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Organic Blue Light-Emitting Diodes and Field-Effect Transistors
Based on Monodisperse Conjugated Oligomers
by
Sean W. Culligan
Submitted in Partial Fulfillment
of the
Requirements for the Degree
Doctor of Philosophy
Supervised by
Professor Shaw H. Chen and Professor Ching W. Tang
Department of Chemical Engineering
The College
School of Engineering and Applied Sciences
University of Rochester
Rochester, New York
2006
ii
To My Wonderful Wife and Family
iii
CURRICULUM VITAE
Sean W. Culligan was born in 1978 in Poughkeepsie, New York. In 2000 he
received a Bachelors of Science degree in Chemical Engineering from the University
of Rochester. He continued his education in the Department of Chemical
Engineering at the University of Rochester and received his Masters of Science
degree in 2003. He then pursued his doctorate in Chemical Engineering under the
supervision of Professor Shaw H. Chen and Professor Ching W. Tang. His research
was in the field of organic electronic materials and devices.
PUBLICATIONS TO DATE IN REFEREED JOURNALS
1. Fan, F. Y., Culligan, S. W., Mastrangelo, J. C., Katsis, D., Chen, S. H. Chem. Mater. 13, 4584 (2001).
2. Geng, Y., Katsis, D., Culligan, S. W., Ou, J. J., Chen, S. H., Rothberg, L. J.
Chem. Mater. 14, 463 (2002). 3. Katsis, D., Geng, Y., Ou, J. J., Culligan, S. W., Trajkovska, A., Chen, S. H.,
Rothberg, L. J. Chem. Mater. 14, 1332 (2002). 4. Geng, Y., Culligan, S. W., Trajkovska, A., Wallace, J. U., Chen. S. H. Chem.
Mater. 15, 542 (2003). 5. Culligan, S. W., Geng, Y., Chen, S. H., Klubek, K. P., Vaeth, K. M., Tang, C.
W. Adv. Mater. 15, 1176 (2003). 6. Geng, Y., Trajkovska, A., Culligan, S. W., Ou, J. J., Chen, H. M. P., Katsis, D.,
Chen, S. H. J. Am. Chem. Soc. 125, 14032 (2003). 7. Chen, A. C.-A., Culligan, S. W., Geng, Y., Chen, S. H., Klubek, K. P., Vaeth, K.
M., Tang, C. W. Adv. Mater. 16, 783 (2004) 8. Culligan, S. W., Chen, A. C.-A., Klubek, K. P., Tang, C. W., Chen, S. H. Adv.
Funct. Mater. (2006, accepted). 9. Culligan, S. W., Yan, F., Yasuda, T., Fujita, K., Mourey, T., Chen, S. H.,
Tsutsui, T., Tang, C. W. Adv. Funct. Mater. (2005, submitted).
iv
ACKNOWLEDGEMENTS
I would like to first thank my thesis advisors Professor Shaw H. Chen and
Professor Ching W. Tang for their support, guidance, and hard work over the course
of our professional relationship. I consider myself fortunate to have worked with two
such exemplary scientists. This thesis is a result of their persistence and leadership.
I also wish to thank Professor Lewis Rothberg of the Department of
Chemistry and Professor Matthew Yates of the Department of Chemical Engineering
for serving as my committee members. Professor Rothberg also deserves my
gratitude for many helpful discussions over the years. It is a privilege to
acknowledge technical discussions and assistance from Professor Steve Jacobs and Dr.
Brian McIntyre of the Institute of Optics, Professor David Harding, Mr. Kenneth
Marshall, and Mr. Mark Bonino of the Laboratory for Laser Energetics, Dr. Jane Ou
of the University of Rochester, and Drs. Kathleen Vaeth, Thomas Mourey, Andrew
Hoteling, Craig Swanson, Ralph Young, Liang-Sheng Liao, Shelby Nelson, Diane
Freeman, and David Levy, all of the Eastman Kodak Company. Experimental
assistance provided by Mike Culver, Kevin Klubek, Dustin Comfort, and Andrea
Childs of Eastman Kodak is also gratefully recognized.
Dr. Dimitris Katsis and Dr. Yanhou Geng also deserve many thanks. Dr.
Katsis served as my experimental mentor early in my graduate career, while the
synthetic talents of Dr. Geng resulted in a majority of the materials used in this thesis.
Thanks also to Dr. Feng Yan for the synthesis of some of the materials used in
Chapter 4. Furthermore, I wish to thank Professor Tetsuo Tsutsui, Dr. Takeshi
Yasuda, and Dr. Katsuhiko Fujita of the Department of Applied Science for
v
Electronics and Materials, Kyushu University, Japan for their efforts in our
collaborative OFET research.
My sincere appreciation also goes to my fellow graduate students Andrew
Chien-An Chen, Anita Trajkovska, Jason Wallace, Prof. Huang-Ming Philip Chen,
Prof. Feng Yu Tsai, Yongfa Fan, Dr. Tanya Kosc, Lichang Zeng, and Ku-Hsieh
Simon Wei for technical discussions, assistance, and general camaraderie in the lab.
My family deserves a great deal of credit for this work, as their love and
encouragement made it possible. I would not have made it through the last few years
without my wife Candice, my parents Dennis and Lorraine, my brother Patrick, and
my grandmother Suzanne DelGiudice. I also wish to thank my in-laws Clinton, Joan,
Shane, and Brent Williams for always making me feel like a part of their family. My
deep appreciation also goes out to my good friends in Rochester, especially Brian
Lachance, who was really there for me in a time of need. My best also to Blake
Hepburn, Matt and Brianne Testa-Wojtezcko, Chris LeFeber and Judy King, Jay
Scherer, the Camaione-Lind family, the Chiumento family, and the Tallarico family
for making my time in Rochester more memorable and enjoyable.
This research was funded by the Eastman Kodak Company, the U.S. Army
Research Office, the New York State Center for Electronic Imaging Systems, the
National Science Foundation, and the U.S. Department of Energy (DOE) through the
Laboratory for Laser Energetics (LLE) and the New York State Energy Research and
Development Authority. The support of DOE does not constitute an endorsement by
DOE of the views expressed herein. The Honorable Frank J. Horton Fellowship
provided by LLE is recognized with deepest appreciation.
vi
ABSTRACT
Monodisperse conjugated oligomers are potentially useful for organic
electronics because of chemical purity and structural uniformity, facilitating the
elucidation of structure-property relationships and the realization of superior device
performance. Nematic liquid crystalline conjugated oligomers with above-ambient
glass transition temperatures represent an attractive approach to prepare uniaxially
aligned thin films across large areas for linearly polarized light emission and
anisotropic charge transport. The temporal stability of blue light-emitting materials is
also a major hurdle to practical applications.
This thesis research, therefore, has focused on organic electronic devices
incorporating monodisperse conjugated oligomers comprising fluorene and other
conjugated units for blue light emission and charge transport. Key results are
summarized as follows:
(1) The thermotropic and optical properties of monodisperse, conjugated
glassy-nematic liquid crystalline oligo(fluorene)s were characterized and related to
the molecular structures. The optical dichroism, birefringence, and polarization ratio
in fluorescence increased with the molecular aspect ratio. Annealed thin films
displayed strongly polarized blue emission with high quantum yield.
(2) The same monodisperse materials were incorporated in strongly polarized,
efficient, deep blue OLEDs. By employing a conductive alignment layer, an
vii
appropriate electron-transport material, and optimizing the light-emitting layer
thickness, world-record device performance was achieved.
(3) A comparative study of carrier transport in organic field-effect transistors
(OFETs) was undertaken to explore the effects of chain length, pendant structure, and
backbone composition on field-effect mobility of monodisperse glassy-nematic
conjugated oligomers and polymer analogues. The oligomers’ extended length was
found to dominate the carrier transport properties. In contrast, the charge carrier
mobility did not correlate with the persistence length of polymer analogues.
(4) Monodisperse model compounds were prepared to study the parameters
influencing OLED device stability. Consistent trends across three OLED
configurations were observed, which were attributable to the difference in hole
mobility. The hole mobilities, measured by transient electroluminescence, affected
the recombination zone width and therefore dictated the device performance
including lifetime.
viii
CONTENTS
Curriculum Vitae iii
iv
vi
xi
xii
xiii
Acknowledgements
Abstract
List of Charts and Reaction Schemes
List of Tables
List of Figures
1
1
4
6
9
10
17
18
20
22
24
28
1. Background and Introduction
1. Liquid Crystals
2. Nematic Liquid Crystals in Displays
3. Glassy Liquid Crystals (GLCs)
4. Organic Electronic Materials and Devices
5. Organic Light-Emitting Diodes (OLEDs)
6. Polarized OLEDs
7. OLED Stability
8. Organic Field-Effect Transistors (OFETs)
9. Conjugated Oligomers as an Alternative to Conjugated Polymers
10. Formal Statement of Research
References
2. Thermotropic and Optical Properties of Monodisperse Glassy-
Nematic Oligo(fluorene)s for Linearly Polarized Blue Emission 38
38
41
49
64
67
1. Introduction
2. Experimental
3. Results and Discussion
4. Summary
References
ix
3. Strongly Polarized, Efficient Deep Blue Electroluminescence Based
on Monodisperse Glassy-Nematic Oligo(fluorene)s 71
71
76
84
100
103
1. Introduction
2. Experimental
3. Results and Discussion
4. Summary
References
4. Organic Field-Effect Transistors Comprising Glassy-Nematic
Films of Conjugated Oligomers and Polymers 107
107
110
115
129
132
1. Introduction
2. Experimental
3. Results and Discussion
4. Summary
References
5. Effect of Hole Mobility through the Emissive Layer on Temporal
Stability of Blue Organic Light-Emitting Diodes (OLEDs) 137
137
139
150
164
166
1. Introduction
2. Experimental
3. Results and Discussion
4. Summary
References
6. Summary, Conclusions, and Potential Avenues for Future Work 170
1. Summary and Conclusions 170
2. Potential Avenues for Future Work 175
178 References
x
Appendices 180
Appendix 1: Second heating and cooling DSC thermograms for
monodisperse oligo(fluorene)s described in Chapter 2.
