimal in the polymers, leading to a change in bandgap of at most
118

0.04 eV (Table 5.1), again indicating minimal rearrangement of the DPP molecules in the
solid state. In contrast, the small molecule dtdDPP had a bandgap change of approximately
0.2 eV when going from solution to film.

Figure 5.4. Fluorescence spectra of three clicked polymers and reference small molecule
in chloroform solution.
The polymers were also studied by fluorescence spectroscopy, in which the two azide-
click-DPP polymers (XV-a) and (XV-s) shared approximately the same shape as the
dtdDPP small molecule (XVI), but the alkyne-click-DPP polymer (IX) had a much broader
emission spectrum (Figure 5.4). Though the three polymers have identical electroactive
units, the aPS-alkyne-click-DPP polymer (IX) places the DPP unit 7 atoms from the
polystyrene backbone whereas the azide-click DPP polymers (XV-a) and (XV-s) maintain
a distance of 5 atoms between the DPP unit and the polystyrene backbone. This difference
in spacer length between the electroactive pendant and the polymer backbone could lead
to a difference in DPP pendant packing between the two types of polymers, which in turn
could alter the optoelectronic properties of the polymer. Other differences include the
degree of polymerization, which is approximately an order of magnitude higher in the
azide-click-DPP polymers (XV-a) and (XV-s) than in the alkyne-click-DPP polymer (IX).
The photoluminescence quantum yield (PLQY) was measured in film and solution (Table
119

D2). In solution, the small molecule fluoresced strongly with a yield of 64 %, but in the
polymers fluorescence was largely quenched, with yields below 4 % (below the instrument
error of ± 10 %). When electrons are excited into a higher energy level, they can relax
either by a radiative process, such as fluorescence, or a non-radiative process, such as
charge transfer. Many of the non-radiative relaxation pathways available to a molecule in
its excited state require that there be other molecules in close proximity, which rarely
happens in a dilute solution. For this reason, the small molecule (XVI) had a relatively high
quantum yield. The quenching observed in pendant DPP polymers is indicative of more
non-radiative decay pathways available to the polymers than the small molecule, and could
indicate that the DPP pendants were π-stacked along the polymer backbone. The
fluorescence for the polymers and the small molecule was further diminished in film, with
dtdDPP having a quantum yield of 4 % and the polymers having quantum yields of 0 %,
all of which are below the instrument error of ± 10 %. This indicates that in the solid state,
non-radiative decay processes dominate for the three polymers as well as the small
molecule reference.
5.4.3 Structural Properties
The thermal transitions of the clicked polymers and small molecule reference
compound were investigated using differential scanning calorimetry (DSC). Only the small
molecule dtdDPP exhibited thermal transitions between 0 °C and 300 °C, with a cold
crystallization at 21 °C and a melt at 77 °C (Figure D49). None of the polymers appeared
to undergo any thermal transitions within this temperature range (Figure D50-D52). This
was supported by X-ray diffraction, which showed a 100 peak only for the small molecule
120

(Figure D53). Thus under the conditions in this study, all three polymers appeared to be
totally amorphous.
5.4.4 Electronic properties
The highest occupied molecular orbital (HOMO) energies of the polymers were
approximated using oxidation onset in cyclic voltammetry (CV) (Table 5.1). Interestingly,
the three polymers reached their peak oxidation at the same applied voltage, but their onsets
differed greatly (Figure D50). The two azide-click-DPP polymers (XV-a) and (XV-s) had
nearly identical HOMOs at 5.34 eV and 5.35 eV, whereas the alkyne-click-DPP polymer
(IX) had a deeper HOMO at 5.65 eV. The difference in frontier orbital levels between the
two types of polymers indicated that there may be a difference in intermolecular electron
delocalization between the azide-click-DPP polymers (XV-a) and (XV-s) and the alkyne-
click-DPP polymer (IX). In general, HOMO energies become more shallow as electron
delocalization increases,
45
as would normally be associated with a decrease in the bandgap
with extended conjugation.
46
However, as illustrated in both Figure 5.3 and Table 5.1, the
bandgap does not change appreciably between the three polymers and the small molecule
dtdDPP, and neither the polymers nor the small molecule have any sort of extended
conjugation. Therefore it can be assumed that delocalization in these systems occurs not
along a conjugated backbone, but between π-stacked pendant DPPs. Thus the relatively
shallow HOMOs of the azide-click-DPP polymers (XV-a) and (XV-s) revealed promising
electron delocalization that may lead to improved charge mobilities. The small molecule
dtdDPP had the deepest HOMO at 5.95 eV, as would be expected from minimal electron
delocalization. However, this molecule seemed to undergo electropolymerization and the
HOMO was estimated from the first CV cycle.
121

To date, we have been unable to measure the mobility via the space-charge limited
current (SCLC) technique.
5.5 Summary and Conclusion
In this study, three novel polymers with a polystyrene backbone and electroactive DPP
pendants were synthesized and characterized. Post-polymerization modification was the
only effective method found to synthesize DPP pendant polymers, as two model systems
with the DPP pre-attached to a monomer (a DPP acrylate and a pre-clicked DPP styrene)
could not be polymerized either by radical or anionic means. UV-vis spectroscopy
indicated that all three polymers behaved similarly in solution and the solid state, while the
small molecule analog differed in its solution and solid-state behavior, indicating
chromophore rearrangement. PLQY measurements in film and solution showed
fluorescence quenching in the polymer solutions relative to the small molecule, suggesting
π-π stacking of the electroactive pendants. Electrochemical measurements revealed more
electron delocalization among the DPP pendants of the polymers than in the dtdDPP small
molecule, as expected, while the azide-click-DPP polymers (XV-a) and (XV-s) exhibited
more delocalization than the alkyne-click-DPP polymer (IX).

122

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126

Appendix A: Influence of Systematic Incorporation of Conjugation-Break Spacers
into Semi-Random Polymers on Mechanical and Electronic Properties

A.1 Materials and Methods
All reagents from commercial sources were used without further purification, unless
otherwise noted. For mechanical measurements, PEDOT:PSS (Clevios PH 1000) was
purchased from Heraeus and used as received. For mechanical measurements, chloroform,
acetone, and isopropyl alcohol were purchased from Sigma–Aldrich Co. and used as
received. For synthesis and all other measurements, solvents were purchased from VWR
and used without purification with the exception of acetonitrile, which was distilled from
CaH2 prior to use. All reactions were performed under dry N2, unless otherwise noted. All
reactions were performed with glassware that was oven-dried and then flamed under high
vacuum and backfilled with N2. Flash chromatography was performed on a Teledyne
CombiFlash Rf instrument with RediSep Rf normal phase disposable columns.
1
H NMR
spectra were recorded in CDCl3 on a Varian Mercury 400 NMR spectrometer (small
molecules) or a Varian Mercury 600 NMR spectrometer (polymers). For molecular weight
determination, polymer samples were dissolved in HPLC grade o-dichlorobenzene (o-
DCB) at a concentration of 0.5 mg/mL and filtered through a 0.2 μm PTFE filter. Size
exclusion chromatography (SEC) was performed using HPLC grade o-DCB at a flow rate
of 0.6 mL/min on one 300 × 7.8 mm TSK-Gel GMHHR-H column (Tosoh Corp.) at 60 °C
using a Viscotek GPC Max VE 2001 separation module and a Viscotek TDA 305 RI
127

detector. The instrument was calibrated with polystyrene standards (1,050 – 3,800,000
g/mol), and data were analyzed using OmniSec 4.6.0 software.
For thin film measurements, solutions were spin-cast onto pre-cleaned 2.5 cm
2
glass
slides (sonicated for 10 minutes in water, 5 minutes in acetone, and 5 minutes in isopropyl
alcohol then dried under high N2 flow) from 7 mg mL
-1
o-dichlorobenzene (o-DCB)
solutions. UV-Vis absorption spectra were obtained on a PerkinElmer Lamda 950
spectrophotometer. The thickness of the thin films and grazing-incidence X-ray diffraction
(GIXRD) measurements were obtained using a Rigaku Diffractometer Ultima IV using Cu
Kα radiation source (λ = 1.54 Å) in the reflectivity and grazing-incidence mode,
respectively. Differential scanning calorimetry (DSC) traces were obtained using a Perkin-
Elmer DSC 8000 with a scan rate of 10 °C/min. Sample size was ∼5 mg, and polymers
were used as obtained after purification.
Cyclic voltammetry (CV) was performed on Princeton Applied Research VersaStat3
potentiostat under the control of VersaStudio Software. A standard three-electrode cell
based on a platinum disc working electrode, a silver wire pseudo reference electrode
(calibrated vs Fc/Fc
+
which is taken as 5.1 eV vs vacuum),
1,2
and a Pt wire counter
electrode was purged with nitrogen and maintained under a nitrogen atmosphere during all
measurements. Polymer films were made by repeatedly dripping a 1% (w/w) chloroform
or chlorobenzene solution onto the Pt disc and drying under nitrogen prior to measurement.
Tetrabutylammonium hexafluorophosphate (0.1 M in freshly distilled acetonitrile) was
used as the supporting electrolyte.

128

A.1.1 SCLC Device Fabrication and Characterization
All steps of device fabrication and testing were performed in air. ITO-coated glass
substrates (10 Ω/m, Thin Film Devices Inc.) were sequentially cleaned by sonication in
detergent, de-ionized water, tetrachloroethylene, acetone, and isopropyl alcohol, and dried
in a nitrogen stream. A thin layer of PEDOT:PSS (Baytron® P VP AI 4083, filtered with
a 0.45 μm PVDF syringe filter – Pall Life Sciences) was first spin-coated on the pre-cleaned
ITO-coated glass substrates and baked at 120 ºC for 60 minutes under vacuum. Solutions
of polymers were prepared in o-dichlorobenzene or chloroform solvent at a concentration
of 7 mg/ml and stirred for overnight. Subsequently, the polymer active layer was spin
coated (with a 0.45 μm PTFE syringe filter – Pall Life Sciences) on top of the PEDOT:PSS
layer. Upon spin coating of polymers, films were first placed under N2 for 30 min or were
annealed at 150 ºC for 30 minutes under N2 and then placed in the vacuum chamber for
aluminum deposition. At the final stage, the substrates were pumped down to high vacuum
(< 2.5×10
-6
Torr) and aluminum (100 nm) was thermally evaporated at 3 – 5 Å/sec using a
Denton Benchtop Turbo IV Coating System onto the active layer through shadow masks
to define the active area of the devices as 5.18 mm
2
. The current − voltage (I−V)
characteristics of the devices were measured under ambient conditions using a Keithley
2400 source-measurement unit.
A.1.2 Preparation of Substrates for Mechanical Measurements
Glass slides cut into 2.5 cm × 2.5 cm squares using a diamond–tipped scribe were used
as the substrate for the spin coating of the polymers. The glass slides were cleaned in
successive sonication baths of Alconox dissolved in deionized water, pure deionized water,
acetone, and isopropyl alcohol for 10 min each. Post sonication, the glass slides were dries
129

using a stream of compressed air. In order to improve wettability, activate the surface of
the glass, and remove any residual organic debris, the slides were treated with 30 W air
plasma for 5 min at a base pressure of 200–300 mTorr.
A.1.3 Preparation of Films for Mechanical Measurements
Solutions of pure polymers were mixed with chloroform in concentrations of 10 mg
mL
-1
and left to stir overnight. After mixing overnight, the solutions were filtered with 1
µm glass fiber media syringe filters. A layer of PEDOT:PSS spun at 1000 rpm (500 rpm s
-
1
ramp) for 3 min, followed by a second step at 2000 rpm (1000 rpm s
-1
ramp) for 30 s.
Filtered solutions were all spun at 1000 rpm (500 rpm s
-1
ramp) for 2 min, followed by
2000 rpm (1000 rpm s
-1
ramp) for 30 s. The thicknesses of the films were determined using
a Veeco Dektak stylus profilometer, in which at least 5 measurements were made per
sample.
A.1.4 Preparation of PDMS Elastomers
For tensile (compression) testing, poly(dimethylsiloxane) (PDMS) was chosen as the
substrate for mechanical measurements. To prepare, 20 g (50 g) of Sylgard 184 Silicone
elastomer base was mixed with 2 g (5 g) of Sylgard 184 Silicone curing agent and stirred
until cloudy. The mixture was then spread into a petri dish with a diameter of 15 cm and a
height of 1.5 cm. The PDMS was degassed by placing the petri dish was placed in a
desiccator under vacuum until the bubbles ceased to be visible. The dish was then placed
in an oven preheated to 70 °C for 45 to 50 min, allowing the PDMS to cure. Next, the
PDMS, with an approximate thickness of 1 mm (3 mm), was cut into 1 cm × 9 cm
rectangular slabs. The elastic modulus of PDMS elastomers was determined to be 0.5 ± 0.2
MPa, on average, using an Instron pull tester.
130

A.1.5 Buckling-Based Metrology for Measuring Elastic Moduli
Elastic moduli were estimated by applying Stafford et al.’s buckling-based metrology
to the wrinkling of the films under compression. (The resulting mechanical buckling
instability produces visible patterns of wrinkling from which the elastic modulus may be
determined.) As such, neat PDMS slabs were strained to approximately 5% on a linear
translation stage and fixed to rectangular glass slides. For each of the materials, thin films
prepared at three different thicknesses (using spin speeds of 500, 1000, and 1500 rpm with
ramp rates of 250, 500, and 750 rpm s
−1
, respectively) were scored and transferred to the
pre-strained PDMS substrates. The pre-strain in the PDMS was then released, engendering
a buckling instability and, in turn, wrinkles in the films. For each film, the wrinkles were
imaged under a microscope at numerous, random locations. To count the number of
wrinkles in an image, we used a MATLAB function based on the Savitzky-Golay
smoothing filter and peak finder, which distinguishes between crests and troughs. The
width of the image was then divided by the average number of wrinkles to compute the
buckling wavelength (λ), the thickness of the polymer film (hf) was measured (on glass)
using a profilometer, the elastic modulus of the PDMS substrate (Es) was determined using
a commercial pull tester, and the elastic modulus of the film (Ef) was estimated using
equation 1

𝐸 f
(1 − 𝜈 f
2
)
=
3𝐸 s
(1 − 𝜈 s
2
)
(
𝜆 2𝜋 ℎ
f
)
3
(1)
where νf (= 0.35) and νs (= 0.5) are the Poisson ratios of the film and substrate,
respectively.
5

131

A.1.6 Tensile Testing of Pseudo–Freestanding Films
The prepared sample is transferred to the surface of the water by partially submerging
the glass slide, which causes the layer of PEDOT:PSS to dissolve and the polymer to
delaminate from the surface. Once floating on the surface of the water, the film is brought
into contact with polydimethylsiloxane (PDMS) grips, prepared in the same methodology
as Alkhadra et al.,
6
attached to the load cell, and van der Waals adhesive forces keeps the
two in contact. The force versus displacement plots were obtained by uniaxially straining
the sample at 0.4 mm min
-1
while simultaneously recording the force trace until the sample
fractured. At least three tests were performed for each material. The stress–strain curves
were derived from the force versus displacement curves using the dimensions of the
corresponding sample.

132

A.2 Synthetic Procedures
The synthetic procedures for 2-bromo-5-trimethyltin-3-hexylthiophene (3HT) and 2,5-
diethylhexyl-3,6-bis(5-bromothiophene-2-yl)pyrrolo[3,4-c]-pyrrole-1,4-dione (DPP) were
previously published.
3
1,X-bis((5-trimethylstannyl)thiophen-2-yl)Xane alkyl spacers (T-
X-T, where X = the number of carbons separating the two thiophenes) were synthesized
according to the literature.
4
Typical polymerization (30% T-6-T): 163 mg (0.40 mmol) of
3HT, 205 mg (0.30 mmol) DPP, and 172 mg (0.30 mmol) T-6-T were dissolved in 25 mL
of dry DMF to give a 0.04 M solution. The solution was then degassed for 10 min before
46 mg (4 mol %) of Pd(PPh3)4 was added in one portion. The solution was degassed for an
additional 5 min and then heated for 48 h at 95 °C. Then the reaction mixture was cooled
to room temperature and precipitated into methanol. Purification was achieved via Soxhlet
extraction using methanol, hexanes, and chloroform. The chloroform fraction was then
reprecipitated into methanol, vacuum filtered, and dried. Yield 72% (213 mg), Mn =
14,400; Ð = 5.42.
133

A.3 Synthetic data
Table A1. Number-averaged polymer molecular weights in kDa as measured by SEC
Mn (kDa) T-4-T T-6-T T-8-T T-10-T
10% 10.5 29.0 19.7 14.0
20% 8.4 17.6 14.2 12.4
30% 6.4 14.4 9.9 9.8
40% 7.4 10.5 8.8 12.2
P3HTT-DPP-10% 9.5

Figure A1. Trends in number-averaged polymer molecular weights as measured by SEC

0
5
10
15
20
25
30
35
10% 20% 30% 40%
M
n
(kDa)
Percent Incorporation of CBS
T-4-T
T-6-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-6-T)
Linear (T-8-T)
Linear (T-10-T)
134

Table A2. Polymerization yields from chloroform Soxhlet fraction
Yield (%) T-4-T T-6-T T-8-T T-10-T
10% 78% 62% 78% 70%
20% 86% 80% 78% 72%
30% 79% 72% 63% 76%
40% 61% 50% 64% 68%
P3HTT-DPP-10% 95%

135

A.4 Polymer Nuclear Magnetic Resonance Spectra

Figure A2.
1
H NMR spectrum of 20% T-10-T. The NMR was taken in CDCl3 at 50 °C on
a 600 MHz instrument. This spectrum is being used as an example of how the amount of
monomer in each polymer was calculated. The aromatic peak at 8.88 ppm from the DPP
monomer is highlighted in blue; the aromatic peak at 6.73 ppm from the T-10-T monomer
is in orange; and the aromatic peak at 6.97 ppm from the 3HT monomer is in yellow. The
integrations of the DPP and T-10-T peaks match at 1.00 and 1.01, as they should. Each of
these peaks account for 2 hydrogens on these symmetrical monomers. The 3HT peak in
yellow accounts for only one hydrogen on the asymmetric monomer, therefore its
integration is doubled from 1.46 to 2.92. If these 3 integrations are summed, they give a
value of 4.93. Dividing each integration by this sum gives the percentage of each monomer
in the polymer chain. Therefore in this example, there is 20.3% DPP, 20.5% T-10-T, and
59.2% 3HT incorporated. This holds true to the feed ratio of 20% DPP, 20% T-10-T, and
60% 3HT.