Appendix 2: Selected raw data (with and without polarization analysis)
for the polarized OLEDs reported in Chapter 3.
Appendix 3: Second heating and cooling DSC thermograms for
monodisperse co-oligomers and conjugated polymer analogues employed
in Chapter 4.
Appendix 4: Representative raw OLED data at a current density of 20
mA/cm2 for Chapter 5.
Appendix 5: MALD/I-TOF-MS results for model compounds and 1H-
NMR spectra for final products and key intermediates from Chapter 5.
180
195
214
220
230
xi
LIST OF CHARTS
2.1 Molecular structures of monodisperse oligo(fluorene)s comprising
five or seven fluorene units.
2.2 Molecular structures of monodisperse oligo(fluorene)s comprising
nine or twelve fluorene units.
3.1 Molecular structures of monodisperse glassy-nematic
oligo(fluorene)s and PEDOT/PSS.
4.1 Molecular structures of conjugated oligomers and polymers
accompanied by thermal transition temperatures from differential
scanning calorimetry (DSC) and polarizing optical microscopy
(POM). G, Glassy; N, Nematic; K, Crystalline; I, Isotropic.
5.1 Molecular structures of model compounds. 140
111
77
43
42
LIST OF REACTION SCHEMES
5.1 Synthesis procedures for model compounds. 141
xii
LIST OF TABLES
2.1 Phase transition temperatures (Tg and Tc), absorption and emission
dichroic ratios, orientational order parameters, and
photoluminescence quantum yields of thermally annealed
monodomain spin cast films.
3.1 Device features and polarized OLED performance data at a current
density of 20 mA/cm2 for devices with a 30 nm TPBI electron
transport layer.
4.1 Orientational order parameters of monodomain glassy-nematic films
determined by UV-Vis linear dichroism.
4.2 Saturation regime field-effect mobilities in cm2/V-s for OFETs
comprising monodomain and polydomain glassy-nematic, and
glassy-amorphous films with on/off ratios ranging from 103 to 104.
5.1 Measured oxidation potentials, calculated HOMO energy levels, and
optical bandgaps for model compounds and hole transport material
NPB.
5.2 Summary of drive voltages, EL efficiencies, and CIE coordinates at a
current density of 20 mA/cm2. The device lifetime at a current
density of 40 mA/cm2 is also included. 155
152
123
119
94
50
xiii
LIST OF FIGURES
1.1 Molecular arrangements characteristic of a) nematic and b) chiral-
nematic mesophases.
1.2 Schematic of a single color sub-pixel of a full-color twisted-
nematic liquid crystal (LC) display. The nematic director is
represented by the lines in the LC layer, corresponding to the
transmitting state for Δε > 0 without application of an electric
field. The transmission axes of the linear polarizers are crossed.
1.3 Glassy-nematic liquid crystals comprising mesogenic pendants
chemically attached to a volume-excluding core through an
aliphatic spacer. G, Glassy; Nm, Nematic; I, Isotropic.
1.4 a) Molecular structures of TAPC and Alq3 and b) Device
configuration of the first efficient thin film OLED.
1.5 Molecular structures of PPV and MEH-PPV.
1.6 Molecular structures of example blue light-emitting conjugated
polymers.
1.7 Molecular structures of a) PFO and b) F8T2.
1.8 Molecular structures of representative crystalline organic
semiconductors that display high OFET mobility.
2.1 Typical results for polyimide alignment layer buffing
optimization performed with a 0.8 wt % solution of
F(Pr)5F(MB)2. Rem refers to the dichroic ratio at the first
emission maximum (λ = 420 nm). The micrometer setting that
controls the height of the rotating drum is indicated. Parallel and
Perpendicular refer to the orientation of the first polarizer relative
to the nematic director.
3
6
9
11
12
16
17
21
46
xiv
2.2 Nematic threaded textures observed in hot stage polarizing optical
microscopy (POM) for micrometer-thick films prepared between a
microscope slide and a cover slip without surface treatment. a)
F(Pr)5F(MB)2 and b) F(MB)7F(EH)2.
2.3 Electron diffraction patterns of annealed glassy-nematic solid films
displaying diffuse rings characteristic of a non-crystalline
morphology. a) F(Pr)5F(MB)2 and b) F(MB)7F(EH)2.
2.4 Ellipsometric characterization for a monodomain film of
F(Pr)5F(MB)2. a) Ordinary (no) and extraordinary (ne) refractive
indices; b) Absorption coefficients parallel and perpendicular to the
nematic director defined by the buffing direction. The absorption
coefficients extracted by ellipsometric modeling agree well with
those measured independently by UV-Vis linear dichroism with a
spectrophotometer.
2.5 Anisotropic optical properties for representative oligo(fluorene)s of
increasing chain length. ne and no and subscripts || and ⊥ are as
defined for Figure 2.4. a) F(MB)5; b) F(MB)7F(EH)2; and c)
F(MB)10F(EH)2.
2.6 Normalized absorbance and photoluminescence spectra (with
excitation at 370 nm) of dilute solutions in UV grade chloroform
(10-5 M in fluorene units) and pristine films for representative
oligo(fluorene)s of increasing chain length. a) F(MB)4F(Pr)1 b)
F(Pr)5F(MB)2; c) F(MB)7F(DMO)2; and d) F(MB)10F(EH)2.
2.7 Linear dichroism and linearly polarized photoluminescence spectra
(with unpolarized excitation at 370 nm) of 80 to 90 nm thick,
thermally annealed monodomain films. Subscricpts || and ⊥ specify
polarization directions parallel and perpendicular to the nematic
director defined by the buffing direction (See Experimental). a)
F(MB)5; b) F(Pr)5F(MB)2; c) F(MB)7F(EH)2; d)
F(MB)10F(EH)2.
54
55
56
57
59
61
xv
2.8 UV-Vis absorption polar plot for a ~ 70 nm thick, monodomain
F(Pr)5F(MB)2 film.
2.9 Effects of prolonged thermal annealing of a 70 nm film of
F(MB)8F(EH)1 at a temperature of 120 oC (i.e. 5 oC above Tg)
under Argon. a) Increased π-orbital interactions due to molecular
alignment and close packing manifested by hypochromism; b)
Accompanying reduction in integrated photoluminescence
intensity; c) Normalized photoluminescence spectra of pristine and
96 hour annealed films, displaying the stability of the blue emission
color; and d) Increase in emission dichroic ratio from 1.1 to 14.4
upon annealing for 96 hours. Subscripts || and ⊥ are as defined in
Figure 2.7.
3.1 Figure 3.1: Typical results for PEDOT/PSS alignment layer buffing
optimization performed with a 0.8 wt % solution of F(Pr)5F(MB)2.
Rem refers to the dichroic ratio at the first emission maximum (λ =
420 nm). The micrometer reading that controls the buffing height
is indicated, and Parallel and Perpendicular refer to the orientation
of the first polarizer relative to the nematic director.
3.2 Figure 3.2: OLED device configurations (thicknesses in
Ångstroms) and molecular structures of electron transport
materials. PEDOT and oligo(fluorene) film thicknesses were
determined by spectroscopic ellipsometry after spin coating and
thermal annealing, while other layers were measured with a quartz
crystal microbalance during vacuum deposition. a) Polarized
OLED without an electron-transport layer; b) OLED with an Alq3
ETL; c) Polarized OLED with a TPBI ETL, in which the
oligo(fluorene) thickness was varied; d) Molecular structures of
Alq3 and TPBI.
62
64
79
83
xvi
3.3 Differential scanning calorimetry thermograms for the three
oligo(fluorene)s used in OLED studies. G Glassy; N, Nematic; I,
Isotropic. Clearing points and mesophase behavior were confirmed
by polarizing optical microscopy.
3.4 UV-Vis linear dichroism and linearly polarized photoluminescence
spectra for monodomain, ~ 70 nm thick films: a) F(MB)5 and b)
F(MB)10F(EH)2. The orientational order parameter Sabs and the
emission dichroic ratio at the first PL maximum Rem are also
included. Subscripts || and ⊥ refer to polarization directions relative
to the orientation of the nematic director defined by the buffing
direction.
3.5 Characteristic 400 μm2 tapping mode AFM surface roughness
analysis images for ~ 70 nm thick films of F(Pr)5F(MB)2 on
uniaxially rubbed PEDOT/PSS: a) Pristine film; b) Thermally
annealed and quenched; c) Thermally annealed and cooled at 5 oC/min; d) Thermally annealed and cooled at 5 oC/hr.
3.6 Polarized OLED properties of a device comprising a ~ 50 nm thick,
monodomain film of F(MB)10F(EH)2 without an electron-
transport layer: a) Current density-voltage-luminance (IV-LV)
relationships; b) Linearly polarized EL spectra at a current density
of 20 mA/cm2; subscripts || and ⊥ are as defined in Figure 3.4.
3.7 EL spectrum of an OLED comprising a ~ 70 nm thick,
monodomain film of F(MB)10F(EH)2 and an Alq3 electron-
transport layer (30 nm) at a current density of 20 mA/cm2.
3.8 Polarized EL spectra for Devices A, D, and E containing emissive
layers of ~ 70 nm thick, monodomain films of F(MB)10F(EH)2,
F(Pr)5F(MB)2, and F(MB)5, respectively, at a current density of
20 mA/cm2. Subscripts || and ⊥ are as defined in Figure 3.4.
85
87
89
91
93
96
xvii
3.9 Current density-voltage-luminance (IV-LV) relationships for
Devices A and C comprising monodomain F(MB)10F(EH)2 films
of different thickness.
3.10 Absolute polarized EL spectra for Devices B and C for
monodomain F(MB)10F(EH)2 films of different thickness at a
current density of 20 mA/cm2. Subscripts || and ⊥ are as defined in
Figure 3.4.
3.11 Comparison of EL spectra and performance data for essentially
isotropic and linearly polarized OLEDs for a) F(MB)5 and b)
F(MB)10F(EH)2.
4.1 Top-gate OFET device structure. Source/drain electrodes are gold,
40 nm thick. Channel lengths ranged from 75 to 120 μm, while the
channel width was 5 mm.
4.2 UV-Vis absorption spectra of dilute solutions (10–5 M repeat units
in CH2Cl2) and 50 to 60 nm thick, glassy-amorphous films of (a)
F(MB)10F(EH)2, F(Pr)5F(MB)2, and PFO, and (b) FTO-1
through FTO-3 and F8T2.
4.3 Polarizing optical micrographs of three glassy films of
F(Pr)5F(MB)2. a) Glassy-amorphous; b) Polydomain glassy-
nematic; and c) Monodomain glassy-nematic.
4.4 UV-Vis absorption spectra of 50 to 60 nm thick, glassy-amorphous
and monodomain glassy-nematic films of (a) F(MB)10F(EH)2,
F(Pr)5F(MB)2, and PFO, and (b) FTO-1, FTO-2, FTO-3 and
F8T2.
4.5 Output curves for the 50 to 60 nm thick, monodomain glassy-
nematic films of (a) F8T2, (b) FTO-1, and (c) PFO, all with chain
alignment parallel to current flow.
97
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100
115
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118
120
122
xviii
4.6 A schematic diagram of charge transport in an oriented film
comprising (a) a short, and (b) a long conjugated oligomer; charge
injection, parallel (i.e. intrachain) and perpendicular (i.e. interchain)
transport are denoted 1, 2, and 3, respectively.
4.7 Photoluminescence spectra of dilute (10–5 M in repeat units)
CH2Cl2 solutions and 50 to 60 nm thick, monodomain glassy-
nematic films of (a) PFO and (b) F(MB)10F(EH)2.
5.1 Vacuum evaporation system schematic.
5.2 a) OLED device structures with layer thickness specified in
Ångstroms and rubrene (Rub) doping levels of 5 and 10 % in
Devices II and III, respectively; b) Molecular structures of transport
materials and emitting dopants; c) Transient OLED device structure
with a C-545T doping level of 2.5 %.
5.3 UV-Vis absorption spectra of dilute solutions (10-5 M) of model
compounds in UV-grade CH2Cl2.
5.4 Normalized photoluminescence spectra of vacuum evaporated thin
films (30 or 50 nm thick) of model compounds with unpolarized
excitation at 360 nm.
5.5 Normalized EL spectra at j = 20 mA/cm2 for a) Device I; b) Device
II; and c) Device III comprising ADN, ANF, and ADF, all
providing evidence in support of a varying recombination zone
width.
5.6 a) Normalized Device I EL spectra comprising FFF and FDN
layers at a current density of 20 mA/cm2. The photoluminescence
spectrum of a thin film of NPB is shown for comparison.
5.7 Frequency-Dependent EL Quenching for a Transient OLED
Comprising a 3500 Å film of NPB upon Application of 8 V Square
Wave DC Voltage Pulses at an Increasing Frequency, 50 % Duty
Cycle. The solid line represents a best fit to a polynomial.
126
128
146
148
151
153
156
157
158
xix
5.8 a) Normalized steady-state EL spectrum of a transient OLED
containing 3000 Å of ADN as the model layer; b) Frequency-
dependent EL quenching of the transient OLEDs comprising model
compounds, 8 V DC square wave voltage pulses, 50 % duty cycle.
The solid lines represent best fits to polynomials.
5.9 Current density dependence of Device III EL spectrum for a) ADN;
b) ANF; c) ADF.
A1.1 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(Pr)5. K, Crystalline; I, Isotropic.
A1.2 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(Pe)5. G, Glassy; I, Isotropic.
A1.3 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)5. G, Glassy; N, Nematic; I, Isotropic.
A1.4 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(Pr)4F(EH)1. G, Glassy; N, Nematic; I, Isotropic.
A1.5 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)4F(Pr)1. G, Glassy; N, Nematic; I, Isotropic.
A1.6 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(Pr)5F(MB)2. G, Glassy; N, Nematic; I, Isotropic.
A1.7 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(Pr)5F(EH)2. G, Glassy; N, Nematic; I, Isotropic.
A1.8 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)7. G, Glassy; SmA, Smectic A; N, Nematic; I,
Isotropic.
A1.9 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)6F(EH)1. G, Glassy; N, Nematic; I, Isotropic.
159
160
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182
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184
185
186
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188
189
xx
A1.10 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)9. K, Crystalline; N, Nematic.
A1.11 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)8F(EH)1. G, Glassy; N, Nematic.
A1.12 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)7F(EH)2. G, Glassy; N, Nematic.
A1.13 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)7F(DMO)2. G, Glassy; N, Nematic; I, Isotropic.
A1.14 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F(MB)10F(EH)2. G, Glassy; N, Nematic.
A2.1 Raw data for a polarized OLED comprising a ~ 50 nm thick,
monodomain annealed film of F(MB)10F(EH)2 without an
electron transport layer. The data were collected without
polarization analysis.
A2.2 Raw data for a polarized OLED comprising a ~ 50 nm thick,
monodomain annealed film of F(MB)10F(EH)2 without an
electron transport layer. The data were collected with a linear
polarizer with its transmission axis parallel to the nematic director
defined by the buffing direction.
A2.3 Raw data for a polarized OLED comprising a ~ 50 nm thick,
monodomain annealed film of F(MB)10F(EH)2 without an
electron transport layer. The data were collected with a linear
polarizer with its transmission axis perpendicular to the nematic
director defined by the buffing direction.
A2.4 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected without
polarization analysis.
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xxi
A2.5 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis parallel to the nematic director
defined by the buffing direction.
A2.6 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis perpendicular to the nematic
director defined by the buffing direction.
A2.7 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(Pr)5F(MB)2 with a TPBI
electron-transport layer. The data were collected without
polarization analysis.
A2.8 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(Pr)5F(MB)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis parallel to the nematic director
defined by the buffing direction.
A2.9 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(Pr)5F(MB)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis perpendicular to the nematic
director defined by the buffing direction.
A2.10 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(MB)5 with a TPBI electron-
transport layer. The data were collected without polarization
analysis.
200
201
202
203
204
205
xxii
A2.11 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(MB)5 with a TPBI electron-
transport layer. The data were collected with a linear polarizer with
its transmission axis parallel to the nematic director defined by the
buffing direction.
A2.12 Raw data for a polarized OLED comprising a ~ 70 nm thick,
monodomain annealed film of F(MB)5 with a TPBI electron-
transport layer. The data were collected with a linear polarizer with
its transmission axis perpendicular to the nematic director defined
by the buffing direction.
A2.13 Raw data for a polarized OLED comprising a ~ 50 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected without
polarization analysis.
A2.14 Raw data for a polarized OLED comprising a ~ 50 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis parallel to the nematic director
defined by the buffing direction.
A2.15 Raw data for a polarized OLED comprising a ~ 50 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis perpendicular to the nematic
director defined by the buffing direction.
A2.16 Raw data for a polarized OLED comprising a ~ 35 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected without
polarization analysis.
206
207
208
209
210
211
xxiii
A2.17 Raw data for a polarized OLED comprising a ~ 35 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis parallel to the nematic director
defined by the buffing direction.
A2.18 Raw data for a polarized OLED comprising a ~ 35 nm thick,
monodomain annealed film of F(MB)10F(EH)2 with a TPBI
electron-transport layer. The data were collected with a linear
polarizer with its transmission axis perpendicular to the nematic
director defined by the buffing direction.
A3.1 Second heating and cooling DSC thermograms at ± 20 oC per
minute for FTO-1. G, Glassy; N, Nematic; I, Isotropic.
A3.2 Second heating and cooling DSC thermograms at ± 20 oC per
minute for FTO-2. G, Glassy; N, Nematic; I, Isotropic.
A3.3 Second heating and cooling DSC thermograms at ± 20 oC per
minute for FTO-3. G, Glassy; N, Nematic; I, Isotropic.
A3.4 Second heating and cooling DSC thermograms at ± 20 oC per
minute for F8T2. G, Glassy; K, Crystalline; N, Nematic; I,
Isotropic.
A3.5 Second heating and cooling DSC thermograms at ± 20 oC per
minute for PFO. G, Glassy; K, Crystalline; N, Nematic; I,
Isotropic.
A.4.1 Raw OLED data for Device I comprising ADN.
A4.2 Raw OLED data for Device I comprising ANF.
A4.3 Raw OLED data for Device I comprising ADF.
A4.4 Raw OLED data for Device II comprising ADN.
A4.5 Raw OLED data for Device II comprising ANF.
212
213
215
216
217
218
219
221
222
223
224
225
xxiv
226 A4.6 Raw OLED data for Device II comprising ADF.
A4.7 Raw OLED data for Device III comprising ADN.
A4.8 Raw OLED data for Device III comprising ANF.
A4.9 Raw OLED data for Device III comprising ADF.
A5.1 Positive ion MALD/I-TOF mass spectrum for ADN using DCTB as
the matrix.
A5.2 Positive ion MALD/I-TOF mass spectrum for ANF using DCTB as
the matrix.
A5.3 Positive ion MALD/I-TOF mass spectrum for ADF using DCTB as
the matrix.
A5.4 Positive ion MALD/I-TOF mass spectrum for FFF using DCTB as
the matrix.
A5.5 Positive ion MALD/I-TOF mass spectrum for FDN using DCTB as
the matrix.
A5.6 1H NMR spectrum of ADN in CDCl3.
A5.7 1H NMR spectrum of intermediate 5 in CDCl3.
A5.8 1H NMR spectrum of intermediate 6 in CDCl3.
A5.9 1H NMR spectrum of ANF in CDCl3.
A5.10 1H NMR spectrum of ADF in CDCl3.
A5.11 1H NMR spectrum of FFF in CDCl3.
A5.12 1H NMR spectrum of FDN in CDCl3.
227
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1
Chapter 1
Background and Introduction
1. Liquid Crystals
Materials that exist in the solid state at a given temperature and pressure may
be further categorized by the degree of order possessed by the system. In single
crystals, for example, atoms or molecules are arranged in a lattice structure with high
degrees of positional and, in some cases, orientational order. On the contrary, some
solids exist in an amorphous phase lacking long range order and exhibiting properties
that are isotropic, or independent of the direction in which they are measured. While
most liquids are also isotropic, liquid crystals are a class of ordered fluids that flow
like liquids, but possess orientational and/or positional order similar to crystals1.