136

Figure A3. Stacked
1
H NMR spectra of 10% T-4-T, 20% T-4-T, 30% T-4-T, and 40% T-4-
T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from
DPP monomer are highlighted in blue; T-4-T in orange; and 3HT in yellow.

137

Figure A4. Stacked
1
H NMR spectra of 10% T-6-T, 20% T-6-T, 30% T-6-T, and 40% T-6-
T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from
DPP monomer are highlighted in blue; T-6-T in orange; and 3HT in yellow.

138

Figure A5. Stacked
1
H NMR spectra of 10% T-8-T, 20% T-8-T, 30% T-8-T, and 40% T-8-
T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from
DPP monomer are highlighted in blue; T-8-T in orange; and 3HT in yellow.

139

Figure A6. Stacked
1
H NMR spectra of 10% T-10-T, 20% T-10-T, 30% T-10-T, and 40%
T-10-T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks
from DPP monomer are highlighted in blue; T-10-T in orange; and 3HT in yellow.

140

A.5 Film UV-Vis Absorption Spectra

Figure A7. UV-Vis absorption spectra of 10% CBS subfamily thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference.

141

Figure A8. UV-Vis absorption spectra of 20% CBS subfamily thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference.

142

Figure A9. UV-Vis absorption spectra of 30% CBS subfamily thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference.

143

Figure A10. UV-Vis absorption spectra of 40% CBS subfamily thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference.
144

Figure A11. UV-Vis absorption spectra of the T-4-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference.
145

Figure A12. UV-Vis absorption spectra of the T-6-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference.
146

Figure A13. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference.
147

Figure A14. UV-Vis absorption spectra of the T-10-T CBS subfamily thin films spin-cast
from o-dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference.

148

Table A3. Optical properties of neat polymers in thin films spin-cast from o-DCB and
placed in a N2 cabinet for 30 minutes.
Polymer λmax, abs
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
Eg
(nm / eV)
a

P3HTT-DPP-10% 680 0.449 826 / 1.50
10% T-4-T 654 0.474 803 / 1.54
10% T-6-T 652 0.461 796 / 1.56
10% T-8-T 656 0.530 799 / 1.55
10% T-10-T 652 0.440 809 / 1.53
20% T-4-T 670 0.769 787 / 1.58
20% T-6-T 658 0.780 784 / 1.58
20% T-8-T 668 0.616 771 / 1.61
20% T-10-T 650 0.490 775 / 1.60
30% T-4-T 672 0.649 762 / 1.63
30% T-6-T 664 0.878 761 / 1.65
30% T-8-T 682 0.636 749 / 1.66
30% T-10-T 674 0.572 754 / 1.64
40% T-4-T 676 0.659 729 / 1.70
40% T-6-T 672 0.870 748 / 1.68
40% T-8-T 680 0.601 739 / 1.68
40% T-10-T 678 0.783 736 / 1.68
a
Calculated from the absorption band edge in thin films, where Eg = 1240/λedge.

149

Table A4. Optical bandgaps of polymer family calculated from absorption band edge.
Eg (eV) T-4-T T-6-T T-8-T T-10-T
10% 1.54 1.56 1.55 1.53
20% 1.58 1.58 1.61 1.60
30% 1.63 1.65 1.66 1.64
40% 1.70 1.68 1.68 1.68
P3HTT-DPP-10% 1.50

Figure A15. Trends in optical bandgaps of polymer family calculated from absorption
band edge.

1.52
1.54
1.56
1.58
1.6
1.62
1.64
1.66
1.68
1.7
1.72
10% 20% 30% 40%
E
g
(eV)
Percent Incorporation of CBS
T-4-T
T-6-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-6-T)
Linear (T-8-T)
Linear (T-10-T)
150

Table A5. Solvent effects on optical bandgap for T-6-T subfamily of polymers.
Eg (eV)
T-6-T
CF
a
CB
b
CB
anneal
o-DCB
c
o-DCB
anneal

10% -
d
1.57 1.57 1.56 1.57
20% 1.59 1.59 1.63 1.61 1.63
30% 1.63 1.64 1.65 1.65 1.67
40% 1.66 1.65 1.65 1.68 1.71
a) Chloroform; b) Chlorobenzene; c) ortho-dichlorobenzene; d) Data not obtained

Figure A16. Trends in solvent effects on optical bandgap for T-6-T subfamily of polymers.

1.54
1.56
1.58
1.6
1.62
1.64
1.66
1.68
1.7
1.72
1.74
10% 20% 30% 40%
E
g
(eV)
Percent Incorporation of CBS
CF
CB
CB anneal
o-DCB
o-DCB anneal
Linear (CF)
Linear (CB)
Linear (CB anneal)
Linear (o-DCB)
Linear (o-DCB anneal)
151

A.6 Cyclic Voltammograms

Figure A17. Cyclic voltammograms of the 10% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference.

Figure A18. Cyclic voltammograms of the 20% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere.

152

Figure A19. Cyclic voltammograms of the 30% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere.

Figure A20. Cyclic voltammograms of the 40% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere.

153

Figure A21. Cyclic voltammograms of the T-4-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference.

Figure A22. Cyclic voltammograms of the T-6-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference.

154

Figure A23. Cyclic voltammograms of the T-8-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference.

Figure A24. Cyclic voltammograms of the T-10-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference.

155

Table A6. Highest occupied molecular orbitals of polymer family calculated from
oxidation onset versus ferrocene.
HOMO (eV) T-4-T T-6-T T-8-T T-10-T
10% 5.45 5.51 5.41 5.43
20% 5.41 5.41 5.49 5.44
30% 5.42 5.40 5.43 5.48
40% 5.43 5.41 5.45 5.50
P3HTT-DPP-10% 5.52

Figure A25. Trends in highest occupied molecular orbitals of polymer family calculated
from oxidation onset versus ferrocene.

5.38
5.4
5.42
5.44
5.46
5.48
5.5
5.52
10% 20% 30% 40%
HOMO (eV)
Percent Incorporation of CBS
T-4-T
T-6-T
T-8-T
T-10-T
156

A.7 Differential Scanning Calorimetry

Figure A26. DSC curve for 3.5 mg of P3HTT-DPP-10% using a scan rate of 10 °C/min.
The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions).

157

Figure A27. DSC curve for 5.0 mg of 30% T-8-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).

Figure A28. DSC curve for 4.2 mg of 30% T-10-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).
158

Figure A29. DSC curve for 4.3 mg of 40% T-6-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).

Figure A30. DSC curve for 4.3 mg of 40% T-8-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).
159

Figure A31. DSC curve for 3.8 mg of 40% T-10-T using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).

160

Table A7. Melting points for polymer family obtained from DSC curves.
Tm (°C) T-4-T T-6-T T-8-T T-10-T
10% -
a
-
a
-
a
-
a

20% -
a
-
a
-
a
-
a

30% -
a
-
a
90 75
40% -
a
143 125 83
P3HTT-DPP-10% 208
a) No thermal transitions observed

Figure A32. Trends in melting points for polymer family obtained from DSC curves.

50
70
90
110
130
150
170
4 6 8 10
T
m
(°C)
Spacer Length (# of carbons)
30%
40%
Linear (30%)
Linear (40%)
161

Table A8. Crystallization points for polymer family obtained from DSC curves.
Tc (°C) T-4-T T-6-T T-8-T T-10-T
10% -
a
-
a
-
a
-
a

20% -
a
-
a
-
a
-
a

30% -
a
-
a
81 73
40% -
a
112 85 76
P3HTT-DPP-10% 182
a) No thermal transitions observed

Figure A33. Trends in crystallization points for polymer family obtained from DSC curves.

50
70
90
110
130
150
170
4 6 8 10
T
c
(°C)
Spacer Length (# of carbons)
30%
40%
Linear (30%)
Linear (40%)
162

A.8 GIXRD Patterns
Thin films were spin-cast from o-DCB as described above and annealed for 30 minutes at
150 °C under N2. The thickness of films and GIXRD measurements were obtained using
Rigaku Diffractometer Ultima IV using Cu Kα radiation source (λ = 1.54 Å) in the
reflectivity and grazing-incidence X-ray diffraction mode, respectively.
Crystallite size was estimated using Scherrer’s equation:
𝜏 =
𝐾𝜆
𝛽 𝑐𝑜𝑠𝜃 (2)
where τ is the mean size of the ordered domains, K is the dimensionless shape factor (K
= 0.9), λ is the x-ray wavelength, β is the line broadening at half the maximum intensity
(FWHM) in radians, and θ is the Bragg angle.

163

Figure A34. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-coated from o-dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for
30 minutes.

164

Table A9. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation.
7,8

Polymer 2θ
(deg.)
d100
(Å)
Intensity
(counts)
FWHM
(deg.)
Crystallite
Size (nm)
P3HTT-DPP-10% 5.61 15.7 61 0.69 11.6
10% T-4-T 5.51 16.0 27 0.82 9.7
10% T-6-T 5.35 16.5 14 0.81 9.9
10% T-8-T 5.37 16.4 12 0.82 9.7
10% T-10-T 5.45 16.2 117 0.57 14.0
20% T-10-T 5.52 16.0 8 0.99 8.1
40% T-6-T 5.66 15.6 6 -
a
-
a

a) Unable to fit peak with Gaussian curve

165

A.9 Mobility Measurements
Mobility was measured using a hole-only device configuration of
ITO/PEDOT:PSS/Polymer/Al in the space charge limited current regime (SCLC).
9
The
dark current was measured under ambient conditions. At sufficient potential the mobilities
of charges in the device can be determined by fitting the dark current to the model of SCL
current and described by equation 3:
2
0 3
9
8
S C L C R
V
J
L
   
(3),
where JSCLC is the current density, ε0 is the permittivity of space, εR is the dielectric
constant of the polymer (assumed to be 3), μ is the zero-field mobility of the majority
charge carriers, V is the effective voltage across the device (V = Vapplied – Vbi – Vr), and L
is the polymer layer thickness. The series and contact resistance of the hole-only device
(18 – 23 Ω) was measured using a blank (ITO/PEDOT/Al) configuration and the voltage
drop due to this resistance (Vr) was subtracted from the applied voltage. The built-in
voltage (Vbi), which is based on the relative work function difference of the two electrodes,
was also subtracted from the applied voltage. The built-in voltage can be determined from
the transition between the ohmic region and the SCL region and is found to be about 1 V.
Polymer film thicknesses were measured using GIXRD in the reflectivity mode.

166

Table A10. Hole mobilities of eight polymers in thin films spin-cast from chloroform.
Results averaged over a minimum of 8 pixels.
Polymer μh, as-cast
a
(cm
2
V
-1
s
-1
)
μh, annealed
b

(cm
2
V
-1
s
-1
)
P3HTT-DPP-10% 1.73 ± 0.57 E-4 9.29 ± 0.53 E-4
10% T-6-T 2.54 ± 0.53 E-6 2.90 ± 0.87 E-6
10% T-8-T 3.37 ± 0.13 E-6 2.08 ± 0.64 E-5
10% T-10-T 1.49 ± 0.81 E-5 2.53 ± 1.02 E-4
20% T-8-T 1.75 ± 0.24 E-6 6.49 ± 2.34 E-6
20% T-10-T 2.51 ± 0.67 E-6 1.06 ± 0.29 E-5
30% T-10-T 5.66 ± 0.78 E-7 3.55 ± 1.94 E-6
40% T-10-T 9.04 ± 0.06 E-7 -
c

a) Films spin-cast from 7 mg mL
-1
solutions of polymers in chloroform and kept in a N2
box for 30 min; b) Films spin-cast from 7 mg mL
-1
solutions of polymers in chloroform
and annealed under N2 for 30 min at 150 °C; c) Not enough material to complete annealing
studies.

Table A11. Hole mobilities of eight polymers in thin films spin-cast from o-DCB. Results
averaged over a minimum of 8 pixels.
Polymer μh, as-cast
a
(cm
2
V
-1
s
-1
)
P3HTT-DPP-10% -
b

10% T-6-T 6.52 E-6
10% T-8-T 2.67 E-6
10% T-10-T 7.05 E-6
20% T-8-T 1.39 E-6
20% T-10-T 4.06 E-6
30% T-10-T 3.75 E-7
40% T-10-T 3.13 E-6
a) Films spin-cast from 7 mg mL
-1
solutions of polymers in o-DCB and kept in a N2 box
for 30 min; b) Not tested.

167

Table A12. Average hole mobilities for as-cast films.
μh E-5
(cm
2
V
-1
s
-1
)
T-4-T T-6-T T-8-T T-10-T
10% -
a
0.254 0.337 1.49
20% -
a
-
a
0.175 0.251
30% -
a
-
a
-
a
0.057
40% -
a
-
a
-
a
0.090
P3HTT-DPP-10% 17.3
a) Measurement unattainable due to poor film quality.

Figure A35. Trends in average hole mobilities for as-cast films.

0
0.2
0.4
0.6
0.8
1
1.2
1.4
1.6
10% 20% 30% 40%
μ
h
E-5 (cm
2
V
-1
s
-1
)
Percent Incorporation of CBS
T-6-T
T-8-T
T-10-T
Linear (T-8-T)
Linear (T-10-T)
168

Table A13. Average hole mobilities for annealed films.
μh E-5
(cm
2
V
-1
s
-1
)
T-4-T T-6-T T-8-T T-10-T
10% -
a
0.290 2.08 25.3
20% -
a
-
a
0.649 1.06
30% -
a
-
a
-
a
0.355
40% -
a
-
a
-
a
-
b

P3HTT-DPP-10% 92.9
a) Measurement unattainable due to poor film quality; b) Not enough material to complete
annealing studies.

Figure A36. Trends in average hole mobilities for annealed films.

0
5
10
15
20
25
30
10% 20% 30%
μ
h
E-5 (cm
2
V
-1
s
-1
)
Percent Incorporation of CBS
T-6-T
T-8-T
T-10-T
Linear (T-8-T)
Linear (T-10-T)
169

A.10 Film-on-Elastomer Mechanical Data
Table A14. Elastic moduli (film-on-elastomer) for polymer family, averaged over 3 trials.
Elastic Modulus
(GPa)
T-4-T T-6-T T-8-T T-10-T
10% 0.33 ± 0.11 0.15 ± 0.02 0.14 ± 0.06 0.15 ± 0.04
20% 1.30 ± 0.42 0.51 ± 0.16 0.65 ± 0.32 0.52 ± 0.12
30% 0.15 ± 0.09 -
a
0.59 ± 0.15 0.60 ± 0.07
40% -
a
-
a
-
a
0.60 ± 0.30
P3HTT-DPP-10% 0.32 ± 0.20
a) Measurement unattainable due to poor film quality.

Figure A37. Trends in elastic moduli for polymer family.

0
0.2
0.4
0.6
0.8
1
1.2
1.4
10% 20% 30% 40%
Elastic Modulus (GPa)
Percent Incorporation of CBS
T-4-T
T-6-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-6-T)
Linear (T-8-T)
Linear (T-10-T)
170

Table A15. Crack onset strain (COS) percentages for polymer family.
COS (%) T-4-T T-6-T T-8-T T-10-T
10% 27 > 80 > 80 > 80
20% 2 37 > 80 > 80
30% < 1 -
a
< 1 > 80
40% 1 -
a
1 77
P3HTT-DPP-10% 10
a) Measurement unattainable due to poor film quality

Figure A38. Trends in COS percentages for polymer family.