Liquid crystals may also be referred to as mesophases or mesomorphic materials.
The presence of order and the resulting anisotropy in optical and electronic properties
is based on the shape of the liquid crystalline molecule, allowing liquid crystals to be
classified accordingly. For example, mesomorphic materials comprised of molecules
with a rod-like shape are known as calamitic liquid crystals, while materials with a
disk or coin-like shape may form so-called discotic mesophases. Calamitic and
discotic liquid crystals were discovered nearly 90 years apart by Reinitzer2 and
Chandrasekhar et al3, respectively.
2
Nematic liquid crystals are a subset of calamitic materials exhibiting only
orientational ordering that have been widely explored for information display
applications. A schematic representation of the molecular arrangement observed in a
nematic liquid crystal domain is shown in Figure 1.1a4. Within the domain, the long
molecular axes tend to point in a preferred direction. By considering the nematic
molecules as vectors, the direction of the molecular long axes and the average angle θ
between the long axes and this direction can be characterized. The preferred direction
is referred to as the nematic director, while the orientational order parameter S for the
arrangement can be defined according to Equation 1-15.
1cos321 2 −= θS (1-1)
Different domains in a nematic liquid crystalline sample will show varying
director orientations, resulting in domain boundaries and characteristic textures that
can be observed between crossed polarizers in polarizing optical microscopy (POM)6.
A sample consisting of many domains can strongly scatter incident light depending
on the domain size and will show a macroscopic order parameter S = 0. For practical
applications in optical devices, light scattering at domain boundaries must be
alleviated by preparing a so-called monodomain, in which all domains in the sample
share a common director orientation. A monodomain film can be realized through the
use of an alignment layer such as uniaxially rubbed polyimide or poly(vinyl alcohol)7.
A monodomain may be formed with the director oriented either parallel or
perpendicular to the sample surface, known as homogeneous or homeotropic
orientation, respectively. Other classes of calamitic liquid crystals with higher
degrees of ordering compared to nematic mesophases are also well-known. Smectic
3
liquid crystals, for example, possess positional ordering in a variety of layered
structures in addition to orientational order8. Smectic liquid crystals were not
explored in detail in this thesis research and therefore will not be discussed further.
Figure 1.1: Molecular arrangements characteristic of a) nematic and b) chiral-nematic
mesophases4.
a) b)
If a chiral moiety is introduced into a nematic liquid crystal, an optically
uniaxial chiral-nematic mesophase may result. Homogeneous alignment of a chiral-
nematic liquid crystal (CLC) is characterized by the molecular arrangement depicted
in Figure 1.1b4, in which uniaxial order is maintained in each molecular layer while
the nematic director rotates along a helix in the direction normal to the film surface.
The distance over which the director rotates 360 degrees is referred to as the helical
pitch length, with the rotation direction defining the handedness. This class of
materials shows selective reflection of circularly polarized light, a unique optical
property9. Circularly polarized light incident on a chiral-nematic liquid crystalline
film will be reflected and maintain its polarization state if it has the same handedness
as the film and the wavelength matches the product of the helical pitch length and the
4
average refractive index of the CLC, while other wavelengths and polarization states
are unaffected. Selective reflection from a CLC film can be useful for passive optics
such as polarizers and reflectors with low losses.
Liquid crystalline mesophases may appear in response to changes in
temperature or in concentrated solutions, so-called thermotropic and lyotropic liquid
crystals10, respectively. Since lyotropic liquid crystals were not studied in this work,
they will not be discussed further. Thermotropic liquid crystals exhibiting mesophase
behavior upon both heating and cooling are termed enantiotropic, while monotropic
materials show liquid crystallinity upon either heating or cooling. All of the liquid
crystals explored in this thesis research exhibited enantiotropic nematic mesophases.
2. Nematic Liquid Crystals in Displays
Flat panel displays (FPDs) represent a worldwide technology market that
continues to grow as a result of the proliferation of personal electronics such as
cellular phones, personal digital assistants, and portable media players, along with the
continued displacement of cathode ray tubes (CRTs) in desktop computing and
television applications. The world-wide FPD market is projected to surpass US $50
billion in 200511. Despite growing competition from other classes of FPDs, the
dominant technology to date in terms of market share remains the liquid crystal
display (LCD)11. Nematic liquid crystals form the backbone of twisted-nematic
(TN)12 and super-twisted nematic (STN)13 displays, as well as many other display
modes. Electro-optic switching of nematic liquid crystals takes advantage of the
dielectric anisotropy Δε and optical birefringence Δn of the ordered fluid state. The
5
application of an electric field can result in the re-orientation of aligned nematic
liquid crystals to a direction either parallel or perpendicular to the field depending on
the sign of Δε14. Monodomain, homeotropically aligned films are optically isotropic
in the plane parallel to the film surface, which is also the plane of polarization for
light propagating along the surface normal. In a simple twisted-nematic cell, two
alignment layers prepared for homogeneous liquid crystal orientation are combined
with mutually perpendicular alignment directions, resulting in a 90 degree rotation of
the nematic director along the cell thickness d. When the retardance of the liquid
crystal layer Δn×d is much greater than the wavelength of light propagating through
the cell, the so-called Mauguin condition is met and waveguiding takes place, rotating
the plane of polarization of a linearly polarized beam by 90 degrees as it traverses the
sample14. Placing such a TN cell comprising a liquid crystal with Δε > 0 between
crossed polarizers results in transmission when no voltage is applied and extinction
when homeotropic alignment is obtained by application of the proper electric field. A
combination of individually addressable twisted-nematic pixels, a pair of crossed
polarizers, a backlight, and an array of absorbing color filters can be used in a full-
color LCD, as depicted schematically for a single-color sub-pixel in Figure 1.2.
Nematic LCDs suffer from image sticking, difficulties in achieving video refresh
rates, high power consumption due to the incorporation of absorbing elements, and
viewing angle dependence15. Significant research and development efforts are
underway to address these issues for improved display performance and reduced
cost16.
6
Figure 1.2: Schematic of a single color sub-pixel of a full-color twisted-nematic
liquid crystal (LC) display. The nematic director is represented by the lines in the LC
layer, corresponding to the transmitting state for Δε > 0 without application of an
electric field. The transmission axes of the linear polarizers are crossed.
Backlight Diffuser
Linear Polarizer Glass Substrate
Transparent Electrode Alignment Layer
LC
Alignment Layer Transparent Electrode
Glass Substrate Linear Polarizer
Color Filter
3. Glassy Liquid Crystals (GLCs)
While liquid crystalline fluids are uniquely suited for electro-optic switching,
their optical properties combined with ease of fabrication across large apertures are
also desirable for passive optics including waveguides and waveplates, polarizers,
reflectors, and color filters17. For practical applications, solid films are preferred over
fluids to reduce the potentially deleterious effects of mechanical or thermal shock.
Furthermore, solid films offer the possibility to decrease the number of bulky and
expensive substrates needed for film preparation. However, most liquid crystals tend
to crystallize upon cooling from the mesomorphic melt, resulting in polycrystalline
7
solid films that scatter light from grain boundaries, limiting applicability in optics or
photonics. Approaches to preserve self-assembled liquid crystalline order in the solid
state have therefore attracted research interest in recent years. Liquid crystalline
polymers can provide stable morphologies and possess the requisite physical
properties for film and fiber formation, but are often intractable for devices because
of high melt viscosity and chain entanglements that limit alignment18. Photo-
crosslinkable liquid crystals have been developed in an effort to address this
limitation19. Consisting of liquid crystalline monomers with photoreactive end
groups, facile processing to well-ordered films can be accomplished in a manner
identical to low molecular weight liquid crystals. Exposing the film to UV irradiation
cross-links the photoreactive groups, preserving the self-assembled, macroscopically
oriented film in a solid network. This approach can be applied to prepare solid chiral-
nematic films with a gradient pitch for broadband selective reflection and circular
polarization19, 20.
Molecular design of materials to encourage vitrification as opposed to
crystallization upon cooling from the mesomorphic melt represents another strategy
to preserve the orientation of liquid crystalline fluids in the solid state. Substantial
efforts to synthesize organic materials with elevated glass transition temperatures (Tg)
relative to the desired application temperatures have resulted in an array of successful
structural motifs including bulky or unusual molecular shapes21, starburst or
dendrimeric structures22, and spiro-configured systems23. All of these approaches
targeted amorphous solids with presumably isotropic properties as opposed to
mesomorphic solids. The first low molar mass liquid crystals capable of vitrification
8
exhibited glass transition temperatures of 22 and – 65 oC for chiral-nematic and
nematic materials, respectively, along with poor morphological stability against
crystallization24. While various material design strategies have been employed over
the past 3 decades targeting glassy liquid crystals (GLCs) with high transition
temperatures and resistance to thermally activated crystallization25, the approach of
Chen et al has been arguably the most successful26-28. By chemically attaching liquid
crystalline pendants to a volume-excluding core through a flexible aliphatic spacer,
hybrid molecular systems comprised of three distinct structural elements that will
each crystallize readily as separate entities form morphologically stable glass-forming
liquid crystals with high transition temperatures. Figure 1.3 shows exemplary
nematic GLCs prepared following this molecular design strategy with high glass
transition temperatures, relatively broad nematic fluid temperature ranges, and large
Δn values at 632.8 nm29. Glassy liquid crystals also possess enhanced chemical
purity achieved through column chromatography or recrystallization and superior
rheological properties compared to liquid crystalline polymers30. Optical elements
including high-performance circular polarizers, reflectors, and notch filters have been
prepared from solid films of chiral-nematic GLCs designed following this core-
pendant approach31. Furthermore, functionality can be introduced through the
volume-excluding core and/or the mesogenic pendant, as demonstrated recently32.
Liquid crystals with above-ambient glass transition temperatures also offer the
potential for temperature-gated responses33, allowing property changes induced by
external stimuli such as light or temperature in the mesomorphic fluid state to be
frozen in a solid film by cooling to below Tg.
9
Figure 1.3: Glassy-nematic liquid crystals comprising mesogenic pendants chemically
attached to a volume-excluding core through an aliphatic spacer. G, Glassy; Nm,
Nematic; I, Isotropic29.