-20
0
20
40
60
80
100
4 6 8 10
Crack Onset Strain (%)
Length of Spacer (# of Carbons)
10%
20%
30%
40%
Linear (10%)
Linear (20%)
Linear (30%)
Linear (40%)
171

Table A16. Film thicknesses for film-on-elastomer measurements (500 rpm | 1000 rpm |
1500 rpm).
Thickness (nm) T-4-T T-6-T T-8-T T-10-T
10%
101 ± 4
77 ± 6
75 ± 6
290 ± 13
212 ± 1
183 ± 3
155 ± 5
126 ± 5
120 ± 6
256 ± 19
163 ± 11
168 ± 11
20%
203 ± 14
159 ± 8
133 ± 7
140 ± 6
94 ± 7
95 ± 6
203 ± 6
166 ± 6
139 ± 6
127 ± 6
113 ± 9
96 ± 5
30%
146 ± 14
130 ± 10
120 ± 6
140 ± 6
106 ± 9
79 ± 10
152 ± 3
107 ± 8
109 ± 4
154 ± 13
110 ± 2
101 ± 2
40% -
a
-
a
-
a

131 ± 27
95 ± 10
78 ± 11
P3HTT-DPP-10%
144 ± 7
100 ± 3
103 ± 9

a) Measurement unattainable due to poor film quality
172

Figure A39. Micrograph of P3HTT-DPP-10% film on PDMS at 24% elongation.
173

Figure A40. Micrograph of 10% T-4-T film on PDMS at 48% elongation.

174

Figure A41. Micrograph of 20% T-6-T film on PDMS at 52% elongation.

Table A17. Mode of failure obtained from micrographs of polymer films, where cracks
perpendicular to strain direction indicate brittle failure and pinholes that do not propagate
indicate ductile failure.
Mode of Failure T-4-T T-6-T T-8-T T-10-T
10% Ductile Ductile Ductile Ductile
20% Brittle Ductile Ductile Ductile
30% Brittle -
a
Brittle Ductile
40% Brittle -
a
Brittle Ductile
P3HTT-DPP-10% Brittle
a) Measurement unattainable due to poor film quality.

175

Table A18. Film thicknesses for film-on-water measurements (1000 rpm).
Thickness (nm) T-4-T T-6-T T-8-T T-10-T
10% -
a
167 ± 7 179 ± 11 117 ± 5
20% -
a
-
a
122 ± 7 145 ± 10
30% -
a
-
a
-
a
115 ± 6
40% -
a
-
a
-
a
93 ± 18
P3HTT-DPP-
10%
-
a

a) Not tested

176

A.11 References
1. Cardona, C. M.; Li, W.; Kaifer, A. E.; Stockdale, D.; Bazan, G. C. Electrochemical
Considerations for Determining Absolute Frontier Orbital Energy Levels of
Conjugated Polymers for Solar Cell Applications. Adv. Mater. 2011, 23, 2367-2371.
2. Thompson, B. C.; Kim, Y.; McCarley, T. D.; Reynolds, J. R. Soluble Narrow Band
Gap and Blue Propylenedioxythiophene-Cyanovinylene Polymers as Multifunctional
Materials for Photovoltaic and Electrochromic Applications. J. Am. Chem. Soc.
2006, 128, 12714-12725.
3. Khlyabich, P. P.; Burkhart, B.; Ng, C. F.; Thompson, B. C. Efficient Solar Cells from
Semi-Random P3HT Analogues Incorporating Diketopyrrolopyrrole.
Macromolecules 2011, 44, 5079-5084.
4. Zhao, X.; Zhao, Y.; Ge, Q.; Butrouna, K.; Diao, Y.; Graham, K. R.; Mei, J.
Complementary Semiconducting Polymer Blends: The Influence of Conjugation-
Break Spacer Length in Matrix Polymers. Macromolecules 2016, 49, 2601-2608.
5. Tahk, D.; Lee, H. H.; Khang, D. Elastic Moduli of Organic Electronic Materials by the
Buckling Method. Macromolecules 2009, 42, 7079-7083.
6. Alkhadra, M. A.; Root, S. E.; Hilby, K. M.; Rodriquez, D.; Sugiyama, F.; Lipomi, D. J.
Quantifying the Fracture Behavior of Brittle and Ductile Thin Films of
Semiconducting Polymers. Chem. Mater. 2017, 29, 10139–10149.
7. Ryu, C.; LeFevre, S.; Bao, Z.; Yang, H. Solubility-Driven Thin Film Structures of
Regioregular Poly(3-Hexyl Thiophene) using Volatile Solvents. Appl. Phys. Lett.
2007, 90, 172116.
8. Sharenko, A.; Treat, N. D.; Love, J. A.; Toney, M. F.; Stingelin, N.; Nguyen, T. Use of
a Commercially Available Nucleating Agent to Control the Morphological
Development of Solution-Processed Small Molecule Bulk Heterojunction Organic
Solar Cells. J. Mater. Chem. A 2014, 2, 15717-15721.
9. Kokil, A.; Yang, K.; Kumar, J. Techniques for Characterization of Charge Carrier
Mobility in Organic Semiconductors. J. Polym. Sci., Part B: Polym. Phys. 2012, 50,
1130-1144.

177

Appendix B: Influence of Acceptor Side-Chain Length on Mechanical and
Electronic Properties of Semi-Random Polymers with Conjugation-Break Spacers

B.1 Materials and Methods
All reagents from commercial sources were used without further purification, unless
otherwise noted. For mechanical measurements, PEDOT:PSS (Clevios PH 1000) was
purchased from Heraeus and used as received. For mechanical measurements, chloroform,
acetone, and isopropyl alcohol were purchased from Sigma–Aldrich Co. and used as
received. For synthesis and all other measurements, solvents were purchased from VWR
and used without purification with the exception of acetonitrile, which was distilled from
CaH2 prior to use. All reactions were performed under dry N2, unless otherwise noted. All
reactions were performed with glassware that was oven-dried and then flamed under high
vacuum and backfilled with N2. Flash chromatography was performed on a Teledyne
CombiFlash Rf instrument with RediSep Rf normal phase disposable columns.
1
H NMR
spectra were recorded in CDCl3 on a Varian Mercury 400 NMR spectrometer (small
molecules) or a Varian Mercury 600 NMR spectrometer (polymers). For molecular weight
determination, polymer samples were dissolved in HPLC grade o-dichlorobenzene (o-
DCB) at a concentration of 0.5 mg mL
-1
and filtered through a 0.2 μm PTFE filter. Size
exclusion chromatography (SEC) was performed using HPLC grade o-DCB at a flow rate
of 0.6 mL min
-1
on one 300 × 7.8 mm TSK-Gel GMHHR-H column (Tosoh Corp.) at 60
°C using a Viscotek GPC Max VE 2001 separation module and a Viscotek TDA 305 RI
detector. The instrument was calibrated with polystyrene standards (1,050 – 3,800,000
g/mol), and data were analyzed using OmniSec 4.6.0 software.
178

For thin film measurements, solutions were spin-cast onto pre-cleaned 2.5 cm
2
glass
slides (sonicated for 10 minutes in water, 5 minutes in acetone, and 5 minutes in isopropyl
alcohol then dried under high N2 flow) from 7 mg mL
-1
chloroform solutions. UV-Vis
absorption spectra were obtained on a PerkinElmer Lamda 950 spectrophotometer. The
thickness of the thin films and grazing-incidence X-ray diffraction (GIXRD) measurements
were obtained using a Rigaku Diffractometer Ultima IV using Cu Kα radiation source (λ =
1.54 Å) in the reflectivity and grazing-incidence mode, respectively. Differential scanning
calorimetry (DSC) traces were obtained using a Perkin-Elmer DSC 8000 with a scan rate
of 10 °C min
-1
. Sample size was ∼5 mg, and polymers were used as obtained after
purification.
Cyclic voltammetry (CV) was performed on Princeton Applied Research VersaStat3
potentiostat under the control of VersaStudio Software. A standard three-electrode cell
based on a platinum disc working electrode, a silver wire pseudo reference electrode
(calibrated vs Fc/Fc
+
which is taken as 5.1 eV vs vacuum),
1,2
and a Pt wire counter
electrode was purged with nitrogen and maintained under a nitrogen atmosphere during all
measurements. Polymer films were made by repeatedly dripping a 1% (w/w) chloroform
solution onto the Pt disc and drying under nitrogen prior to measurement.
Tetrabutylammonium hexafluorophosphate (0.1 M in freshly distilled acetonitrile) was
used as the supporting electrolyte.
B.1.1 SCLC Device Fabrication and Characterization
All steps of device fabrication and testing were performed in air. ITO-coated glass
substrates (10 Ω/m, Thin Film Devices Inc.) were sequentially cleaned by sonication in
detergent, de-ionized water, tetrachloroethylene, acetone, and isopropyl alcohol, and dried
179

in a nitrogen stream. A thin layer of PEDOT:PSS (Baytron® P VP AI 4083, filtered with
a 0.45 μm PVDF syringe filter – Pall Life Sciences) was first spin-coated on the pre-cleaned
ITO-coated glass substrates and baked at 120 ºC for 60 minutes under vacuum. Solutions
of polymers were prepared in chloroform at a concentration of 7 mg mL
-1
and stirred
overnight. Subsequently, the polymer active layer was spin coated (with a 0.45 μm PTFE
syringe filter – Pall Life Sciences) on top of the PEDOT:PSS layer. Upon spin coating of
polymers, films were first placed under N2 for 30 min or were annealed at 150 ºC for 30
minutes under N2 and then placed in the vacuum chamber for aluminum deposition. At the
final stage, the substrates were pumped down to high vacuum (< 2.5×10
-6
Torr) and
aluminum (100 nm) was thermally evaporated at 3 – 5 Å/sec using a Denton Benchtop
Turbo IV Coating System onto the active layer through shadow masks to define the active
area of the devices as 5.18 mm
2
. The current − voltage (I−V) characteristics of the devices
were measured under ambient conditions using a Keithley 2400 source-measurement unit.
B.1.2 Preparation of Substrates for Mechanical Measurements
Glass slides cut into 2.5 cm × 2.5 cm squares using a diamond–tipped scribe were used
as the substrate for the spin coating of the polymers. The glass slides were cleaned in
successive sonication baths of Alconox dissolved in deionized water, pure deionized water,
acetone, and isopropyl alcohol for 10 min each. Post sonication, the glass slides were dried
using a stream of compressed air. In order to improve wettability, activate the surface of
the glass, and remove any residual organic debris, the slides were treated with 30 W air
plasma for 5 min at a base pressure of 200–300 mTorr.

180

B.1.3 Preparation of Films for Mechanical Measurements
Solutions of pure polymers were mixed with chloroform in concentrations of 10 mg
mL
-1
and left to stir overnight. After mixing overnight, the solutions were filtered with 1
µm glass fiber media syringe filters. A layer of PEDOT:PSS spun at 1000 rpm (500 rpm s
-
1
ramp) for 3 min, followed by a second step at 2000 rpm (1000 rpm s
-1
ramp) for 30 s.
Filtered solutions were all spun at 1000 rpm (500 rpm s
-1
ramp) for 2 min, followed by
2000 rpm (1000 rpm s
-1
ramp) for 30 s. The thicknesses of the films were determined using
a Veeco Dektak stylus profilometer, in which at least 5 measurements were made per
sample.
B.1.4 Tensile Testing of Pseudo–Freestanding Films
The prepared sample is transferred to the surface of the water by partially submerging
the glass slide, which causes the layer of PEDOT:PSS to dissolve and the polymer to
delaminate from the surface. Once floating on the surface of the water, the film is brought
into contact with polydimethylsiloxane (PDMS) grips, prepared in the same methodology
as Alkhadra et al.,
3
attached to the load cell, and van der Waals adhesive forces keeps the
two in contact. The force versus displacement plots were obtained by uniaxially straining
the sample at 0.4 mm min
-1
while simultaneously recording the force trace until the sample
fractured. At least three tests were performed for each material. The stress–strain curves
were derived from the force versus displacement curves using the dimensions of the
corresponding sample.
B.2 Synthetic Procedures
The synthetic procedures for 2-bromo-5-trimethyltin-3-hexylthiophene (3HT) and 2,5-
didecyltetradecyl-3,6-bis(5-bromothiophene-2-yl)pyrrolo[3,4-c]-pyrrole-1,4-dione (DPP)
181

were previously published.
4,5
1,X-bis((5-trimethylstannyl)thiophen-2-yl)Xane alkyl
spacers (T-X-T, where X = the number of carbons separating the two thiophenes) were
synthesized according to the literature.
6
Typical polymerization (30% T-8-T/30% dtdDPP):
83 mg (0.20 mmol) of 3HT, 170 mg (0.15 mmol) DPP, and 90 mg (0.15 mmol) T-8-T were
dissolved in 12.5 mL of dry DMF to give a 0.04 M solution. The solution was then degassed
for 10 min before 25 mg (4 mol %) of Pd(PPh3)4 was added in one portion. The solution
was degassed for an additional 5 min and then heated for 48 h at 95 °C. Then the reaction
mixture was cooled to room temperature and precipitated into methanol. Purification was
achieved via Soxhlet extraction using methanol, hexanes, and chloroform. The chloroform
fraction was then reprecipitated into methanol, vacuum filtered, and dried. Yield 70% (153
mg), Mn = 47,600; Ð = 2.37.

182

B.3 Synthetic data
Table B1. Number-averaged polymer molecular weights in kDa as measured by SEC.
Mn (kDa) T-4-T T-8-T T-10-T
10% 29.9 26.1 33.0
20% - 35.5 31.1
30% - 47.6 -
40% 20.9 44.1 -
P3HTT-dtdDPP-10% 14.9

Figure B1. Trends in number-averaged polymer molecular weights as measured by SEC.

20
25
30
35
40
45
50
10% 15% 20% 25% 30% 35% 40%
M
n
(kDa)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
183

Table B2. Polymerization yields from chloroform Soxhlet fraction.
Yield (%) T-4-T T-8-T T-10-T
10% 48% 34% 38%
20% - 41% 44%
30% - 70% -
40% 69% 76% -
P3HTT-dtdDPP-10% 64%

Figure B2. Trends in polymerization yields from chloroform Soxhlet fraction.

20%
30%
40%
50%
60%
70%
80%
90%
10% 20% 30% 40%
Yield (%)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
184

B.4 Polymer Nuclear Magnetic Resonance Spectra

Figure B3.
1
H NMR spectrum of 40% T-4-T. The NMR was taken in CDCl3 at 50 °C on a
600 MHz instrument. This spectrum is being used as an example of how the amount of
monomer in each polymer was calculated. The aromatic peak at 8.86 ppm from the DPP
monomer is highlighted in blue; the aromatic peak at 6.75 ppm from the T-4-T monomer
is in orange; and the aromatic peak at 6.98 ppm from the 3HT monomer is in yellow. The
integrations of the DPP and T-4-T peaks match at 1.00 and 1.01, as they should. Each of
these peaks account for 2 hydrogens on these symmetrical monomers. The 3HT peak in
yellow accounts for only one hydrogen on the asymmetric monomer, therefore its
integration is doubled from 0.27 to 0.54. If these 3 integrations are summed, they give a
value of 2.55. Dividing each integration by this sum gives the percentage of each monomer
in the polymer chain. Therefore in this example, there is 39.2% DPP, 39.6% T-8-T, and
21.2% 3HT incorporated. This holds true to the feed ratio of 40% DPP, 40% T-8-T, and
20% 3HT.

185

Figure B4. Stacked
1
H NMR spectra of 10% T-4-T and 40% T-4-T. NMRs were taken in
CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from DPP monomer are
highlighted in blue; T-4-T in orange; and 3HT in yellow.

186

Figure B5. Stacked
1
H NMR spectra of 10% T-8-T, 20% T-8-T, 30% T-8-T, and 40% T-8-
T. NMRs were taken in CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from
DPP monomer are highlighted in blue; T-8-T in orange; and 3HT in yellow.

187

Figure B6. Stacked
1
H NMR spectra of 10% T-10-T and 20% T-10-T. NMRs were taken in
CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from DPP monomer are
highlighted in blue; T-10-T in orange; and 3HT in yellow.

188

B.5 Film UV-Vis Absorption Spectra

Figure B7. UV-Vis absorption spectra of 10% CBS subfamily thin films spin-cast from
chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10% absorption
spectrum provided for fully conjugated reference.

189

Figure B8. UV-Vis absorption spectra of the T-4-T CBS subfamily thin films spin-cast from
chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10% absorption
spectrum provided for fully conjugated reference.

190

Figure B9. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast from
chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10% absorption
spectrum provided for fully conjugated reference.

191

Figure B10. UV-Vis absorption spectra of the T-10-T CBS subfamily thin films spin-cast
from chloroform and placed in a N2 cabinet for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference.
192

Figure B11. UV-Vis absorption spectra of 10% CBS subfamily thin films spin-cast from
chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference.
193

Figure B12. UV-Vis absorption spectra of the T-4-T CBS subfamily thin films spin-cast
from chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference.
194

Figure B13. UV-Vis absorption spectra of the T-8-T CBS subfamily thin films spin-cast
from chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference.
195

Figure B14. UV-Vis absorption spectra of the T-10-T CBS subfamily thin films spin-cast
from chloroform and annealed under N2 at 150 °C for 30 minutes. P3HTT-dtdDPP-10%
absorption spectrum provided for fully conjugated reference.

196

Table B3. Optical properties of neat polymers in thin films spin-cast from chloroform and
placed in a N2 cabinet for 30 minutes.
Polymer
λ max, abs
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
E g
(nm / eV)
a

P3HTT-dtdDPP-10% 672 0.403 815 / 1.52
10% T-4-T 642 0.376 800 / 1.55
10% T-8-T 508 0.286 801 / 1.55
10% T-10-T 622 0.272 804 / 1.54
20% T-8-T 622 0.377 771 / 1.61
20% T-10-T 626 0.335 776 / 1.60
30% T-8-T 608 0.397 746 / 1.66
40% T-4-T 611 0.436 739 / 1.68
40% T-8-T 602 0.402 733 / 1.69
a
Calculated from the absorption band edge in thin films, where Eg = 1240/λedge.

Table B4. Optical properties of neat polymers in thin films spin-cast from chloroform and
annealed under N2 at 150 °C for 30 minutes.
Polymer
λ max, abs
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
E g
(nm / eV)
a

P3HTT-dtdDPP-10% 508 0.370 823 / 1.51
10% T-4-T 498 0.308 807 / 1.54
10% T-8-T 478 0.246 794 / 1.56
10% T-10-T 478 0.211 796 / 1.56
20% T-8-T 626 0.221 767 / 1.62
20% T-10-T 626 0.172 760 / 1.63
30% T-8-T 608 0.242 744 / 1.67
40% T-4-T 622 0.353 740 / 1.68
40% T-8-T 594 0.258 727 / 1.71
a
Calculated from the absorption band edge in thin films, where Eg = 1240/λedge.

197

Table B5. Optical bandgaps of as-cast polymer family calculated from absorption band
edge.
Eg, as cast (eV) T-4-T T-8-T T-10-T
10% 1.55 1.55 1.54
20% - 1.61 1.60
30% - 1.66 -
40% 1.68 1.69 -
P3HTT-dtdDPP-10% 1.52

Figure B15. Trends in optical bandgaps of as-cast polymer family calculated from
absorption band edge.

1.5
1.55
1.6
1.65
1.7
1.75
10% 20% 30% 40%
E
g
(eV)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
198

Table B6. Optical bandgaps of annealed polymer family calculated from absorption band
edge.
Eg, annealed (eV) T-4-T T-8-T T-10-T
10% 1.54 1.56 1.56
20% - 1.62 1.63
30% - 1.67 -
40% 1.68 1.71 -
P3HTT-dtdDPP-10% 1.51

Figure B16. Trends in optical bandgaps of annealed polymer family calculated from
absorption band edge.

1.5
1.55
1.6
1.65
1.7
1.75
1.8
10% 20% 30% 40%
E
g
(eV)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
199

B.6 Cyclic Voltammograms

Figure B17. Cyclic voltammograms of the 10% CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference.

200

Figure B18. Cyclic voltammograms of the T-4-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference.
201

Figure B19. Cyclic voltammograms of the T-8-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference.

202

Figure B20. Cyclic voltammograms of the T-10-T CBS subfamily of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-dtdDPP-10% cyclic voltammogram
provided for fully conjugated reference.

203

Table B7. Highest occupied molecular orbitals of polymer family calculated from
oxidation onset versus ferrocene.
HOMO (eV) T-4-T T-8-T T-10-T
10% 5.55 5.59 5.57
20% - 5.56 5.59
30% - 5.52 -
40% 5.57 5.54 -
P3HTT-dtdDPP-10% 5.62

Figure B21. Trends in highest occupied molecular orbitals of polymer family calculated
from oxidation onset versus ferrocene.

5.3
5.35
5.4
5.45
5.5
5.55
5.6
5.65
10% 20% 30% 40%
HOMO (eV)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
204

B.7 Differential Scanning Calorimetry
Figure B22. DSC curve for P3HTT-dtdDPP-10% using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).
205

Figure B23. DSC curve for 10% T-4-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

206

Figure B24. DSC curve for 40% T-4-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

Figure B25. DSC curve for 10% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

207

Figure B26. DSC curve for 20% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

Figure B27. DSC curve for 30% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

208

Figure B28. DSC curve for 40% T-8-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

Figure B29. DSC curve for 10% T-10-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

209

Figure B30. DSC curve for 20% T-10-T using a scan rate of 10 °C/min. The top curve is
heating (endothermic transitions) and the bottom curve is cooling (exothermic transitions).

210

Table B8. Melting points for polymer family obtained from DSC curves.
Tm (°C) T-4-T T-8-T T-10-T
10% 108 -
a
-
a

20% - 72 56
30% - 103 -
40% 164 120 -
P3HTT-dtdDPP-10% 150
a) No thermal transitions observed

Figure B31. Trends in melting points for polymer family obtained from DSC curves.

50
70
90
110
130
150
170
10% 20% 30% 40%
T
m
(°C)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
211

Table B9. Crystallization points for polymer family obtained from DSC curves.
Tc (°C) T-4-T T-8-T T-10-T
10% 93 -
a
-
a

20% - 50 33
30% - 64 -
40% 140 86 -
P3HTT-dtdDPP-10% 143
a) No thermal transitions observed

Figure B32. Trends in crystallization points for polymer family obtained from DSC curves.

30
50
70
90
110
130
150
10% 20% 30% 40%
T
c
(°C)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
212

B.8 GIXRD Patterns
For thin film measurements, solutions were spin-coated onto pre-cleaned 2.5 cm
2

glass slides (sonicated for 10 minutes in water, acetone, and isopropyl alcohol then dried
under high N2 flow) from chloroform solutions. The thickness of films and GIXRD
measurements were obtained using Rigaku Diffractometer Ultima IV using Cu Kα
radiation source (λ = 1.54 Å) in the reflectivity and grazing-incidence X-ray diffraction
mode, respectively.
Crystallite size was estimated using Scherrer’s equation:
τ = Kλ/(β cosθ) (1)
where τ is the mean size of the ordered domains, K is the dimensionless shape factor (K =
0.9), λ is the x-ray wavelength, β is the line broadening at half the maximum intensity
(FWHM) in radians, and θ is the Bragg angle.

213

Figure B33. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-cast from chloroform and placed in a N2 cabinet for 30 minutes.

Table B10. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
7,8
for as-
cast films.
Polymer
2θ
(deg.)
d 100
(Å)
Intensity
(counts)
FWHM
(deg.)
Crystallite
Size
(nm)
40% T-4-T/40% dtdDPP 4.33 20.4 34 0.84 9.48
20% T-8-T/20% dtdDPP 4.61 19.2 8 -
a
-
a

30% T-8-T/30% dtdDPP 4.77 18.5 10 -
a
-
a

40% T-8-T/40% dtdDPP 5.10 17.3 10 -
a
-
a

10% T-10-T/10% dtdDPP 4.61 19.2 9 1.28 6.23
20% T-10-T/20% dtdDPP 4.60 19.2 9 -
a
-
a

a) Unable to fit peak with Gaussian curve
214

Figure B34. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-cast from chloroform and annealed under N2 at 150 °C for 30 minutes.

Table B11. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
7,8
for
annealed films.
Polymer
2θ
(deg.)
d 100
(Å)
Intensity
(counts)
FWHM
(deg.)
Crystallite Size
(nm)
P3HTT-dtdDPP-10% 4.44 19.9 8 1.17 6.81
40% T-4-T/40% dtdDPP 4.40 20.1 268 0.57 13.98
40% T-8-T/40% dtdDPP 5.58 15.8 177 0.55 14.51
215

B.9 Mobility Measurements
Mobility was measured using a hole-only device configuration of
ITO/PEDOT:PSS/Polymer/Al in the space charge limited current regime (SCLC).
9
The
dark current was measured under ambient conditions. At sufficient potential the mobilities
of charges in the device can be determined by fitting the dark current to the model of SCL
current and described by equation 3:
2
0 3
9
8
S C L C R
V
J
L
   
(3),
where JSCLC is the current density, ε0 is the permittivity of space, εR is the dielectric constant
of the polymer (assumed to be 3), μ is the zero-field mobility of the majority charge
carriers, V is the effective voltage across the device (V = Vapplied – Vbi – Vr), and L is the
polymer layer thickness. The series and contact resistance of the hole-only device (18 – 23
Ω) was measured using a blank (ITO/PEDOT/Al) configuration and the voltage drop due
to this resistance (Vr) was subtracted from the applied voltage. The built-in voltage (Vbi),
which is based on the relative work function difference of the two electrodes, was also
subtracted from the applied voltage. The built-in voltage can be determined from the
transition between the ohmic region and the SCL region and is found to be about 1 V.
Polymer film thicknesses were measured using GIXRD in the reflectivity mode.

216

Table B12. Hole mobilities of eight CBS polymers and fully conjugated reference polymer
in thin films spin-cast from chloroform and left in a N2 box for 30 minutes. Results averaged
over at least 4 pixels.
μh E-5
(cm
2
V
-1
s
-1
)
T-4-T T-8-T T-10-T
10% 0.162 0.137 0.089
20% - 0.044 0.033
30% - 0.021 -
40% 0.041 0.017 -
P3HTT-dtdDPP-10% 1.45

Figure B35. Trends in as-cast hole mobilities of eight CBS polymers.

0
0.02
0.04
0.06
0.08
0.1
0.12
0.14
0.16
0.18
10% 20% 30% 40%
μ
h
E-5 (cm
2
V
-1
s
-1
)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
217

Table B13. Hole mobilities of eight CBS polymers and fully conjugated reference polymer
in thin films spin-cast from chloroform and annealed under N2 at 150 °C for 30 minutes.
Results averaged over at least 4 pixels.
μh E-5
(cm
2
V
-1
s
-1
)
T-4-T T-8-T T-10-T
10% 0.057 0.047 0.047
20% - 0.038 0.014
30% - 0.020 -
40% 0.021 0.013 -
P3HTT-dtdDPP-10% 2.97

Figure B36. Trends in annealed hole mobilities of eight CBS polymers.

0
0.01
0.02
0.03
0.04
0.05
0.06
10% 20% 30% 40%
μ
h
E-5 (cm
2
V
-1
s
-1
)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
218

B.10 Film-on-Water Mechanical Data
Table B14. Elastic moduli for polymer family.
Elastic Modulus (GPa) T-4-T T-8-T T-10-T
10% -
a
0.023 ± 0.003 0.016 ± 0.002
20% - 0.026 ± 0.002 0.006 ± 0.001
30% - 0.023 ± 0.004 -
40% 0.038 ± 0.001 0.036 ± 0.004 -
P3HTT-dtdDPP-10% 0.167 ± 0.016
a) Measurement unattainable due to film quality

Figure B37. Trends in elastic moduli for polymer family.

0
0.005
0.01
0.015
0.02
0.025
0.03
0.035
0.04
10% 20% 30% 40%
Elastic Modulus (GPa)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-4-T)
Linear (T-8-T)
Linear (T-10-T)
219

Table B15. Toughness values for polymer family.
Toughness (MPa) T-4-T T-8-T T-10-T
10% -
a
4.99 ± 0.26 4.13 ± 1.06
20% - 3.98 ± 1.20 4.15 ± 0.20
30% - 9.05 ± 2.04 -
40% 0.79 ± 0.06 20.59 ± 1.33 -
P3HTT-dtdDPP-10% 4.95 ± 1.00
a) Measurement unattainable due to film quality

Figure B38. Trends in toughness for polymer family.

0
5
10
15
20
25
10% 20% 30% 40%
Toughness (MPa)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-8-T)
Linear (T-10-T)
220

Table B16. UTS for polymer family.
UTS (MPa) T-4-T T-8-T T-10-T
10% -
a
3.61 ± 0.27 2.15 ± 0.50
20% - 2.63 ± 0.52 1.75 ± 0.33
30% - 4.49 ± 0.55 -
40% 1.77 ± 0.00 8.45 ± 0.11 -
P3HTT-dtdDPP-10% 8.68 ± 0.02
a) Measurement unattainable due to film quality

Figure B39. Trends in UTS for polymer family.

0
1
2
3
4
5
6
7
8
9
10% 20% 30% 40%
UTS (MPa)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-8-T)
Linear (T-10-T)
221

Table B17. Fracture strength for polymer family.
Fracture Strength (MPa) T-4-T T-8-T T-10-T
10% -
a
3.57 ± 0.32 1.90 ± 0.80
20% - 2.51 ± 0.55 1.74 ± 0.32
30% - 4.38 ± 0.67 -
40% 1.58 ± 0.05 8.45 ± 0.04 -
P3HTT-dtdDPP-10% 8.20 ± 0.55
a) Measurement unattainable due to film quality

Figure B40. Trends in fracture strength for polymer family.

0
1
2
3
4
5
6
7
8
9
10% 20% 30% 40%
Fracture Strength (MPa)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-8-T)
Linear (T-10-T)
222

Table B18. Fracture strain for polymer family.
Fracture Strain (%) T-4-T T-8-T T-10-T
10% -
a
185 ± 33 217 ± 6
20% - 190 ± 35 373 ± 67
30% - 323 ± 46 -
40% 53 ± 4 398 ± 32 -
P3HTT-dtdDPP-10% 63 ± 11
a) Measurement unattainable due to film quality

Figure B41. Trends in fracture strain for polymer family.

0
100
200
300
400
500
600
700
800
10% 20% 30% 40%
Fracture Strain (%)
Percent Incorporation of CBS
T-4-T
T-8-T
T-10-T
Linear (T-8-T)
Linear (T-10-T)
223

B.11 References
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Considerations for Determining Absolute Frontier Orbital Energy Levels of
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2. Thompson, B. C.; Kim, Y.; McCarley, T. D.; Reynolds, J. R. Soluble Narrow Band
Gap and Blue Propylenedioxythiophene-Cyanovinylene Polymers as Multifunctional
Materials for Photovoltaic and Electrochromic Applications. J. Am. Chem. Soc.
2006, 128, 12714-12725.
3. Alkhadra, M. A.; Root, S. E.; Hilby, K. M.; Rodriquez, D.; Sugiyama, F.; Lipomi, D. J.
Quantifying the Fracture Behavior of Brittle and Ductile Thin Films of
Semiconducting Polymers. Chem. Mater. 2017, 29, 10139–10149.
4. Khlyabich, P. P.; Burkhart, B.; Ng, C. F.; Thompson, B. C. Efficient Solar Cells from
Semi-Random P3HT Analogues Incorporating Diketopyrrolopyrrole.
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Content on Semi‐random Conjugated Polymers. J. Polym. Sci. , Part A: Polym.
Chem. 2017, 55, 3884-3892.
6. Zhao, X.; Zhao, Y.; Ge, Q.; Butrouna, K.; Diao, Y.; Graham, K. R.; Mei, J.
Complementary Semiconducting Polymer Blends: The Influence of Conjugation-
Break Spacer Length in Matrix Polymers. Macromolecules 2016, 49, 2601-2608.
7. Ryu, C.; LeFevre, S.; Bao, Z.; Yang, H. Solubility-Driven Thin Film Structures of
Regioregular Poly(3-Hexyl Thiophene) using Volatile Solvents. Appl. Phys. Lett.
2007, 90, 172116.
8. Sharenko, A.; Treat, N. D.; Love, J. A.; Toney, M. F.; Stingelin, N.; Nguyen, T. Use of
a Commercially Available Nucleating Agent to Control the Morphological
Development of Solution-Processed Small Molecule Bulk Heterojunction Organic
Solar Cells. J. Mater. Chem. A 2014, 2, 15717-15721.
9. Kokil, A.; Yang, K.; Kumar, J. Techniques for Characterization of Charge Carrier
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1130-1144.