N1 : CN
(CH2)3O N2 : (CH2)3O CN
ON2
N2OOC COON2
COON1
N1OOC COON1
COON1
O
NO2
N1OOC
COON1
COON1
N1OOC
N1OOC
N1OOC
OOC COO
COO
COON1
COON1
COON2
O
NO2
N2OOC
COON2
COON2
N2OOC
N2OOC
N2OOC
OOC COO
COO
COON2
COON2
G 75oC Nm 235oC I G 82oC Nm 347oC I
G 86oC Nm 288oC IG 76oC Nm 153oC I
G 109oC Nm 183oC I G 127oC Nm 308oC I
4. Organic Electronic Materials and Devices
The field of organic electronic materials and devices has been intensively
explored over the past twenty years34. Both charge-transporting organic materials
doped into polymer films for application in xerography35 and electronic processes in
aromatic organic crystals34c have been studied for decades. The field was re-
invigorated in 1977 by the discovery of high conductivity in iodine-doped
polyacetylene36, which also contributed to MacDiarmid, Heeger, and Shirakawa
sharing the Nobel Prize in Chemistry in 2000. In particular, materials and device
10
science for applications in organic light-emitting diodes (OLEDs)37 and organic field-
effect transistors (OFETs)38 have advanced rapidly since their respective
discoveries39, 40. A brief overview of OLEDs and OFETs is provided in the following
sections.
5. Organic Light-Emitting Diodes (OLEDs)
Electroluminescence (EL) from an organic material was first observed by
Bernanose et al in 1953 for an acridine orange dye41. This phenomenon did not
receive considerable research attention until the thin film OLED was reported in
198739, exhibiting display-level brightness at DC drive voltages of less than 10 V. In
principle, this device was similar to a conventional semiconductor p-n junction, with
sequentially vacuum-evaporated thin films of an arylamine derivative and an
aluminum complex incorporated as hole- and electron-transporting layers,
respectively. A transparent anode of indium-tin oxide (ITO) and a magnesium:silver
alloy cathode completed the OLED. Like the inorganic device, the OLED exhibited
current rectification, with green light emission from the electron-transporting material
observed only for forward biases larger than a turn-on voltage of approximately 2.5
Volts. The external quantum efficiency was around 1 %, defined in terms of the
photons collected through the transparent anode per injected electron. Furthermore,
at an initial luminance of 50 cd/m2, the device half-life, or the time required for the
brightness to reach 25 cd/m2, was nearly 100 hours. Figure 1.4 shows the molecular
structures of the hole-transporting 4,4'-cyclohexylidenebis[N,N-di-p-tolylaniline]
11
(TAPC), the electron-transporting and green light-emitting tris-(8-hydroxyquinoline)
aluminum (Alq3), and the device configuration employed in this breakthrough work.
Figure 1.4: a) Molecular structures of TAPC and Alq3 and b) Device configuration
of the first efficient thin film OLED39.
OLEDs with an emissive layer based on conjugated polymers, which offer the
additional advantage of solution processability, were realized independently by
Burroughes et al in 199042 and Braun and Heeger in 199143. The preparation of
polymeric LEDs (PLEDs) from solution provides enhanced potential for large area,
inexpensive flat panel displays on flexible substrates44. The original reports described
PLEDs comprising green light-emitting poly(p-phenylene vinylene) (PPV) or a
soluble, orange-emitting derivative thereof, poly(2-methoxy-5-(2’-ethylhexyloxy)-
1,4-phenylene vinylene) (MEH-PPV), with external quantum efficiencies up to 0.05
and 1 %, respectively. The much higher quantum efficiency for the MEH-PPV
device may have been due to improved electron injection from calcium, a low work
function metal cathode. The structures of PPV and MEH-PPV are depicted in Figure
1.5.
N N
N
ON
OAl
NO
Mg:Ag a) b)
Alq3 (60 nm)
TAPC (75 nm)
ITO Alq3
TAPC
12
Figure 1.5: Molecular structures of PPV and MEH-PPV42,43.
n
O
Om
PPV MEH-PPV
These discoveries demonstrated that organic materials could be useful for flat
panel displays with potential advantages over LCDs, including reduced thickness and
weight, faster response times, and wider viewing angles. The realization of highly
efficient emission in each of the three primary colors is necessary for OLEDs to
compete with LCDs in the flat panel display market. OLEDs emitting white light
combined with color filters could also be used for full-color displays. Potential
applications in solid-state lighting45 provide additional motivation for the
development of white OLEDs. For any practical device, the temporal stability under
continual forward bias is a critical issue46, which will be discussed in a subsequent
section. Some of the key developments in materials and device science for vacuum
evaporable molecular materials and conjugated polymers are briefly reviewed in the
following.
The efficiency of the first OLED capable of achieving display-level brightness
at practically useful driving voltages was hindered by the relatively low solid-state
photoluminescence quantum yield of Alq3. The introduction of the three-layer
device47 and the incorporation of light-emitting dopants in the electron-transporting
layer48 were two early developments that aimed to overcome this limitation. The
13
three-layer approach separated electron-transporting and light-emitting functions,
permitting each to be optimized independently and allowing the emission color to be
tuned by varying the structure of the emissive material. Including a light-emitting
dopant in the electron-transporting layer improved the OLED efficiency from 1 % to
2.5 % because of the enhanced PL quantum yield of the light-emitting solid solution,
while providing another approach to adjust the emission color. Furthermore, the
device lifetime was extended to several hundred hours. Blue OLEDs were explored
in detail for the first time in 199049, completing the color gamut. The most efficient
blue device in this work employed a three-layer structure, with the blue emitter
sandwiched between a hole-transporting aromatic diamine and an electron-
transporting layer consisting of an oxadiazole derivative. The stability of OLEDs
containing a doped emitting layer was dramatically improved to beyond 4000 hours
from an initial luminance of 500 cd/m2 by incorporating a copper phthalocyanine hole
injection layer, using N,N'-Bis(1-naphthyl)-N,N'-diphenyl-4,4'-benzidine (NPB) as a
hole transport layer, and applying alternating, rather than direct, current to drive the
device50.
The work described above showed that OLEDs based on vacuum-evaporated
thin films of molecular materials could exhibit the required efficiencies and lifetimes
for display applications, paving the way for an outpouring of research efforts with
further improvements in device performance as a primary objective. The synthesis
and characterization of many classes of amorphous materials for charge transport22a, b,
51 or efficient emission of red52, 53, green54, 55, or blue56, 57 light in neat or doped films
has been reported. White OLEDs have also been pursued, with a focus on simple
14
architectures that can achieve high efficiencies along with appropriate and current-
density independent chromaticity coordinates58. Electrophosphorescence has been
explored in an effort to further increase efficiencies59. For EL to be observed,
injected electrons and holes must recombine to form emissive excitons. In
fluorescent materials, a singlet π-π* transition is typically the lowest optically
allowed electronic excitation, whereas either singlet or triplet excitons may be formed
by recombining charge carriers in OLEDs. Spin statistics would predict the
formation of three triplets for each singlet based on the energy degeneracy of the
triplet state60, limiting the fraction of excitons contributing to the emission to 1 out of
4 because the triplet excited state to ground state transition is optically forbidden.
The design of materials capable of efficient phosphorescent emission from the triplet
excited state was therefore a logical approach to improve OLED efficiency. OLEDs
containing iridium complexes as phosphorescent dopants with record-high external
quantum efficiencies up to 20 % have been demonstrated61-63. Despite the impressive
performance of red, green, and white phosphorescent devices, deep blue fluorescent
OLEDs still hold an advantage in efficiency over the phosphorescent devices reported
to date with comparable chromaticity coordinates64, for which the stability has also
not been disclosed. Phosphorescent OLEDs will not be discussed further since the
OLED research in this thesis focused on deep blue, linearly polarized OLEDs and the
stability of blue fluorescent OLEDs.
In parallel to the efforts to enhance the performance of OLEDs based on
vacuum-evaporable small molecules, substantial research has focused on improving
PLEDs based on conjugated polymers. The inclusion of transparent thin films of
15
conducting polymers such as poly(aniline) (PANI) or poly(ethylenedioxythiophene)
(PEDOT) doped with polycations such as poly(styrene sulfonic acid) (PSS) as hole-
injection layers improved the efficiency and lifetime of MEH-PPV devices65.
Because charge transport in PPV devices was found to be hole dominated66,
conjugated polymers with increased electron affinities or higher electron mobilities
received considerable attention. Substitutions of the polymer backbone with strong
electron-withdrawing groups67 or co-polymerizations with electron-transporting units
such as oxadiazole68 have been successfully implemented. In particular, luminance
efficiencies over 20 cd/A have been realized for oxadiazole-substituted PPVs in
PLEDs with air-stable aluminum cathodes68c. Obtaining deep blue emission proved
challenging with PPVs69, providing additional motivation for research into other
classes of conjugated polymers capable of generating efficient EL across the visible
spectrum.
An unsubstituted poly(p-phenylene) (PPP) film prepared from soluble
precursors was employed as the light-emitting layer in the first blue PLED, with a
quantum efficiency of 0.05 %70. Attaching appropriate pendants to phenylene
monomers prior to polymerization was shown to solubilize the resulting polymers,
but also produced an undesirable blue shift of the emission color because of steric
interactions between the substituents on adjacent phenylene rings71. The preparation
of ladder-type poly(phenylene)s (LPPPs) has been demonstrated as an alternative to
improve the solubility while planarizing the conjugated backbone to maintain the blue
emission color72. Various other classes of conjugated polymers have been explored
for blue light emission, including poly(carbazole)s73, poly(dibenzosilole)74, and
16
poly(p-phenylene ethnylene)s75, for example. Polymer blends have also been used as
the active layer in blue PLEDs with limited success76. The structures of several blue
light-emitting conjugated polymers are presented in Figure 1.6.
Figure 1.6: Molecular structures of example blue light-emitting conjugated
polymers70, 72c, 73a.
n
C10H21
C10H21
H13C6
C6H13
N
OC8H17
H17C8OC8H17
m
PPP PCMEHP
MeLPPP
Poly(fluorene)s, an emerging class of conjugated polymers pursued for blue
emission, possess an intermediate degree of backbone planarization compared to PPP
and its ladder-type analogue provided by a carbon atom fusing each biphenyl group77.