224

Appendix C: Effects of Modulating Content of Conjugation-Break Spacers in Semi-
Random Polymers on Mechanical and Electronic Properties
C.1 Materials and Methods
All reagents from commercial sources were used without further purification, unless
otherwise noted. For mechanical measurements, PEDOT:PSS (Clevios PH 1000) was
purchased from Heraeus and used as received. For mechanical measurements, chloroform,
acetone, and isopropyl alcohol were purchased from Sigma–Aldrich Co. and used as
received. For synthesis and all other measurements, solvents were purchased from VWR
and used without purification with the exception of acetonitrile, which was distilled from
CaH2 prior to use. All reactions were performed under dry N2, unless otherwise noted. All
reactions were performed with glassware that was oven-dried and then flamed under high
vacuum and backfilled with N2. Flash chromatography was performed on a Teledyne
CombiFlash Rf instrument with RediSep Rf normal phase disposable columns.
1
H NMR
spectra were recorded in CDCl3 on a Varian Mercury 400 NMR spectrometer (small
molecules) or a Varian Mercury 600 NMR spectrometer (polymers). For molecular weight
determination, polymer samples were dissolved in HPLC grade o-dichlorobenzene (o-
DCB) at a concentration of 0.5 mg mL
-1
and filtered through a 0.2 μm PTFE filter. Size
exclusion chromatography (SEC) was performed using HPLC grade o-DCB at a flow rate
of 0.6 mL min
-1
on one 300 × 7.8 mm TSK-Gel GMHHR-H column (Tosoh Corp.) at 60
°C using a Viscotek GPC Max VE 2001 separation module and a Viscotek TDA 305 RI
detector. The instrument was calibrated with polystyrene standards (1,050 – 3,800,000
g/mol), and data were analyzed using OmniSec 4.6.0 software.
225

For thin film measurements, solutions were spin-cast onto pre-cleaned 2.5 cm
2
glass
slides (sonicated for 10 minutes in water, 5 minutes in acetone, and 5 minutes in isopropyl
alcohol then dried under high N2 flow) from 7 mg mL
-1
o-dichlorobenzene (o-DCB)
solutions. UV-Vis absorption spectra were obtained on a PerkinElmer Lamda 950
spectrophotometer. The thickness of the thin films and grazing-incidence X-ray diffraction
(GIXRD) measurements were obtained using a Rigaku Diffractometer Ultima IV using Cu
Kα radiation source (λ = 1.54 Å) in the reflectivity and grazing-incidence mode,
respectively. Differential scanning calorimetry (DSC) traces were obtained using a Perkin-
Elmer DSC 8000 with a scan rate of 10 °C min
-1
. Sample size was ∼5 mg, and polymers
were used as obtained after purification.
Cyclic voltammetry (CV) was performed on Princeton Applied Research VersaStat3
potentiostat under the control of VersaStudio Software. A standard three-electrode cell
based on a platinum disc working electrode, a silver wire pseudo reference electrode
(calibrated vs Fc/Fc
+
which is taken as 5.1 eV vs vacuum),
1,2
and a Pt wire counter electrode
was purged with nitrogen and maintained under a nitrogen atmosphere during all
measurements. Polymer films were made by repeatedly dripping a 1% (w/w) chloroform
solution onto the Pt disc and drying under nitrogen prior to measurement.
Tetrabutylammonium hexafluorophosphate (0.1 M in freshly distilled acetonitrile) was
used as the supporting electrolyte.
C.1.1 SCLC Device Fabrication and Characterization
All steps of device fabrication and testing were performed in air. ITO-coated glass
substrates (10 Ω/m, Thin Film Devices Inc.) were sequentially cleaned by sonication in
detergent, de-ionized water, tetrachloroethylene, acetone, and isopropyl alcohol, and dried
226

in a nitrogen stream. A thin layer of PEDOT:PSS (Baytron® P VP AI 4083, filtered with
a 0.45 μm PVDF syringe filter – Pall Life Sciences) was first spin-coated on the pre-cleaned
ITO-coated glass substrates and baked at 120 ºC for 60 minutes under vacuum. Solutions
of polymers were prepared in hot chloroform at a concentration of 7 mg mL
-1
and stirred
overnight. Subsequently, the polymer active layer was spin-cast from hot solvent (with a
0.45 μm PTFE syringe filter – Pall Life Sciences) on top of the PEDOT:PSS layer. Upon
spin coating of polymers, films were first placed under N2 for 30 min or were annealed at
150 ºC for 30 minutes under N2 and then placed in the vacuum chamber for aluminum
deposition. At the final stage, the substrates were pumped down to high vacuum (< 2.5×10
-
6
Torr) and aluminum (100 nm) was thermally evaporated at 3 – 5 Å/sec using a Denton
Benchtop Turbo IV Coating System onto the active layer through shadow masks to define
the active area of the devices as 5.18 mm
2
. The current − voltage (I−V) characteristics of
the devices were measured under ambient conditions using a Keithley 2400 source-
measurement unit.
C.1.2 Preparation of Substrates for Mechanical Measurements
Glass slides cut into 2.5 cm × 2.5 cm squares using a diamond–tipped scribe were used
as the substrate for the spin coating of the polymers. The glass slides were cleaned in
successive sonication baths of Alconox dissolved in deionized water, pure deionized water,
acetone, and isopropyl alcohol for 10 min each. Post sonication, the glass slides were dried
using a stream of compressed air. In order to improve wettability, activate the surface of
the glass, and remove any residual organic debris, the slides were treated with 30 W air
plasma for 5 min at a base pressure of 200–300 mTorr.

227

C.1.3 Preparation of Films for Mechanical Measurements
Solutions of pure polymers were mixed with chloroform in concentrations of 10 mg
mL
-1
and left to stir overnight. After mixing overnight, the solutions were filtered with 1
µm glass fiber media syringe filters. A layer of PEDOT:PSS spun at 1000 rpm (500 rpm s
-
1
ramp) for 3 min, followed by a second step at 2000 rpm (1000 rpm s
-1
ramp) for 30 s.
Filtered solutions were all spun at 1000 rpm (500 rpm s
-1
ramp) for 2 min, followed by
2000 rpm (1000 rpm s
-1
ramp) for 30 s. The thicknesses of the films were determined using
a Veeco Dektak stylus profilometer, in which at least 5 measurements were made per
sample.
C.1.4 Tensile Testing of Pseudo–Freestanding Films
The prepared sample is transferred to the surface of the water by partially submerging
the glass slide, which causes the layer of PEDOT:PSS to dissolve and the polymer to
delaminate from the surface. Once floating on the surface of the water, the film is brought
into contact with polydimethylsiloxane (PDMS) grips, prepared in the same methodology
as Alkhadra et al.,
5
attached to the load cell, and van der Waals adhesive forces keeps the
two in contact. The force versus displacement plots were obtained by uniaxially straining
the sample at 0.4 mm min
-1
while simultaneously recording the force trace until the sample
fractured. At least three tests were performed for each material. The stress–strain curves
were derived from the force versus displacement curves using the dimensions of the
corresponding sample.
C.2 Synthetic Procedures
The synthetic procedures for 2-bromo-5-trimethyltin-3-hexylthiophene (3HT) and 2,5-
diethylhexyl-3,6-bis(5-bromothiophene-2-yl)pyrrolo[3,4-c]-pyrrole-1,4-dione (DPP) were
228

previously published.
3
1,8-bis((5-trimethylstannyl)thiophen-2-yl)octane alkyl spacers
were synthesized according to the literature.
4

1,8-bis((5-bromo)thiophen-2-yl)octane (dibromo(T-8-T)): 0.947 g (3.40 mmol) 1,8-
(thiophen-2-yl)octane (synthesized according to the literature)
4
was dissolved in about 25
mL of chloroform, followed by 1.20 g (6.85 mmol) of N-bromosuccinimide. The reaction
was kept in the dark by wrapping with aluminum foil and stirred for 48 h at room
temperature. The solution was loaded directly onto a silica-packed cartridge and a column
was run in hexane. The eluent was recrystallized from hot ethanol, collected via vacuum
filtration, and dried. Yield 67% (1.002 g off-white crystals).
1
H NMR (600 MHz, CDCl3)
δ 6.84 (d, J = 3.7 Hz, 2H), 6.52 (dd, J = 3.6, 1.1 Hz, 2H), 2.73 (t, J = 7.6 Hz, 4H), 1.62 (p,
J = 7.6 Hz, 4H), 1.37 – 1.27 (m, 8H).
13
C NMR (151 MHz, CDCl3) δ 147.71, 129.54,
124.52, 108.73, 31.54, 30.45, 29.30, 29.03.
Typical polymerization (30% T-8-T/10% DPP): 260 mg (0.63 mmol) 3HT, 69 mg (0.10
mmol) DPP, 121 mg (0.20 mmol) distannyl(T-8-T), and 44 mg (0.10 mmol) dibromo(T-8-
T) were dissolved in 25 mL of dry DMF to give a 0.04 M solution. The solution was then
degassed for 10 min before 45 mg (4 mol %) of Pd(PPh3)4 was added in one portion. The
solution was degassed for an additional 5 min and then heated for 48 h at 95 °C. Then the
reaction mixture was cooled to room temperature and precipitated into methanol.
Purification was achieved via Soxhlet extraction using methanol, hexanes, and chloroform.
The chloroform fraction was then reprecipitated into methanol, vacuum filtered, and dried.
Yield 67% (160 mg), Mn = 12,800; Ð = 3.75.

229

C.3 Synthetic data
Table C1. Number-averaged polymer molecular weights in kDa as measured by SEC.

Mn (kDa)
10% T-8-T/10% DPP 19.7
20% T-8-T/10% DPP 8.5
30% T-8-T/10% DPP 12.8
40% T-8-T/10% DPP 12.4
50% T-8-T/10% DPP 10.2
P3HTT-DPP-10% 10.5

Figure C1. Trends in number-averaged polymer molecular weights as measured by SEC.

5
7
9
11
13
15
17
19
21
23
10% 15% 20% 25% 30% 35% 40% 45% 50%
M
n
(kDa)
Percent Incorporation of CBS
230

Table C2. Polymerization yields from chloroform Soxhlet fraction.

% Yield
10% T-8-T/10% DPP 78%
20% T-8-T/10% DPP 67%
30% T-8-T/10% DPP 67%
40% T-8-T/10% DPP 67%
50% T-8-T/10% DPP 48%
P3HTT-DPP-10% 68%

Figure C2. Trends in polymerization yields from chloroform Soxhlet fraction.

40%
45%
50%
55%
60%
65%
70%
75%
10% 20% 30% 40% 50%
Yield (%)
Percent Incorporation of CBS
T-8-T
Linear (T-8-T)
231

C.4 Small Molecule Nuclear Magnetic Resonance Spectra

Figure C3.
1
H NMR spectrum for 1,8-bis((5-bromo)thiophen-2-yl)octane taken in CDCl3
at 25 °C on a 600 MHz instrument.

232

Figure C4.
13
C NMR spectrum for 1,8-bis((5-bromo)thiophen-2-yl)octane taken in CDCl3
at 25 °C on a 600 MHz instrument.

233

C.5 Polymer Nuclear Magnetic Resonance Spectra

Figure C5.
1
H NMR spectrum of 20% T-8-T/10% DPP. The NMR was taken in CDCl3 at
50 °C on a 600 MHz instrument. This spectrum is being used as an example of how the
amount of monomer in each polymer was calculated. The aromatic peak at 8.88 ppm from
the DPP monomer is highlighted in blue; the aromatic peaks from 6.63-6.73 ppm from the
T-8-T monomer are in orange; and the aromatic peak at 6.96 ppm from the 3HT monomer
is in yellow. The integration of the DPP peak is 1.00 and that of the T-8-T peak is 2.10.
Each of these peaks account for 2 hydrogens on these symmetrical monomers. The 3HT
peak in yellow accounts for only one hydrogen on the asymmetric monomer, therefore its
integration is doubled from 3.53 to 7.06. If these 3 integrations are summed, they give a
value of 10.16. Dividing each integration by this sum gives the percentage of each
monomer in the polymer chain. Therefore in this example, there is 9.8% DPP, 20.7% T-8-
T, and 69.5% 3HT incorporated. This holds true to the feed ratio of 10% DPP, 20% T-8-
T, and 70% 3HT.

234

Figure C6. Stacked
1
H NMR spectra of 10% T-8-T/10% DPP, 20% T-8-T/10% DPP, 30%
T-8-T/10% DPP, 40% T-8-T/10% DPP, and 50% T-8-T/10% DPP. NMRs were taken in
CDCl3 at 50 °C on a 600 MHz instrument. Distinctive peaks from DPP monomer are
highlighted in blue; T-8-T in orange; and 3HT in yellow.

235

C.6 Film UV-Vis Absorption Spectra

Figure C7. UV-Vis absorption spectra of the T-8-T CBS family thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference.
Table C3. Optical properties of neat polymers in thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes.
Polymer
λ max, abs
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
λ max, abs
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
E g
(nm / eV)
a

P3HTT-DPP-10% 488 0.364 682 0.613 833 / 1.49
10% T-8-T/10% DPP 478 0.286 654 0.425 806 / 1.54
20% T-8-T/10% DPP 432 0.267 648 0.373 790 / 1.57
30% T-8-T/10% DPP 412 0.271 656 0.278 760 / 1.63
40% T-8-T/10% DPP 372 0.286 656 0.240 743 / 1.67
50% T-8-T/10% DPP 354 0.339 658 0.267 729 / 1.70
a
Calculated from the absorption band edge in thin films, where Eg = 1240/λedge.
236

Figure C8. UV-Vis absorption spectra of the T-8-T CBS family thin films spin-cast from o-
dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for 30 minutes. P3HTT-DPP-
10% absorption spectrum provided for fully conjugated reference.

Table C4. Optical properties of neat polymers in thin films spin-cast from o-
dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for 30 minutes.
Polymer
λ max, abs
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
λ max, abs
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
E g
(nm / eV)
a

P3HTT-DPP-10% 490 0.366 682 0.538 832 / 1.49
10% T-8-T/10% DPP 448 0.255 654 0.311 794 / 1.56
20% T-8-T/10% DPP 426 0.245 646 0.277 780 / 1.59
30% T-8-T/10% DPP 398 0.267 652 0.189 753 / 1.65
40% T-8-T/10% DPP 380 0.288 654 0.176 739 / 1.68
50% T-8-T/10% DPP 354 0.331 656 0.214 728 / 1.70
a
Calculated from the absorption band edge in thin films, where Eg = 1240/λedge.
237

Table C5. Optical bandgaps of unannealed polymer family calculated from absorption
band edge.

Eg (eV)
10% T-8-T/10% DPP 1.54
20% T-8-T/10% DPP 1.57
30% T-8-T/10% DPP 1.63
40% T-8-T/10% DPP 1.67
50% T-8-T/10% DPP 1.70
P3HTT-DPP-10% 1.49

Figure C9. Trends in optical bandgaps of unannealed polymer family calculated from
absorption band edge.

1.5
1.55
1.6
1.65
1.7
1.75
10% 20% 30% 40% 50%
E
g
(eV)
Percent Incorporation of CBS
238

Table C6. Optical bandgaps of annealed polymer family calculated from absorption band
edge.

Eg (eV)
10% T-8-T/10% DPP 1.56
20% T-8-T/10% DPP 1.59
30% T-8-T/10% DPP 1.65
40% T-8-T/10% DPP 1.68
50% T-8-T/10% DPP 1.70
P3HTT-DPP-10% 1.49

Figure C10. Trends in optical bandgaps of annealed polymer family calculated from
absorption band edge.

1.5
1.55
1.6
1.65
1.7
1.75
10% 20% 30% 40% 50%
E
g
(eV)
Percent Incorporation of CBS
239

C.7 Cyclic Voltammograms

Figure C11. Cyclic voltammograms of the T-8-T CBS family of films taken in dry
acetonitrile under nitrogen atmosphere. P3HTT-DPP-10% cyclic voltammogram provided
for fully conjugated reference.

240

Table C7. Highest occupied molecular orbitals of polymer family calculated from
oxidation onset versus ferrocene.

HOMO (eV)
10% T-8-T/10% DPP 5.52
20% T-8-T/10% DPP 5.48
30% T-8-T/10% DPP 5.50
40% T-8-T/10% DPP 5.63
50% T-8-T/10% DPP 5.48
P3HTT-DPP-10% 5.39

Figure C12. Trends in highest occupied molecular orbitals of polymer family calculated
from oxidation onset versus ferrocene.

5.3
5.35
5.4
5.45
5.5
5.55
5.6
5.65
10% 20% 30% 40% 50%
HOMO (eV)
Percent Incorporation of CBS
241

C.8 Differential Scanning Calorimetry

Figure C13. DSC curve for P3HTT-DPP-10% using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).

242

Figure C14. DSC curve for 5.6 mg of 10% T-8-T/10% DPP using a scan rate of 10 °C/min.
The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions).

Figure C15. DSC curve for 20% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).
243

Figure C16. DSC curve for 30% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).

Figure C17. DSC curve for 40% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).
244

Figure C18. DSC curve for 50% T-8-T/10% DPP using a scan rate of 10 °C/min. The top
curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).

245

Table C8. Melting points for polymer family obtained from DSC curves.

Tm (°C)
10% T-8-T/10% DPP -
a

20% T-8-T/10% DPP -
a

30% T-8-T/10% DPP -
a

40% T-8-T/10% DPP -
a

50% T-8-T/10% DPP 78
P3HTT-DPP-10% 211
a) No thermal transitions observed

Table C9. Crystallization points for polymer family obtained from DSC curves.

Tc (°C)
10% T-8-T/10% DPP -
a

20% T-8-T/10% DPP -
a

30% T-8-T/10% DPP -
a

40% T-8-T/10% DPP -
a

50% T-8-T/10% DPP 42
P3HTT-DPP-10% 207
a) No thermal transitions observed

246

C.9 GIXRD Patterns
For thin film measurements, solutions were spin-coated onto pre-cleaned 2.5 cm
2

glass slides (sonicated for 10 minutes in water, acetone, and isopropyl alcohol then dried
under high N2 flow) from o-dichlorobenzene solutions. The thickness of films and GIXRD
measurements were obtained using Rigaku Diffractometer Ultima IV using Cu Kα
radiation source (λ = 1.54 Å) in the reflectivity and grazing-incidence X-ray diffraction
mode, respectively.
Crystallite size was estimated using Scherrer’s equation:
τ = Kλ/(β cosθ) (1)
where τ is the mean size of the ordered domains, K is the dimensionless shape factor (K =
0.9), λ is the x-ray wavelength, β is the line broadening at half the maximum intensity
(FWHM) in radians, and θ is the Bragg angle.

247

Figure C19. Grazing-incidence X-ray diffraction patterns of thin films spin-cast from o-
dichlorobenzene (o-DCB) and placed in a N2 cabinet for 30 minutes. P3HTT-DPP-10%
absorption spectrum provided for fully conjugated reference.

Table C10. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
6,7
for as-
cast films.
Polymer
2θ
(deg.)
d 100
(Å)
Intensity
(counts)
FWHM
(deg.)
Crystallite Size
(nm)
P3HTT-DPP-10% 5.82 15.17 177 0.983 8.12
10% T-8-T/10% DPP 6.24 14.15 389 0.817 9.78
20% T-8-T/10% DPP 6.25 14.13 204 0.895 8.93
30% T-8-T/10% DPP 5.81 15.20 27 1.51 5.29

248

Figure C20. Grazing-incidence X-ray diffraction patterns of thin films of polymer thin
films spin-cast from o-dichlorobenzene (o-DCB) and annealed under N2 at 150 °C for 30
minutes. P3HTT-DPP-10% absorption spectrum provided for fully conjugated reference.

Table C11. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation
6,7
for
annealed films.
Polymer
2θ
(deg.)
d 100
(Å)
Intensity
(counts)
FWHM
(deg.)
Crystallite Size
(nm)
P3HTT-DPP-10% 5.77 15.30 250 0.851 9.38
10% T-8-T/10% DPP 5.66 15.60 47 1.06 7.53

249

C.10 Mobility Measurements
Mobility was measured using a hole-only device configuration of
ITO/PEDOT:PSS/Polymer/Al in the space charge limited current regime (SCLC).
8
The
dark current was measured under ambient conditions. At sufficient potential the mobilities
of charges in the device can be determined by fitting the dark current to the model of SCL
current and described by equation 3:
2
0 3
9
8
S C L C R
V
J
L
   
(3),
where JSCLC is the current density, ε0 is the permittivity of space, εR is the dielectric constant
of the polymer (assumed to be 3), μ is the zero-field mobility of the majority charge
carriers, V is the effective voltage across the device (V = Vapplied – Vbi – Vr), and L is the
polymer layer thickness. The series and contact resistance of the hole-only device (18 – 23
Ω) was measured using a blank (ITO/PEDOT/Al) configuration and the voltage drop due
to this resistance (Vr) was subtracted from the applied voltage. The built-in voltage (Vbi),
which is based on the relative work function difference of the two electrodes, was also
subtracted from the applied voltage. The built-in voltage can be determined from the
transition between the ohmic region and the SCL region and is found to be about 1 V.
Polymer film thicknesses were measured using GIXRD in the reflectivity mode.

250

Table C12. Hole mobilities of polymers in thin films spin-cast from hot chloroform and left
in a N2 box for 30 minutes. Results averaged over at least 4 pixels.
Polymer Average Hole Mobility, μ h, (cm
2
V
-1
s
-1
)
P3HTT-DPP-10% 1.02 E-4
10% T-8-T/10% DPP 3.02 E-6
20% T-8-T/10% DPP 2.20 E-6
30% T-8-T/10% DPP 2.02 E-7
40% T-8-T/10% DPP 3.63 E-8
50% T-8-T/10% DPP 6.82 E-9

Figure C21. Trends in hole mobility of polymer family in as-cast films.

0.0001
0.001
0.01
0.1
1
10% 20% 30% 40% 50%
μ
h
E-5 (cm
2
V
-1
s
-1
)
Percent Incorporation of CBS
251

Table C13. Hole mobilities of polymers in thin films spin-cast from hot chloroform and
annealed under N2 at 150 °C for 30 minutes. Results averaged over at least 4 pixels.
Polymer Average Hole Mobility, μ h, (cm
2
V
-1
s
-1
)
P3HTT-DPP-10% -
10% T-8-T/10% DPP 1.35 E-5
20% T-8-T/10% DPP 3.35 E-6
30% T-8-T/10% DPP 1.06 E-7
40% T-8-T/10% DPP 1.19 E-8
50% T-8-T/10% DPP 3.15 E-9

Figure C22. Trends in hole mobility of polymer family in annealed films.

0.0001
0.001
0.01
0.1
1
10
10% 20% 30% 40% 50%
μ
h
E-5 (cm
2
V
-1
s
-1
)
Percent Incorporation of CBS
252

C.11 Film-on-Water Mechanical Data
Table C14. Elastic moduli for polymer family.

Elastic Modulus (GPa)
10% T-8-T/10% DPP 0.140 ± 0.030
20% T-8-T/10% DPP 0.069 ± 0.006
30% T-8-T/10% DPP 0.009 ± 0.004
40% T-8-T/10% DPP 0.002 ± 0.000
50% T-8-T/10% DPP -
a

a) Measurement unattainable due to film quality

Figure C23. Trends in elastic moduli for polymer family.

0.001
0.01
0.1
1
10% 20% 30% 40%
Elastic Modulus (GPa)
Percent Incorporation of CBS
253

Table C15. Toughness for polymer family.

Toughness (MPa)
10% T-8-T/10% DPP 1.5 ± 0.5
20% T-8-T/10% DPP 23.65 ± 0.89
30% T-8-T/10% DPP 32.89 ± 4.43
40% T-8-T/10% DPP 14.99 ± 3.14
50% T-8-T/10% DPP -
a

a) Measurement unattainable due to film quality

Figure C24. Trends in toughness for polymer family.

0
5
10
15
20
25
30
35
10% 20% 30% 40%
Toughness (MPa)
Percent Incorporation of CBS
254

Table C16. Ultimate tensile strain (UTS) for polymer family.

UTS (MPa)
10% T-8-T/10% DPP 6.9 ± 0.9
20% T-8-T/10% DPP 18.99 ± 1.26
30% T-8-T/10% DPP 13.58 ± 1.92
40% T-8-T/10% DPP 13.07 ± 4.54
50% T-8-T/10% DPP -
a

a) Measurement unattainable due to film quality

Figure C25. Trends in UTS for polymer family.

6
8
10
12
14
16
18
20
10% 20% 30% 40%
UTS (MPa)
Percent Incorporation of CBS
255

Table C17. Fracture strength for polymer family.

Fracture Strength (MPa)
10% T-8-T/10% DPP 6.6 ± 0.5
20% T-8-T/10% DPP 18.73 ± 1.25
30% T-8-T/10% DPP 12.74 ± 0.62
40% T-8-T/10% DPP 9.97 ± 4.78
50% T-8-T/10% DPP -
a

a) Measurement unattainable due to film quality

Figure C26. Trends in fracture strength for polymer family.

0
2
4
6
8
10
12
14
16
18
20
10% 20% 30% 40%
Fracture Strength (MPa)
Percent Incorporation of CBS
256

Table C18. Fracture strain for polymer family.

Fracture Strain (%)
10% T-8-T/10% DPP 39 ± 9
20% T-8-T/10% DPP 200 ± 5
30% T-8-T/10% DPP 448 ± 29
40% T-8-T/10% DPP 293 ± 25
50% T-8-T/10% DPP -
a

a) Measurement unattainable due to film quality

Figure C27. Trends in fracture strain for polymer family.

0
50
100
150
200
250
300
350
400
450
500
10% 20% 30% 40%
Fracture Strain (%)
Percent Incorporation of CBS
257

C.12 References
1. Cardona, C. M.; Li, W.; Kaifer, A. E.; Stockdale, D.; Bazan, G. C. Electrochemical
Considerations for Determining Absolute Frontier Orbital Energy Levels of
Conjugated Polymers for Solar Cell Applications. Adv. Mater. 2011, 23, 2367-2371.
2. Thompson, B. C.; Kim, Y.; McCarley, T. D.; Reynolds, J. R. Soluble Narrow Band
Gap and Blue Propylenedioxythiophene-Cyanovinylene Polymers as Multifunctional
Materials for Photovoltaic and Electrochromic Applications. J. Am. Chem. Soc.
2006, 128, 12714-12725.
3. Khlyabich, P. P.; Burkhart, B.; Ng, C. F.; Thompson, B. C. Efficient Solar Cells from
Semi-Random P3HT Analogues Incorporating Diketopyrrolopyrrole.
Macromolecules 2011, 44, 5079-5084.
4. Zhao, X.; Zhao, Y.; Ge, Q.; Butrouna, K.; Diao, Y.; Graham, K. R.; Mei, J.
Complementary Semiconducting Polymer Blends: The Influence of Conjugation-
Break Spacer Length in Matrix Polymers. Macromolecules 2016, 49, 2601-2608.
5. Alkhadra, M. A.; Root, S. E.; Hilby, K. M.; Rodriquez, D.; Sugiyama, F.; Lipomi, D. J.
Quantifying the Fracture Behavior of Brittle and Ductile Thin Films of
Semiconducting Polymers. Chem. Mater. 2017, 29, 10139–10149.
6. Ryu, C.; LeFevre, S.; Bao, Z.; Yang, H. Solubility-Driven Thin Film Structures of
Regioregular Poly(3-Hexyl Thiophene) using Volatile Solvents. Appl. Phys. Lett.
2007, 90, 172116.
7. Sharenko, A.; Treat, N. D.; Love, J. A.; Toney, M. F.; Stingelin, N.; Nguyen, T. Use of
a Commercially Available Nucleating Agent to Control the Morphological
Development of Solution-Processed Small Molecule Bulk Heterojunction Organic
Solar Cells. J. Mater. Chem. A 2014, 2, 15717-15721.
8. Kokil, A.; Yang, K.; Kumar, J. Techniques for Characterization of Charge Carrier
Mobility in Organic Semiconductors. J. Polym. Sci. , Part B: Polym. Phys. 2012, 50,
1130-1144.

258

Appendix D: Diketopyrrolopyrrole (DPP) pendant polymers
D.1 Materials and Methods
All reagents from commercial sources were used without further purification, unless
otherwise noted. Solvents were purchased from VWR and used without purification with
the exception of acetonitrile, which was distilled from CaH2 prior to use. All reactions were
performed under dry N2, unless otherwise noted. All reactions were performed with
glassware that was oven-dried and then flamed under high vacuum and backfilled with N2.
Flash chromatography was performed on a Teledyne CombiFlash Rf instrument with
RediSep Rf normal phase disposable columns.
1
H NMR spectra were recorded in CDCl3
on a Varian Mercury 400 NMR spectrometer (small molecules) or a Varian Mercury 600
NMR spectrometer (polymers). For molecular weight determination, polymer samples
were dissolved in HPLC grade o-dichlorobenzene (o-DCB) at a concentration of 0.5
mg/mL and filtered through a 0.2 μm PTFE filter. Size exclusion chromatography (SEC)
was performed using HPLC grade o-DCB at a flow rate of 0.6 mL/min on one 300 × 7.8
mm TSK-Gel GMHHR-H column (Tosoh Corp.) at 60 °C using a Viscotek GPC Max VE
2001 separation module and a Viscotek TDA 305 RI detector. The instrument was
calibrated with polystyrene standards (1,050 – 3,800,000 g/mol), and data were analyzed
using OmniSec 4.6.0 software.
For thin film measurements, solutions were spin-cast onto pre-cleaned 2.5 cm
2
glass
slides (sonicated for 10 minutes in water, 5 minutes in acetone, and 5 minutes in isopropyl
alcohol then dried under high N2 flow) from 7 mg mL
-1
chloroform solutions. UV-Vis
absorption spectra were obtained on a PerkinElmer Lamda 950 spectrophotometer. The
steady-state emission at room temperature was measured with a Photon Technology
259

International QuantaMaster QM-400 spectrofluorometer. The thickness of the thin films
and grazing-incidence X-ray diffraction (GIXRD) measurements were obtained using a
Rigaku Diffractometer Ultima IV using Cu Kα radiation source (λ = 1.54 Å) in the
reflectivity and grazing-incidence mode, respectively. Differential scanning calorimetry
(DSC) traces were obtained using a Perkin-Elmer DSC 8000 with a scan rate of 10 °C/min.
Sample size was ∼5 mg, and polymers were used as obtained after purification. Cyclic
voltammetry (CV) was performed on Princeton Applied Research VersaStat3 potentiostat
under the control of VersaStudio Software. A standard three-electrode cell based on a
platinum disc working electrode, a silver wire pseudo reference electrode (calibrated vs
Fc/Fc
+
which is taken as 5.1 eV vs vacuum),
1,2
and a Pt wire counter electrode was purged
with nitrogen and maintained under a nitrogen atmosphere during all measurements.
Polymer films were made by repeatedly dripping a 1% (w/w) chloroform solution onto the
Pt disc and drying under nitrogen prior to measurement. Tetrabutylammonium
hexafluorophosphate (0.1 M in freshly distilled acetonitrile) was used as the supporting
electrolyte.
D.2 Synthesis
All reactions were performed on a Schlenk line under nitrogen gas in flame-dried glassware
unless otherwise noted. All reagents were used as purchased from commercial sources
unless otherwise noted. 4-methylstyrene (X) was passed through an alumina plug to
remove radical inhibitor and subsequently stored in the dark in a refrigerator. 2,2'-
azobisisobutyronitrile (AIBN) was recrystallized from methanol and stored in the dark in
a refrigerator. N-bromosuccinimide (NBS) was recrystallized from water and stored in the
dark in a refrigerator. Dry toluene was obtained from a solvent system and stored in a
260

Straus flask. Sodium azide and subsequent organic azides were handled with caution,
minimizing exposure to heat, friction, light, metals, and impact. Azide compounds were
weighed out using static-free plastic spatulas. Dri-solv dimethyl formamide (DMF) was
used. 2,5-diethylhexyl-3,6-bis(5-bromothiophene-2-yl)pyrrolo[3,4-c]-pyrrole-1,4-dione
(DPP) and decyltetradecyl (dtd) bromide were synthesized according to literature
procedures.
3,4
Potassium carbonate was oven-dried and stored in a desiccator. Dry
tetrahydrofuran (THF) was distilled from sodium/benzophenone. A 57% by weight
suspension of sodium hydride in oil was used.