This carbon atom at the 9-position of the resulting fluorene unit can be readily
functionalized with alkyl, aryl, or other substituents to increase solubility, promote
liquid crystalline mesomorphism78, or incorporate additional functionalities such as
charge injection or transport79. Copolymers of fluorene with undisclosed chemical
structures have been used as the emissive layer in blue PLEDs with high efficiencies
up to 5 cd/A and promising temporal stabilities80. Mesomorphic poly(fluorene)s have
been employed in both linearly and circularly polarized PLEDs81, while nematic
poly(fluorene)s have been utilized in organic field-effect transistors (OFETs) with
17
anisotropic mobilities82. The structures of a liquid crystalline fluorene homopolymer
with linear octyl pendants (PFO) for blue light emission78, 81c and a copolymer of
fluorene and bithiophene (F8T2) used in OFETs82a are shown in Figure 1.7.
Figure 1.7: Molecular structures of a) PFO and b) F8T2.
nS
S
m
a) b)
In poly(fluorene) homopolymers with linear alkyl chains, early work showed
that the purity of the blue emission color was degraded by the appearance of a yellow
emission band upon exposure to UV light, air, or thermal annealing, which was
attributed to the formation of excimers and/or oxidation at the chain ends83. More
recently, the formation of ketonic defects on single polymer chains has gained
support as the primary mechanism responsible for this observation84. Appropriate
functionalization at the 9, 9-position suppressed the yellow emission band79, which
could also be prevented by purification of fluorene monomers85. Beyond the
considerable efforts to improve blue light-emitting poly(fluorene)s, copolymers of
fluorene and physically doped poly(fluorene) films have been reported in highly
efficient green, red, and white PLEDs86-87.
6. Polarized OLEDs
Most of the OLEDs and PLEDs discussed in the previous section emit
unpolarized light because the emission transition dipole moment is randomly oriented
18
in the plane of the light-emitting organic thin film prepared by vacuum evaporation or
solution processing. For many organic materials with extended π-electronic systems,
the emission transition dipole moment is approximately parallel to the long molecular
axis, suggesting that macroscopic uniaxial alignment of the long axes could result in
the emission of linearly polarized light88. An efficient and stable linearly polarized
light source would benefit a variety of display technologies89, providing motivation
for research into materials and devices capable of linearly polarized emission.
Molecular self-assembly mediated by nematic mesomorphism appears to be the most
promising method to prepare macroscopically ordered films that can emit linearly
polarized light directly90. Poly(fluorene)s as described above are an important class
of main-chain conjugated polymers that can show liquid crystalline mesomorphism
depending on the substituent at the 9,9-position81, 91. Thermal annealing of a
poly(fluorene) film in the mesomorphic fluid state followed by cooling to room
temperature on an orienting surface such as a rubbed PPV film promotes macroscopic
orientation of the conjugated backbones, resulting in linearly polarized blue EL with
polarization ratios up to 2581b but generally low efficiencies. Further improvements
in materials and device design are needed to increase the competitiveness of polarized
OLEDs relative to their unpolarized equivalents.
7. OLED Stability
OLEDs based on fluorescent and phosphorescent small molecules and
conjugated polymers have progressed substantially since their respective discoveries
and are poised to compete in the flat panel display market because of improvements
19
in device efficiency and lifetime. However, the issue of operational stability is still a
major challenge to OLED display technology46. Two primary mechanisms for the
decrease in OLED luminance under prolonged driving conditions have been reported
for Alq392. This material is widely used as a host for efficient green or red OLEDs
and as an electron-transport layer, prompting the efforts to understand and improve
the lifetime for Alq3-containing OLEDs. Mixed host OLEDs, in which the abrupt
organic/organic heterojunction between hole- and electron-transporting layers is
removed in favor of a single layer of uniformly93 or graded94 mixed transport
materials, have also shown better temporal stability than conventional OLEDs. The
improved lifetime has generally been attributed to an enhanced recombination zone
width and/or better charge balance. Either mechanism of Alq3 device degradation
could be supported by the mixed host OLED results, suggesting that more research is
needed to resolve this issue. In contrast to green-emitting Alq3 OLEDs, the stability
of blue devices has remained largely unexplored despite the need for blue light-
emitting materials in full-color OLED displays. Although blue OLEDs with
promising efficiencies and lifetimes have been reported57, the temporal stability of
blue devices is poor compared to state-of-the-art green and red OLEDs. Efficient and
stable blue materials are also necessary for the production of white OLEDs capable of
meeting display and solid-state lighting requirements. With an ultimate goal of
superior materials and devices, identification of the key factors that influence OLED
stability for blue emitters is crucial to the continued development of OLED
technologies.
20
8. Organic Field-Effect Transistors (OFETs)
The transistor was discovered by researchers at Bell Labs and reported in
194895, and few electronic devices have had as substantial an impact on the
technology used in everyday life, with the development and miniaturization of
transistors based on silicon providing the driving force for the evolution of personal
computers. Despite very high charge carrier mobilities in crystalline samples,
inorganic semiconductors are challenging to implement in applications requiring
large area coverage, low cost, low temperature processing, or flexibility. Since the
discovery of the OFET in 198640, organic semiconductors with high mobilities have
been pursued for a variety of device applications96. In the first device, a hole mobility
of 10-5 cm2/V-s was achieved with an unsubstituted polythiophene as the organic
semiconductor. In subsequent years, organic transistors employing polycrystalline
thin films attracted attention because of promising results for oligo- and
poly(thiophene)s38a, 97, 98 and pentacene38b, 99. In general, high mobility is obtained in
single- or polycrystalline organic semiconductors through charge transport along
highly ordered π-stacks100. Organic semiconductors with high mobilities could be
useful as an inexpensive alternative to amorphous silicon for applications such as
active matrix LCD backplanes, where field-effect mobilities on the order of 1 cm2/V-
s are needed. Mobilities reaching and exceeding this level have been achieved with
crystalline organic semiconductors. However, single crystals are limited to small
areas and are challenging to prepare and implement in practical devices, while the
mobility of a polycrystalline film is known to vary with grain size, crystallographic
defects, and orientation and order within the π-stacks38, 97b, 98. Figure 1.8 shows the
21
structures of several organic semiconductors that can be prepared as poly- or single-
crystalline films with high mobilities in OFETs.
Figure 1.8: Molecular structures of representative crystalline organic semiconductors
that display high OFET mobility.
S
S
S
S
S
S
Pentacene Sexithiophene
Rubrene
As a class of ordered organic semiconductors, liquid crystals have also been
attempted to achieve high charge carrier mobilities101. Smectic and discotic liquid
crystalline fluids have displayed mobilities on the order of 0.1 cm2/V-s based on time
of flight transient photocurrent measurements, competitive with the best solution-
processable conjugated polymer OFET results reported to date98, 102, 103. To fabricate
practically useful electronic devices from liquid crystals, the mesophase should be
preserved in the solid state without encountering crystallization. Glassy liquid
crystalline conjugated polymers, such as poly[(9,9-dioctylfluorenyl-2,7-diyl)-co-
bithiophene] (F8T2), have demonstrated anisotropic field-effect mobilities on the
order of 10−2 cm2/V-s82a, the highest value reported to date for a nematic material. π-
Conjugated liquid crystals that can achieve high mobilities in the solid state without
crystallization may represent an avenue towards large area, ordered films for OFETs
prepared by solution methods. Furthermore, the influence of factors such as
molecular length, contact resistance, and molecular aggregation on the charge-carrier
22
transport properties of liquid crystalline semiconductors remains practically
unexplored.
9. Conjugated Oligomers as an Alternative to Conjugated Polymers
Conjugated polymers have been employed in a broad array of device
applications because their properties can be readily tuned through molecular design
and synthesis. However, the synthesis procedures typically followed to prepare
conjugated polymers result in distributions of molecular weight (polydispersity),
conjugation length, and potentially, in the case of copolymers, chemical composition.
Furthermore, polymer chain defects such as kinks, bends, or chain ends cannot be
avoided and may have profound effects on device performance. Conjugated
polymers are generally purified by precipitation and extraction, which are unlikely to
remove high molecular weight polymer chains that contain defects or traces of
impurities. The purity level of organic semiconductors is vital for photonic and
electronic applications104, as impurities may act as charge traps, emission quenching
centers, or non-radiative recombination sites.
The preparation of monodisperse conjugated oligomers is a potentially elegant
approach to overcome the difficulties associated with polydispersity and purity in
conjugated polymers105. In a monodisperse system, each molecule is comprised of an
identical number of conjugated units, in contrast to conjugated polymers which
contain a distribution of chain lengths. Conjugated oligomers are an interesting class
of advanced materials for fundamental study and device applications in organic
electronics106 because of this structural regularity along with enhanced chemical
23
purity achieved through re-crystallization, column chromatography, or temperature
gradient vacuum sublimation. A substantial amount of theoretical work to predict or
describe the optical and electronic properties of conjugated oligomers has been
performed100, 107. From an experimental point of view, conjugated oligomers have
often served as model systems for polymer analogues105, although many instances of
disagreement between observed polymer properties and those predicted by the
extension of oligomer results have been noted107e, 108. Furthermore, the photophysical
behavior of some conjugated oligomers suggests that fundamental differences exist
between rod-like oligomers and polymers containing a distribution of conjugation
lengths along with kinks and chain folds109. In addition to the scientific interest in the
optical and electronic properties of well-defined organic semiconductors,
monodisperse conjugated oligomers have demonstrated superior performance to
polymer analogues in several photonic and electronic device applications110. The fact
that the performance of conjugated polymer devices varies with the polymer’s
molecular weight111 makes a comparison of device results questionable, even for
different synthetic batches of the same conjugated polymer, in the absence of strict
controls over this parameter. All of these factors provide motivation to explore
device applications with conjugated oligomers as an alternative to polymers.