Scheme D1. Synthesis of protected TIPS ethynylstyrene monomer (II), polymerization of
TIPS ethynylstyrene (III), and deprotection to atactic poly(ethynylstyrene) (IV)
(II) Protected alkyne monomer
To a 100 mL 3-neck flask was added 0.392 g (0.56 mmol) PdCl2(PPh3)2, 5.719 g (31.24
mmol) 4-bromostyrene (I), 12.874 g (70.59 mmol) triisopropylsilyl (TIPS) acetylene, and
41 mL triethylamine. The reaction was heated at 50 °C for 5 minutes, followed by the
addition of 0.149 g (0.78 mmol) copper iodide, at which point the reaction turned from
yellow to dark red. The reaction was stirred in the dark at 50 °C for 16 hours, after which
a grey salt was filtered off. The filtrate was loaded onto a silica cartridge and purified over
a hexane column. The clear liquid monomer was stored in the dark in a refrigerator. Yield:
7.608 g (86 %)
1
H NMR (400 MHz, Chloroform-d) δ 7.48 – 7.39 (m, 2H), 7.37 – 7.27 (m,
261

2H), 6.67 (ddd, J = 17.6, 14.1, 10.9 Hz, 1H), 5.75 (ddt, J = 17.4, 5.7, 0.7 Hz, 1H), 5.28 (dt,
J = 10.8, 1.0 Hz, 1H), 1.13 (m, 3H), 1.11 – 1.06 (m, 18H).
(III) aPS polymerization
To a 25 mL 3-neck flask was added 4.257 g (14.96 mmol) TIPS ethynylstyrene (II), 4 mg
(0.02 mmol) AIBN, and 3 mL dry toluene. The flask was degassed three times using the
freeze-pump-thaw technique, then heated to 70 °C and stirred for 10 hours. The solution
was precipitated into cold methanol and the precipitate was filtered and dried in a desiccator
overnight. Yield: 0.817 g (19 %); Mn: 13.3 kDa, Ð: 1.81
(IV) Deprotection
To a 100 mL 3-neck flask was added 0.786 g (2.76 mmol) poly(TIPS ethynylstyrene) (III)
and 40 mL of dry THF. After the polymer dissolved, the reaction was cooled to 0 °C and
the tetrabutylammonium fluoride solution was added slowly. The pale yellow solution
turned brown and darkened with the addition of TBAF. After the addition, the reaction was
allowed to warm to room temperature and was stirred overnight, after which it was
precipitated into cold methanol, filtered, and dried in a desiccator overnight. Yield: 0.311
g (88 %); Mn: 8.3 kDa, Ð: 1.84

Scheme D2. Synthesis of monoalkylated dtdDPP (VI), synthesis of bromobutyl/dtdDPP
(VII), conversion of bromine to azide (VIII)
(VI) DPP monoalkylation
To a 100 mL 3-neck flask was added 2.433 g (8.10 mmol) DPP (V), 2.812 g (20.25 mmol)
potassium carbonate, and 47 mL dry DMF. The reaction was heated to 120 °C and stirred
262

for 30 minutes (until the solution took on a blue hue) before the addition of 2.501 g (5.99
mmol) decyltetradecyl bromide. The reaction was stirred at 120 °C overnight, then
extracted with chloroform and the organic fraction washed with water and brine. The
organic fraction was dried over magnesium sulfate, filtered, and evaporated, then loaded
onto a silica cartridge. A column was run with a hexane:DCM gradient to remove the
disubstituted dtdDPP, then a hexane:ethyl acetate gradient was used to elute the
monosubstituted dtdDPP as the second peak. Yield: 1.176 g (31 %)
1
H NMR (400 MHz,
CDCl3) δ 8.78 (d, J = 2.7 Hz, 1H), 8.36 (d, J = 2.7 Hz, 1H), 7.63 (dd, J = 21.9, 3.9 Hz, 2H),
7.30 – 7.27 (m, 1H), 7.24 (d, J = 5.0 Hz, 1H), 4.02 (d, J = 7.7 Hz, 2H), 1.90 (m, 1H), 1.13-
1.33 (m, 40H), 0.83-0.93 (m, 6H).
(VII) DPP monoalkylation
To a 50 mL 3-neck flask was added 1.176 g (1.85 mmol) monoalkylated dtdDPP (VI) and
25 mL dry DMF. The solution was heated to 110 °C and 0.222 g (57 % by weight in oil,
5.27 mmol) sodium hydride was added gradually, turning the solution blue and releasing
hydrogen gas. When the bubbling stopped, 0.89 mL (7.45 mmol) 1,4-dibromobutane was
added and the reaction stirred overnight at 110 °C. The reaction extracted with chloroform
and the organic fraction washed with water and brine. The organic fraction was dried over
magnesium sulfate, filtered, and evaporated, then loaded onto a silica cartridge. The
compound was purified on a column with a gradient of hexane:DCM. Yield: 0.768 g (54
%)
1
H NMR (400 MHz, CDCl3) δ 8.94 (d, J = 3.9 Hz, 1H), 8.83 (d, J = 3.9 Hz, 1H), 7.64
(d, J = 5.0 Hz, 2H), 7.30 (d, J = 9.3 Hz, 2H), 4.38 (m, 2H), 4.02 (d, J = 7.7 Hz, 2H), 3.43
(t, 2H), 1.86 (m, 3H), 1.72 (m, 2H), 1.13-1.33 (m, 40H), 0.83-0.93 (m, 6H).
(VIII) Azide conversion
263

To a 50 mL 3-neck flask was added 0.768 g (1.00 mmol) bromoDPP (VII) and 20 mL dry
DMF. When the reaction mixture appeared homogenous, 0.669 g (10.29 mmol) sodium
azide was added. The reaction was stirred in the dark at room temperature for 4 days, then
extracted with ethyl acetate. The organic fraction was washed with water and brine, then
dried over magnesium sulfate, filtered and evaporated. Yield: 0.735 g (100 %)
1
H NMR
(400 MHz, CDCl3) δ 8.94 (d, J = 3.9 Hz, 1H), 8.83 (d, J = 3.9 Hz, 1H), 7.64 (d, J = 5.0 Hz,
2H), 7.30 (d, J = 9.3 Hz, 2H), 4.13 (m, 2H), 4.02 (d, J = 7.7 Hz, 2H), 3.35 (t, 2H), 1.86 (m,
3H), 1.72 (m, 2H), 1.13-1.33 (m, 40H), 0.83-0.93 (m, 6H).

Scheme D3. Copper-catalyzed azide-alkyne “click” reaction to join azidobutyl/dtdDPP
(VIII) with atactic poly(ethynylstyrene) (IV) to yield aPS-alkyne-click-DPP (IX)
(IX) Click reaction
To a 25 mL 3-neck flask was added 9 mg (0.03 mmol) copper sulfate pentahydrate
dissolved in 1 mL water, followed by 0.238 g (0.32 mmol) azide DPP (VIII), 34 mg (0.26
mmol) deprotected alkyne polymer (IV), and 8 mL DMF. The reagents were stirred until
it appeared the polymer had dissolved, then 19 mg (0.09 mmol) sodium ascorbate dissolved
in 1 mL water was injected. This was accompanied by an immediate darkening and
warming of the reaction mixture. The reaction was stirred in the dark at room temperature
264

for 1 day, then the solution was precipitated into cold methanol and the pink precipitate
was filtered and dried in a desiccator overnight. To remove excess small molecule azide, a
short column was run in chloroform. Yield: 55 mg (25 %); Mn: 10.9 kDa, Ð: 1.60

Scheme D4. Synthesis of atactic, syndiotactic, and isotactic poly(4-methylstyrene) (XI-a,
XI-s, and XI-i), bromination of atactic and syndiotactic poly(4-methylstyrene) (XII-a and
XII-s), and azide conversion to atactic and syndiotactic poly(4-azidomethylstyrene) (XIII-
a and XIII-s)
(XI-a) aPS
7.25 mL (55.03 mmol) of 4-methylstyrene (X), 20 mL of dry toluene, and 3.1 mg (0.04
mmol) AIBN were added to a 25 mL 3-neck flask. The system was evacuated using freeze-
pump-thaw methods three times before being lowered into a preheated oil bath. The
reaction mixture was stirred at 90 °C for 2 hours, then precipitated into cold methanol. The
resulting fluffy white precipitate was filtered and dried in a desiccator overnight. Yield:
0.998 g (15 %); Mn: 66.3 kDa, Ð: 1.89
(XI-s) sPS
Into a 50 mL 3-neck flask was added 10 mL of dry toluene and 0.38 mL (0.005 mmol) of
a 13 mM solution of CpTiCl3 in dry toluene. This was followed by 3.00 mL (2.99 mmol)
of MMAO-12 and 9.90 mL (75.14 mmol) of 4-methylstyrene (X). The reaction mixture
was stirred at room temperature for 2 days and the reaction was terminated with an injection
of acidified methanol. The mixture was precipitated into cold methanol and the resulting
fluffy white precipitate was filtered and dried in a desiccator overnight. Crude yield: 5.334
265

g (60 %); Mn: 36.3 kDa, Ð: 3.69 The solids were washed with methyl ethyl ketone in a
Soxhlet extractor for 24 hours, removing the atactic polymer and leaving syndiotactic
polymer behind in the thimble. Syndiotactic yield: 1.759 g (20 %); Mn: 36.9 kDa, Ð: 3.44
(XII) Bromination
To a 50 mL 3-neck flask equipped with a condenser was added 0.501 g (4.24 mmol) poly(4-
methylstyrene) (XI-a), 0.764 g (4.29 mmol) NBS, 0.037 g (0.22 mmol) AIBN, and 15 mL
of CCl4. The system was refluxed at 85 °C for 1 hour until orange coloration faded. The
reaction mixture was precipitated into cold methanol and the precipitate was filtered and
dried in a desiccator overnight. Yield: 0.682 g (82 %)
(XIII) Azide conversion
To a 50 mL 3-neck flask was added 0.497 g (2.52 mmol) of s-poly(4-bromomethylstyrene)
(XII-s), 1.574 g (24.21 mmol) sodium azide, and 20 mL dry DMF. The reaction was stirred
at 55 °C in the dark for 4 days, after which the mixture was precipitated into water. The
precipitate was filtered and dried in a desiccator overnight, and the crumbly light brown
solids were later stored in the dark in a freezer. Yield: 0.380 g (95 %)

Scheme D5. Synthesis of unsymmetrical propyne/dtdDPP (XIV)
(XIV) Unsymmetrical DPP alkylation
To a 100 mL 3-neck flask was added 1.095 g (3.64 mmol) DPP (V) and 25 mL of DMF.
The solution was heated to 110 °C and 0.486 g (57 % by weight in oil, 11.5 mmol) sodium
hydride was added gradually, turning the solution blue and releasing hydrogen gas. When
the bubbling stopped, after about 30 minutes, 2.146 g (5.14 mmol) decyltetradecyl bromide
266

was added. The reaction was stirred for another 30 minutes before the addition of 0.33 mL
(4.36 mmol) propargyl bromide, resulting in a red colored solution. After stirring
overnight, the solution was extracted with chloroform and the organic layer washed with
water and brine. The organic fraction was dried over magnesium sulfate, filtered,
evaporated, and loaded onto a silica cartridge. The compound was purified over a gradient
hexane:DCM column, eluting as the second peak (the first being disubstituted
decyltetradecylDPP). Yield: 0.540 g (22 %)
1
H NMR (400 MHz, CDCl3) δ 8.86 (d, J = 5.1
Hz, 1H), 8.74 (d, J = 5.0 Hz, 1H), 7.66 (d, J = 15.7 Hz, 2H), 7.30 (s, 2H), 4.85 (d, J = 2.5
Hz, 2H), 4.02 (d, J = 7.7 Hz, 2H), 2.31 (s, 1H), 1.90 (s, 1H), 1.10-1.35 (m, 40H), 0.80-0.93
(m, 6H).

Scheme D6. Copper-catalyzed azide-alkyne “click” reaction to join propyne/dtdDPP
(XIV) with atactic or syndiotactic poly(4-azidomethylstyrene) (XIII-a and XIII-s) to yield
aPS-azide-click-DPP and sPS-azide-click-DPP (XV-a and XV-s)
(XV) Click reaction
To a 50 mL 3-neck flask was added 10 mg (0.04 mmol) copper sulfate pentahydrate
dissolved in 1 mL water. This was followed by 0.051 g (0.31 mmol) a-poly(4-
azidomethylstyrene) (XIII-a), 0.292 g (0.43 mmol) propyne/decyltetradecylDPP (XIV), 8
267

mL DMF, and 13 mL toluene. The reagents were stirred until it appeared the polymer had
dissolved, then 84 mg (0.42 mmol) sodium ascorbate dissolved in 1 mL water was injected.
This was accompanied by an immediate darkening and warming of the reaction mixture.
The reaction was monitored by IR for the disappearance of the azide peak, in pursuit of
which two more aliquots of 76 mg and 60 mg of sodium ascorbate were added. After 3
days of stirring at room temperature, the solution was precipitated into cold methanol and
the resulting fluffy pink precipitate was filtered and dried in a desiccator overnight. Yield:
0.179 g (67 %)

Scheme D7. Synthesis of dtdDPP reference compound (XVI).

Scheme D8. Synthesis of "pre-clicked" styrene-based DPP pendant monomer (XVIII).

268

Scheme D9. Synthesis of acrylate-based DPP pendant monomer (XX).

269

D.3 Nuclear Magnetic Resonance Spectroscopy

Figure D1.
1
H NMR spectrum of poly(TIPS ethynylstyrene) (III) taken at 50 °C in CDCl3
on a 600 MHz instrument.

270

Figure D2.
13
C NMR spectrum of poly(TIPS ethynylstyrene) (III) taken at 50 °C in CDCl3
on a 600 MHz instrument.

271

Figure D3.
1
H NMR spectrum of poly(ethynylstyrene) (IV) taken at 50 °C in CDCl3 on a
600 MHz instrument.

272

Figure D4.
13
C NMR spectrum of poly(ethynylstyrene) (IV) taken at 50 °C in CDCl3 on a
600 MHz instrument.

273

Figure D5.
1
H NMR spectrum of monoalkylated dtdDPP (VI) taken at 25 °C in CDCl3 on
a 400 MHz instrument.

274

Figure D6.
1
H NMR spectrum of azidobutyl dtdDPP (VIII) taken at 25 °C in CDCl3 on a
400 MHz instrument.

275

Figure D7.
1
H NMR spectrum of clicked polymer (IX) taken at 50 °C in CDCl3 on a 600
MHz instrument.

276

Figure D8.
13
C NMR spectrum of clicked polymer (IX) taken at 50 °C in CDCl3 on a 600
MHz instrument.

277

Figure D9. Stacked
1
H NMR spectra of (III, VI, IX) taken in CDCl3 at 50 °C on a 600
MHz instrument. Distinctive peaks are highlighted from each successive reaction. The peak
highlighted in red in the alkyl region of the bottom spectrum is from the triisopropyl
protecting group and it disappears entirely in the next spectrum, being replaced by the
deprotected alkyne peak highlighted in green in the middle spectrum. That peak then
disappears entirely and new DPP resonances appear in the top spectrum, highlighted in
blue, indicating that the reaction proceeded to completion.

278

Figure D10. Stacked
13
C NMR spectra of (III, VI, IX) taken in CDCl3 at 50 °C on a 600
MHz instrument. Distinctive peaks are highlighted from each successive reaction. The
peaks highlighted in red on the bottom spectrum are from the triisopropyl protecting group
and it disappears entirely in the next spectrum, being replaced by the deprotected alkyne
peak highlighted in green in the middle spectrum. That peak then disappears entirely and
new DPP resonances appear in the top spectrum, highlighted in blue, indicating that the
reaction proceeded to completion.
279

Figure D11.
1
H NMR spectrum of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in CDCl3
on a 600 MHz instrument.

280

Figure D12.
13
C NMR spectrum of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in CDCl3
on a 600 MHz instrument.

Figure D13. Deconvolution of the aromatic peaks at ~143 ppm on the
13
C NMR spectrum
of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals 26.4% mm, 39.0% mr/rm, and 34.6% rr triads.
5-7

281

Figure D14. Deconvolution of the methine peaks at ~42-46 ppm on the
13
C NMR spectrum
of a-poly(4-methylstyrene) (XI-a) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals pentad peaks, but we were unable to fully assign the pentads.

Figure D15.
1
H NMR spectrum of s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3
on a 600 MHz instrument.

282

Figure D16.
13
C NMR spectrum of s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3
on a 600 MHz instrument.

Figure D17. Deconvolution of the aromatic peak at ~143 ppm on the
13
C NMR spectrum
of s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals only one peak, indicating rr triads > 95%.
8

283

Figure D18. Deconvolution of the methine peak at ~44 ppm on the
13
C NMR spectrum of
s-poly(4-methylstyrene) (XI-s) taken at 50 °C in CDCl3 on a 600 MHz instrument.
Deconvolution reveals only one peak, indicating rr triads > 95%.
8

Figure D19.
1
H NMR spectrum of a-poly(4-bromomethylstyrene) (XII-a) taken at 50 °C in
CDCl3 on a 600 MHz instrument.

284

Figure D20.
13
C NMR spectrum of a-poly(4-bromomethylstyrene) (XII-a) taken at 50 °C
in CDCl3 on a 600 MHz instrument.

285

Figure D21.
1
H NMR spectrum of s-poly(4-bromomethylstyrene) (XII-s) taken at 50 °C in
CDCl3 on a 600 MHz instrument.

286

Figure D22.
13
C NMR spectrum of s-poly(4-bromomethylstyrene) (XII-s) taken at 50 °C in
CDCl3 on a 600 MHz instrument.

287

Figure D23.
1
H NMR spectrum of a-poly(4-azidomethylstyrene) (XIII-a) taken at 55 °C in
DMSO-d6 on a 600 MHz instrument.

288

Figure D24.
13
C NMR spectrum of a-poly(4-azidomethylstyrene) (XIII-a) taken at 55 °C
in DMSO-d6 on a 600 MHz instrument.

289

Figure D25.
1
H NMR spectrum of s-poly(4-azidomethylstyrene) (XIII-s) taken at 55 °C in
toluene on a 600 MHz instrument.