Conjugated oligomers might be expected to show calamitic mesophases
because of their rod-like molecular shapes. On the contrary, there have been few
reports of liquid crystalline conjugated oligomers112. Furthermore, with a few notable
exceptions110a, 113, a glass transition temperature was not reported for the conjugated
oligomers that do exhibit mesomorphism. Soluble, liquid crystalline conjugated
24
oligomers with above-ambient glass transition temperatures would represent an
attractive approach to prepare macroscopically ordered solid films by spin coating
and thermal annealing above Tg followed by cooling to room temperature, provided
that thermally activated crystallization can be avoided. The structural uniformity and
chemical purity offered by oligomers suggests that the pursuit of structure-property
relationships of fundamental interest and practical utility in the design and
interpretation of organic electronic device experiments may be fruitful.
10. Formal Statement of Research
Organic electronic materials has been one of the most active research areas of
the past two decades because the optical and electronic properties of organic
semiconductors facilitate the preparation of useful devices with the potential for
drastically lower costs compared to inorganic counterparts. Organic light-emitting
diodes based on small molecules and conjugated polymers are currently available in
commercial products ranging from monochromatic cellular phone displays to full-
color viewfinders for digital cameras. However, further advances in OLEDs will
depend on the identification of the critical factors that determine temporal stability
from both the molecular structure and device architecture points of view. This
problem is particularly relevant for blue OLEDs due to their relatively short lifetimes
compared to green and red devices. Despite improvements in OLEDs, liquid crystal
displays are expected to maintain their position as the leader in the flat panel display
industry in the coming years11. Although LCDs are a mature technology, substantial
improvements in thickness, weight, and, particularly, power consumption can still be
25
achieved. The incorporation of absorbing elements significantly reduces the amount
of light generated by the backlight that can reach the viewer. Backlights based on
OLEDs have been proposed to reduce power consumption, but conventional OLEDs
would not alleviate the need for linear polarizers, adding to the current density
required to produce display-level brightness and thus limiting the OLED lifetime.
Obviously, OLEDs that emit linearly polarized light directly are potentially useful for
LCDs, and may find additional applications in other display technologies. A variety
of approaches to achieve linearly polarized electroluminescence have been
demonstrated, with generally poor device performance. Polarized OLEDs based on
liquid crystalline poly(fluorene)s with high polarization ratios81b, e but poor to
moderate efficiencies81d have been realized. However, the efficiency, emission
dichroism, and color purity of polarized OLEDs must be improved if they are to
compete with unpolarized equivalents. Blue light-emitting devices are particularly
attractive, as blue light can be efficiently down-converted to the green or red for full-
color display applications. Further research towards materials and devices capable of
linearly polarized light emission is warranted based on the expected benefits to
display technologies.
Organic field-effect transistors (OFETs) have also reached practical
performance levels for certain device applications in terms of mobility, on/off ratio,
and current capacity. However, most OFETs exhibiting mobilities competitive with
amorphous silicon are comprised of vacuum-evaporated materials with a
polycrystalline or single-crystalline morphology. To reap the full benefits of organic
semiconductors as low-cost alternatives to inorganic analogues, solution-processed
26
materials with high mobilities must be developed. As a class of ordered organic
semiconductors, liquid crystals can show high carrier mobilities, but this property has
remained largely unexplored for transistor applications82 because of the difficulties
associated with preparing electronic devices from liquid crystalline fluids. Soluble,
liquid crystalline conjugated oligomers with above-ambient glass transition
temperatures may represent an approach to obtain high mobilities in ordered solid
films across large areas in solution-processed OFETs, and are further expected to
provide a means to achieve anisotropic charge transport.
In my thesis, therefore, the following objectives were targeted through
materials characterization and organic electronic device design and experimentation:
(1) To elucidate the relationships between molecular structure and
thermotropic and optical properties of blue light-emitting, monodisperse glassy
nematic conjugated oligomers comprised of fluorene units, with an emphasis on the
thermal transition temperatures, anisotropic light absorption and emission, and
photoluminescence quantum yield.
(2) To employ the same materials in linearly polarized blue OLEDs, with
device experiments designed to probe the effects of the molecular structure of the
emissive material and the device configuration on OLED performance. Specifically,
the film thickness and the electron transport layer will be investigated as parameters
for optimization.
27
(3) To incorporate glassy-nematic conjugated oligomers and polymer
analogues in organic transistors to uncover the effects of chain length, monomer
sequence, injection barrier, and molecular aggregation on anisotropic charge carrier
mobilities.
(4) To provide new insight into the factors influencing the OLED lifetime for
blue light-emitting materials through synthesis of model compounds, device
fabrication, OLED characterization, and the measurement of charge carrier transport
properties with the transient OLED technique.
28
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38
Chapter 2
Thermotropic and Optical Properties of Monodisperse Glassy-Nematic
Oligo(fluorene)s for Linearly Polarized Blue Emission
1. INTRODUCTION
Solution-processable organic materials possessing π-conjugation have been
intensively pursued for photonic and electronic devices in recent years1,2 because of
their interesting properties and the potential for low-cost, large area fabrication. In
many cases, controlling molecular orientation and solid-state morphology through
materials design or processing is important in facilitating scientific or technological
advances. For example, uniaxial orientation of π-conjugated systems and the ability
to obtain a preferred morphology across a large area is crucial in the fabrication of
polarized organic light-emitting diodes3, 4 and organic field-effect transistors5, 6,
respectively. Several approaches to achieve uniaxial molecular alignment of
conjugated molecules have been reported, including physical stretching or rubbing,
Langmuir-Blodgett film preparation, the friction transfer technique, and self-assembly
mediated by liquid crystalline mesomorphism7-14. The high degrees of achievable
anisotropy and the relative ease of implementation suggest that uniaxial alignment of
nematic liquid crystals is the most promising method reported to date. Most
thermotropic liquid crystals tend to crystallize upon cooling from the mesomorphic
39
melt, resulting in grain boundaries that can scatter light or limit charge transport in
optical or electronic devices. For practical utility, π-conjugated nematic liquid
crystals that can maintain the molecular alignment characteristic of nematic fluids in
the solid state without crystallization would be very desirable. Some conjugated
polymers display a thermotropic nematic mesophase, for which alignment can be
preserved in the solid state by rapid cooling to suppress crystallization or by slow
cooling through the glass transition temperature (Tg). In particular, poly(fluorene)s
are a very promising class of liquid crystalline, main chain conjugated polymers for
organic electronic devices because of their photophysical and electronic properties3b,
4a,b, 12. Optimization of the mesomorphic properties through variation in aliphatic
pendant structure has been attempted to improve thermal stability and maximize
anisotropy4a,b, 12. However, for developing structure-property relationships,
conjugated oligomers possessing well-defined chain lengths and monomer sequences
are considered to be superior to polymers13.
In addition to the salient features mentioned above, monodisperse conjugated
oligomers are expected to show a higher degree of chemical purity than polymers
because of the applicability of column chromatography, recrystallization, or gradient
sublimation as purification techniques. For electronic applications, purity is critical
as traces of impurities could trap charges or quench excitons, potentially limiting
device performance. Through proper functionalization of conjugated backbones with
alkyl, alkoxy, or other substituents, solubility in common organic solvents can be
maintained, facilitating solution processing. These features enable investigation of
structure-property relationships which are fundamentally useful in materials
40
development and practically important in the design and interpretation of device
experiments. However, oligomers are less likely to undergo glass transition than
polymers, and glassy oligomeric films are generally more likely to suffer from
thermally activated crystallization than their polymeric counterparts. Many small
molecular weight organic materials that can vitrify possess a low glass transition
temperature, hindering applications in electronic devices. The preparation of
conjugated systems with high glass transition temperatures has received a great deal
of attention aimed at overcoming this limitation16, 17. However, there have been few
reports of conjugated oligomers possessing a calamitic liquid crystalline mesophase18-
25, and fewer still that display an above-ambient glass transition temperature26, which
is a potential avenue to explore for the preparation of macroscopically ordered solid
films. Photo-polymerization of π-conjugated nematic fluids has also been employed
to preserve macroscopic alignment facilitated by liquid crystalline self-assembly14, 27,
28. While this approach enables multi-layer fabrication and lithographic patterning28,
exposure to UV light for polymerization as well as the presence of initiators and/or
inhibitors could result in deleterious effects on material properties and device
performance.
Monodisperse π-conjugated oligo(fluorene)s with chiral alkyl chains
possessing high glass transition temperatures and broad mesomorphic temperature
ranges have been demonstrated in self-assembled, helically ordered chiral-nematic
films that exhibit circularly polarized light emission 29. A similar strategy employing
racemic alkyl chains was expected to result in nematic liquid crystalline conjugated
oligomers capable of preserving macroscopic uniaxial alignment by cooling through
41
the glass transition without encountering crystallization, resulting in linearly polarized
fluorescence. The resulting novel materials were characterized in this Chapter with
the following objectives: (1) to investigate the effects of chain length and pendant
structure on thermotropic properties and optical anisotropy in the forms of linear
dichroism, optical birefringence, and linearly polarized photoluminescence; and (2) to
determine photoluminescence quantum yield, and temporal stability of morphology,
emission color, and polarization ratio of aligned thin films.
2. EXPERIMENTAL
Materials System and Morphology
Charts 2.1 and 2.2 depict the molecular structures of the monodisperse
oligo(fluorene)s used in this study. Material synthesis, purification, structure
validation, and chemical purity assessment have been reported elsewhere30.
Morphology and the nature of the phase transition were characterized with hot stage
polarizing optical microscopy (DMLM, Leica, FP90 central processor and FP82 hot
stage, Mettler Toledo). Thermal transition temperatures were determined by
differential scanning calorimetry (Perkin- Elmer DSC-7) with a continuous N2 purge
at 20 mL/min. Samples were preheated to 360 oC except for F(Pr)5F(MB)2, which
was preheated to 375 oC, followed by quenching at – 200 oC/minute to – 30 oC before
collecting the second heating scans at 20 oC/minute for which the phase transition
temperatures are reported. The second heating and cooling DSC thermograms are
compiled in Appendix 1. Samples for polarizing optical microscopy (POM) were
42
heated to appropriate temperatures to observe clearing points below 375 oC, the
temperature limitation of the hot stage.
Chart 2.1: Molecular structures of monodisperse oligo(fluorene)s comprising five or
seven fluorene units.