290

Figure D26.
13
C NMR spectrum of s-poly(4-azidomethylstyrene) (XIII-s) taken at 25 °C in
tetrahydrofuran on a 600 MHz instrument.

291

Figure D27.
1
H NMR spectrum of unsymmetrical propyne/dtdDPP (XIV) taken at 25 °C
in CDCl3 on a 400 MHz instrument.

292

Figure D28.
1
H NMR spectrum of a-poly(styrene-click-DPP) (XV-a) taken at 50 °C in
CDCl3 on a 600 MHz instrument.

293

Figure D29.
13
C NMR spectrum of a-poly(styrene-click-DPP) (XV-a) taken at 50 °C in
CDCl3 on a 600 MHz instrument.

294

Figure D30.
1
H NMR spectrum of s-poly(styrene-click-DPP) (XV-s) taken at 50 °C in
CDCl3 on a 600 MHz instrument.

295

Figure D31.
13
C NMR spectrum of s-poly(styrene-click-DPP) (XV-s) taken at 50 °C in
CDCl3 on a 600 MHz instrument.

296

Figure D32. Stacked
1
H NMR spectra of a-poly(styrene)s (XI-a, XII-a, XIII-a, XV-a)
taken on a 600 MHz instrument. Distinctive peaks are highlighted from each successive
reaction. The peak highlighted in red in the alkyl region of the bottom spectrum is from the
para-methyl group on the styrene and it disappears entirely in the next spectrum, being
replaced by the methylene peak (adjacent to a bromine) highlighted in green in the second
spectrum from the bottom. The peak highlighted in green then shifts to the peak highlighted
in teal as the bromine is converted to an azide in the third spectrum from the bottom. That
peak then disappears entirely and new DPP resonances appear in the top spectrum,
highlighted in purple, indicating that the reaction proceeded to completion.
297

Figure D33. Stacked
13
C NMR spectra of a-poly(styrene)s (XI-a, XII-a, XIII-a, XV-a)
taken on a 600 MHz instrument. Distinctive peaks are highlighted from each successive
reaction. The peak highlighted in red in the alkyl region of the bottom spectrum is from the
para-methyl group on the styrene and it disappears entirely in the next spectrum, being
replaced by the methylene peak (adjacent to a bromine) highlighted in green in the second
spectrum from the bottom. The peak highlighted in green then shifts to the peak highlighted
in teal as the bromine is converted to an azide in the third spectrum from the bottom. That
peak then disappears entirely and new DPP resonances appear in the top spectrum,
highlighted in purple, indicating that the reaction proceeded to completion.

298

Figure D34. Stacked
1
H NMR spectra of s-poly(styrene)s (XI-s, XII-s, XIII-s, XV-s) taken
on a 600 MHz instrument. Distinctive peaks are highlighted from each successive reaction.
The peak highlighted in red in the alkyl region of the bottom spectrum is from the para-
methyl group on the styrene and it disappears entirely in the next spectrum, being replaced
by the methylene peak (adjacent to a bromine) highlighted in green in the second spectrum
from the bottom. The peak highlighted in green then shifts to the peak highlighted in teal
as the bromine is converted to an azide in the third spectrum from the bottom. That peak
then disappears entirely and new DPP resonances appear in the top spectrum, highlighted
in purple, indicating that the reaction proceeded to completion.

299

Figure D35. Stacked
13
C NMR spectra of s-poly(styrene)s (XI-s, XII-s, XIII-s, XV-s) taken
on a 600 MHz instrument. Distinctive peaks are highlighted from each successive reaction.
The peak highlighted in red in the alkyl region of the bottom spectrum is from the para-
methyl group on the styrene and it disappears entirely in the next spectrum, being replaced
by the methylene peak (adjacent to a bromine) highlighted in green in the second spectrum
from the bottom. The peak highlighted in green then shifts to the peak highlighted in teal
as the bromine is converted to an azide in the third spectrum from the bottom. That peak
then disappears entirely and new DPP resonances appear in the top spectrum, highlighted
in purple, indicating that the reaction proceeded to completion.

300

D.4 Infrared Spectroscopy

Figure D36. Click reaction monitored by IR spectroscopy for a-poly(styrene-azide) (XV-
a). The disappearance of the azide stretching peak at 2100 cm
-1
is indicative of the reaction
going to completion.

301

Figure D37. Click reaction monitored by IR spectroscopy for s-poly(styrene-azide) (XV-
s). The disappearance of the azide stretching peak at 2100 cm
-1
is indicative of the reaction
going to completion.

302

D.5 UV-Visible Spectroscopy

Figure D38. UV-Vis absorption spectra of clicked polymers in solution (chloroform).
Small molecule dtdDPP (XVI) absorption spectrum provided for reference.

303

Figure D39. UV-Vis absorption spectra of clicked polymer thin films spin-cast from
chloroform and placed in a N2 cabinet for 30 minutes. Small molecule dtdDPP (XVI)
absorption spectrum provided for reference.

304

Figure D40. Normalized dtdDPP (XVI) absorption spectra in solution (chloroform) and
film

305

Figure D41. Normalized aPS-alkyne-click-DPP (IX) absorption spectra in solution
(chloroform) and film

306

Figure D42. Normalized aPS-azide-click-DPP (XV-a) absorption spectra in solution
(chloroform) and film

307

Figure D43. Normalized sPS-azide-click-DPP (XV-s) absorption spectra in solution
(chloroform) and film

Table D1. Absorption properties of neat polymers in thin films spin-cast from chloroform
and placed in a N2 cabinet for 30 minutes.
Polymer
λmax, film
(nm)
λmax, solution
(nm)
Absorption
Coefficient
(cm
-1
x10
-5
)
Eg, film
(nm / eV)
a

Eg, solution
(nm / eV)
a

dtdDPP (XVI) 512 548 0.496 634 / 1.99 572 / 2.17
aPS-alkyne-click-
DPP (IX)
560 549 0.383 596 / 2.08 579 / 2.12
aPS-azide-click-
DPP (XV-a)
558 548 0.323 592 / 2.09 584 / 2.12
sPS-azide-click-
DPP (XV-s)
558 516 0.262 591 / 2.10 588 / 2.11
a
Calculated from the absorption band edge in thin films, where Eg = 1240/λedge.

308

D.6 Fluorescence Spectroscopy

Figure D44. Normalized solution emission spectra in CHCl3. Small molecule dtdDPP
(XVI) emission spectrum provided for reference.

309

Figure D45. Normalized dtdDPP (XVI) absorption and emission spectra in CHCl3

310

Figure D46. Normalized aPS-alkyne-click-DPP (IX) absorption and emission spectra in
CHCl3

311

Figure D47. Normalized aPS-azide-click-DPP (XV-a) absorption and emission spectra in
CHCl3

312

Figure D48. Normalized sPS-azide-click-DPP (XV-s) absorption and emission spectra in
CHCl3

Table D2. Emission properties of neat polymers in thin films spin-cast from chloroform
and placed in a N2 cabinet for 30 minutes.
Polymer
λmax, solution
(nm)
λonset,solution
(nm / eV)
a

λintersect,solution
(nm / eV)
a

Φsolution
(%)
b

Φfilm
(%)
b

dtdDPP (XVI) 564 543 / 2.28 557 / 2.23 64 % 4 %
aPS-alkyne-click-
DPP (IX)
608 538 / 2.30 557 / 2.23 4 % 0 %
aPS-azide-click-
DPP (XV-a)
560 538 / 2.30 555 / 2.23 2 % 0 %
sPS-azide-click-
DPP (XV-s)
560 537 / 2.31 555 / 2.23 2 % 0 %
a
Calculated from the absorption band edge in thin films, where Eg = 1240/λedge.

313

D.7 Differential Scanning Calorimetry (DSC)

Figure D49. DSC curve for 7.1 mg of dtdDPP (XVI) using a scan rate of 10 °C/min. The
top curve is heating (endothermic transitions) and the bottom curve is cooling (exothermic
transitions).

314

Figure D50. DSC curve for 5.3 mg of aPS-alkyne-click-DPP (IX) using a scan rate of 10
°C/min. The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions).

315

Figure D51. DSC curve for 7.5 mg of aPS-azide-click-DPP (XV-a) using a scan rate of 10
°C/min. The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions).

316

Figure D52. DSC curve for 7.2 mg of sPS-azide-click-DPP (XV-s) using a scan rate of 10
°C/min. The top curve is heating (endothermic transitions) and the bottom curve is cooling
(exothermic transitions).

317

D.8 X-ray Diffraction
Thin films were spin-cast from chloroform as described above and placed in a N2 cabinet
for 30 minutes. The thickness of films and GIXRD measurements were obtained using
Rigaku Diffractometer Ultima IV using Cu Kα radiation source (λ = 1.54 Å) in the
reflectivity and grazing-incidence X-ray diffraction mode, respectively.
Crystallite size was estimated using Scherrer’s equation:
𝜏 =
𝐾𝜆
𝛽 𝑐𝑜𝑠𝜃 (1)
where τ is the mean size of the ordered domains, K is the dimensionless shape factor (K =
0.9), λ is the x-ray wavelength, β is the line broadening at half the maximum intensity
(FWHM) in radians, and θ is the Bragg angle.

Figure D53. Grazing-incidence X-ray diffraction pattern of thin film of dtdDPP (XVI)
spin-cast from chloroform and placed in a N2 cabinet for 30 minutes.
318

Table D3. 2θ, interchain distances (100), GIXRD intensities, full-width at half maximum
(FWHM) values, and crystallite size (nm) calculated from Scherrer’s equation.
9,10

Polymer
2θ
(deg.)
d100
(Å)
Intensity
(counts)
FWHM
(deg.)
Crystallite
Size
(nm)
dtdDPP (XVI) 2.53 34.9 591 0.35 22.8
a) Unable to fit peak with Gaussian curve
319

D.9 Cyclic Voltammetry

Figure D54. Cyclic voltammograms of the clicked polymer films taken in dry acetonitrile
under nitrogen atmosphere. The small molecule dtdDPP (XVI) cyclic voltammogram is
provided for reference.
320

D.10 References
1. Cardona, C. M.; Li, W.; Kaifer, A. E.; Stockdale, D.; Bazan, G. C. Electrochemical
Considerations for Determining Absolute Frontier Orbital Energy Levels of Conjugated
Polymers for Solar Cell Applications. Adv. Mater. 2011, 23, 2367-2371.
2. Thompson, B. C.; Kim, Y.; McCarley, T. D.; Reynolds, J. R. Soluble Narrow Band
Gap and Blue Propylenedioxythiophene-Cyanovinylene Polymers as Multifunctional
Materials for Photovoltaic and Electrochromic Applications. J. Am. Chem. Soc. 2006,
128, 12714-12725.
3. Khlyabich, P. P.; Burkhart, B.; Ng, C. F.; Thompson, B. C. Efficient Solar Cells from
Semi-Random P3HT Analogues Incorporating Diketopyrrolopyrrole. Macromolecules
2011, 44, 5079-5084.
4. Ekiz, S.; Gobalasingham, N. S.; Thompson, B. C. Exploring the Influence of Acceptor
Content on Semi‐random Conjugated Polymers. J. Polym. Sci. , Part A: Polym. Chem.
2017, 55, 3884-3892.
5. Dammert, R. M.; Maunu, S. L.; Maurer, F. H. J.; Neelov, I. M.; Niemelä, S.;
Sundholm, F.; Wästlund, C. Free Volume and Tacticity in Polystyrenes. Macromolecules
1999, 32, 1930-1938.
6. Ziaee, F.; Khoshkhoo, M. Effect of Temperature on Tacticity in Thermal
Polymerization of P-Methylstyrene by 13C NMR Spectroscopy. Iran. Polym. J. 2012, 21,
21-29.
7. Kawamura, T.; Uryu, T.; Matsuzaki, K. Poly(Methy1styrene)s and
Poly(Methoxystyrene)s obtained with Ziegler Catalyst, Cationic Catalyst, Or Radical
Initiators. Makromol. Chem. 1982, 183, 125-141.
8. Grassi, A.; Longo, P.; Proto, A.; Zambelli, A. Reactivity of some Substituted Styrenes
in the Presence of a Syndiotactic Specific Polymerization Catalyst. Macromolecules
1989, 22, 104-108.
9. Ryu, C.; LeFevre, S.; Bao, Z.; Yang, H. Solubility-Driven Thin Film Structures of
Regioregular Poly(3-Hexyl Thiophene) using Volatile Solvents. Appl. Phys. Lett. 2007,
90, 172116.
10. Sharenko, A.; Treat, N. D.; Love, J. A.; Toney, M. F.; Stingelin, N.; Nguyen, T. Use
of a Commercially Available Nucleating Agent to Control the Morphological
Development of Solution-Processed Small Molecule Bulk Heterojunction Organic Solar
Cells. J. Mater. Chem. A 2014, 2, 15717-15721.

321

Biographical Sketch
Betsy was born and raised in Richmond, Virginia, where she
attended Maggie L. Walker Governor’s School for Government
and International Studies. After graduating in 2007, she took a
gap year during which she hiked the Appalachian Trail and
volunteered on a biodynamic farm and taught in a one-room
schoolhouse in Costa Rica. Betsy then attended the University
of North Carolina at Chapel Hill, where she majored in Chemistry and minored in Hispanic
Studies and Linguistics. While at UNC, she researched encapsulation methods for organic
photovoltaics in the lab of Professor Wei You. After earning her B.S. in 2011, Betsy moved
to New Jersey to work at the US Environmental Protection Agency as an Oak Ridge
Institute for Science and Education Fellow. She left this position in 2013 to pursue her PhD
in Chemistry at the University of Southern California.

Asset Metadata

Creator Melenbrink, Elizabeth L. (Betsy) (author)

Core Title Non-traditional architectures for semiconducting polymers

Contributor Electronically uploaded by the author (provenance)

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Publication Date 10/30/2018

Defense Date 08/24/2018

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag conjugation-break spacers,electroactive pendant polymers,mechanical properties,OAI-PMH Harvest,organic electronics,semiconducting polymers

Format application/pdf (imt)

Language English

Advisor Thompson, Barry (committee chair), Thompson, Mark (committee member), Yoon, Jongseung (committee member)

Creator Email b.melenbrink@verizon.net,melenbri@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c89-99585

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Document Type Dissertation

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Rights Melenbrink, Elizabeth L. (Betsy)

Type texts

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Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

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Abstract (if available)

Abstract Electronic devices are becoming progressively smaller and more integrated into daily life. Organic semiconductors offer the possibility of making these devices lightweight, flexible, and easy to manufacture. ❧ Conjugated polymers are widely explored for a variety of organic electronic devices. However, there are some drawbacks to using the traditional linear, fully conjugated polymer. This dissertation explores some alternative architectures for semiconducting polymers in an attempt to overcome some of the challenges facing traditional architectures. ❧ An overview of alternative architectures for semiconducting polymers is presented in Chapter 1. Special attention is paid to linear conjugated polymers with non-conjugated segments interspersed, so-called conjugation-break spacers (CBSs), as well as non-conjugated polymers with electroactive pendants. The history, advantages, disadvantages, and recent trends for these systems are discussed. ❧ In Chapter 2 conjugation-break spacers are applied to the semi-random polymer design developed previously in the Barry C. Thompson research group. The effects of spacer length and content as it varies concurrently with the other monomers in the polymer system are studied. Charge mobility is observed to decrease with increasing CBS incorporation, but increase with increasing CBS length. This is accompanied by an improvement in mechanical properties as measured by film-on-elastomer (FOE) and film-on-water (FOW) techniques. This study results in CBS polymers with comparable electronic properties but greatly improved mechanical properties relative to their fully conjugated analogs. ❧ Some polymers within the family synthesized for Chapter 2 could not be analyzed due to poor solubility. Consequently, a subset of these polymers with longer solubilizing side-chains are synthesized and analyzed in Chapter 3. Solubility and mechanical properties are seen to improve greatly, with some polymers exhibiting properties approaching elastomeric limits. Free-standing films measured by the FOW technique could be stretched to many times their original length before breaking. However, this improvement in mechanical properties is accompanied by a diminishment in electronic properties. ❧ Within the series of CBS polymers synthesized for Chapter 2, the content of all three monomers was changed simultaneously, making it difficult to isolate the effects of the CBS content on electronic and mechanical properties. Therefore, in Chapter 4 a small set of CBS polymers are synthesized and evaluated in which the content of one monomer is held constant and the other two are varied. Similar to Chapter 3, an improvement in the mechanical performance of these polymers is observed along with a decrease in desirable electronic properties. ❧ Finally, in Chapter 5 a set of novel non-conjugated polystyrene-based polymers with electroactive diketopyrrolopyrrole (DPP) pendants is synthesized and characterized. The three polymers differ in tacticity and spacer length linking the DPP pendant to the polystyrene backbone. There is no observed effect of tacticity on optical properties, but a possible influence of spacer length on optical properties.