F(Pe)5
F(Pr)5
F(MB)5
F(Pr)4F(EH)1
F(MB)4F(Pr)1
F(MB)7
F(Pr)5F(MB)2
F(MB)6F(EH)1
F(Pr)5F(EH)2
43
Chart 2.2: Molecular structures of monodisperse oligo(fluorene)s comprising nine or
twelve fluorene units.
F(MB)9
F(MB)8F(EH)1
F(MB)7F(EH)2
F(MB)7F(DMO)2
F(MB)10F(EH)2
Absorption and Fluorescence Spectroscopy of Dilute Solutions
Oligo(fluorene)s were dissolved in UV-grade chloroform at a concentration of
10-5 M of fluorene units. For co-oligomers comprised of fluorene units with varying
alkyl chains on a single molecule, an average molecular weight of the repeat unit was
used to determine the solution concentration. Absorption spectra were collected with
a UV-Vis-NIR diode array spectrophotometer (HP 8453E). Fluorescence spectra
were collected with a spectrofluorimeter (Quanta Master C60SE, Photon Technology
44
International). The excitation wavelength was 370 nm, with emission collected in a
90 o orientation.
Substrate Preparation
Optically flat fused silica substrates (25.4 mm diameter × 3 mm thick,
transparent to 200 nm, Esco Products) were scrubbed with 0.05 micron alumina
micropolish (Gamma Micropolish II, Buehler), rinsed with de-ionized water,
ultrasonicated in a dilute aqueous detergent (Isoclean, Cole Parmer) at 60 oC for one
hour, thoroughly rinsed with NanoPure water, and finally dried overnight in a laminar
flow hood prior to use. Clean substrates were spin-coated with a thin film of a
commercial alignment layer (Nissan SUNEVER grade 610 polyimide varnish, Nissan
Chemical Industries, Japan) following a modified version of previously reported
procedures31. The concentrated polyimide varnish was diluted 8 times by mass with a
mixed solvent comprised of 20/80 v/v butyl cellosolve/ethyl cellosolve prior to spin
coating at 3500 RPM for 2 minutes. The solution was applied using a filtered syringe
(0.45 μm Teflon filters, Millipore). After spin coating, the substrates were soft-baked
at 80 oC for 10 minutes followed by hard curing at 250 oC for 1 hour. To facilitate
liquid crystal alignment, the substrates were uniaxially rubbed (buffed) with a home-
built buffing machine by passing the coated substrate 6 times beneath a velvet cloth
mounted on a drum, which rotated at ~ 1500 rpm. The height of the drum, and hence
the buffing strength, was controlled with a micrometer. The buffing direction defined
the so-called nematic director by imposing a homogeneous boundary condition that
was followed by the liquid crystal molecules32.
45
The buffing strength, determined by the micrometer which defined the height
of the rotating drum, was found to influence the anisotropy obtained for aligned liquid
crystalline thin films. Therefore, care was taken to ensure that the polyimide films
exhibited consistent alignment performance to enable a comparison of material
properties without systematic errors induced by variation in the quality of the
alignment layer. A buffing optimization procedure was developed to accomplish this,
which involved uniaxial rubbing of alignment layers at various micrometer settings
followed by preparation of monodomain glassy nematic films of F(Pr)5F(MB)2, as
described in the following section. Varying the buffing strength allowed the linearly
polarized fluorescence dichroism to be maximized. Subsequent alignment layers
were buffed at the micrometer setting that gave the best results for F(Pr)5F(MB)2.
Linearly polarized photoluminescence spectra of films prepared for buffing
optimization are presented in Figure 2.1, demonstrating the sensitivity of the liquid
crystal alignment to the buffing strength. Through this approach, uniaxial rubbing
can be performed reproducibly and the anisotropic optical properties of
oligo(fluorene)s can be compared across batches of polyimide alignment layers
without introducing systematic errors caused by variation in the alignment ability of
the rubbed polyimide. All substrate preparation was performed in a clean room rated
class 10,000 with fume hoods rated as class 100 environments.
46
Figure 2.1: Typical results for polyimide alignment layer buffing optimization
performed with a 0.8 wt % solution of F(Pr)5F(MB)2. Rem refers to the dichroic
ratio at the first emission maximum (λ = 420 nm). The micrometer setting that
controls the height of the rotating drum is indicated. Parallel and Perpendicular refer
to the orientation of the first polarizer relative to the nematic director.
0.0
1.0
2.0
3.0
4.0
400 500 600
ParallelPerpendicular
ParallelPerpendicular
Em
issi
on In
tens
ity, a
.u.
Wavelength, nm
3.97
4.00
Rem = 16.7
Rem = 13.9
Preparation and Characterization of Thin Films
Thin oligo(fluorene) films were prepared in the clean room by spin coating
from filtered 1 wt % solutions in UV-grade chloroform at a spin rate of 5000 RPM
followed by overnight drying in vacuo. Thermal annealing was performed in an
Argon-purged glassware at temperatures 5 to 20 degrees above the glass transition for
30 minutes or less. Films for electron diffraction (JEM 2000 EX, JEOL USA) were
prepared following the same procedures except single crystalline sodium chloride
substrates (13 mm diameter × 2 mm thickness, International Crystal Laboratories)
without alignment coatings were used. These films were floated off in a trough of de-
ionized water for mounting on copper grids. Electron diffraction measurements were
performed by Dr. Brian McIntyre (University of Rochester, Institute of Optics). The
47
absence of nematic textures or defects under 500 times magnification in POM
revealed that annealed thin films on optimized alignment layers were monodomain.
Absorption and linear dichroism of monodomain films were evaluated using a
UV-Vis-NIR spectrophotometer (Lambda-900, Perkin-Elmer) and linear polarizers
(HNP’B, Polaroid). Fluorescence spectra, with and without polarization analysis,
were collected using the spectrofluorimeter with a liquid light guide directing a 370
nm unpolarized excitation beam onto the samples in a straight-through geometry.
Emitted light normal to the film surface was detected. The linearly polarized
photoluminescence (LPPL) was characterized by controlling the orientation of the
nematic director of the film relative to a pair of linear polarizers placed in front of the
detector. The film was oriented with its nematic director vertical or horizontal. For
each film orientation, the emission spectrum was collected twice, rotating the
transmission axis of the first linear polarizer by 90 o such that polarized emission
parallel and perpendicular to the director was detected. The results from the two film
orientations were averaged to minimize error. Experimental error due to polarization
differences in the detector was further reduced by inserting a second linear polarizer
with its transmission axis at 45° between the first polarizer and the detector for all
measurements.
Optical properties of thin films were determined with transmission mode
spectroscopic ellipsometry (V-VASE, J. A. Woollam Co.). This non-destructive
optical technique allowed the determination of the real and imaginary parts of both
the ordinary and extraordinary refractive indices along with the film thickness
48
following literature procedures33. The refractive index is required for the
determination of thin film quantum yield as described below34.
Determination of the Fluorescence Quantum Yield
The fluorescence quantum yield Φ was measured following a previously
reported procedure35. A relative method employing a standard comprised of a
poly(methyl methacrylate) (PMMA, Polysciences Inc., weight average molecular
weight = 75,000 g/mole) film doped with 10-2 M diphenyl anthracene (Acros
Organics, 99 %) that had been repeatedly re-crystallized from xylenes was used. The
standard was prepared by spin-coating on a clean fused silica substrate followed by
overnight drying in vacuo. A secondary standard consisting of an identical PMMA
film lightly doped with anthracene (Aldrich, 99 %) that had been repeatedly re-
crystallized from ethanol was prepared to validate the experimental method. The
UV-Vis absorption and integrated fluorescence spectra with a normal incidence,
unpolarized excitation beam at 355 nm and detection at 60 degrees off normal were
collected for both the primary and secondary standards. The primary standard was
assigned a quantum yield value of 0.8336 and was considered as a reference (subscript
r) for the determination of Φ for the secondary standard sample (subscript s) using the
following equation35:
2
2
)101()101(
rrA
ssA
rs nBnB
s
r
−
−
−−
Φ=Φ (2-1)
In Equation 2-1, A represents absorbance at 355 nm, while B is defined as the
integrated photoluminescence spectrum from 360 to 600 nm. The factor n2 is
49
included to correct for the refractive index difference between the sample and
reference films as defined in Equation 2-2.
( ) ( )( ) λλ
λλλdI
dnIn∫
∫≡
22 * (2-2)
In this equation, I and n represent the photoluminescence spectrum and the refractive
index, respectively, with limits of integration identical to those in Equation 2-1. The
secondary standard showed a quantum yield of 0.28 ± 0.03, within experimental
uncertainty of the literature value of 0.27 for anthracene in benzene or ethanol37,
validating the method. Aligned thin films prepared as described above were
measured and the quantum yields were calculated in a similar fashion. An average of
two integrated photoluminescence spectra collected with the nematic director oriented
vertically and horizontally was employed in the quantum yield calculation for
macroscopically ordered oligo(fluorene) films.
3. RESULTS AND DISCUSSION
Thermotropic and Morphological Properties
Thermotropic properties, characterized by differential scanning calorimetry
(DSC) and polarizing optical microscopy (POM) with heating rates of 20 oC/min, are
summarized in Table 2.130. For DSC experiments, samples were pre-heated to 360 oC
except for F(Pr)5F(MB)2, which was heated to 375 oC, followed by quenching to –
30 oC before collecting the second heating scan for which the transition temperatures
are reported.
50
Table 2.1: Phase transition temperatures (Tg and Tc), absorption and emission dichroic
ratios, orientational order parameters, and photoluminescence quantum yields of
thermally annealed monodomain spin cast films.
Compound Tg, oC [a] Tc, oC [a] τ, nm [b] Rabs [c] Sabs
[d] Rem[c]
Pentamers F(Pr)5 [f] N/A N/A N/A N/A N/A N/A F(Pe)5 [g] 72 N/A N/A N/A N/A N/A F(M B)5 93 166 85 9.8 ± 0.5 0.75 ± 0.1 10.8 ± 0.5 F(Pr)4F(EH)1 [h] 114 178 N/A N/A N/A N/A F(M B)4F(Pr)1 83 123 88 7.0 ± 0.4 0.67 ± 0.01 6.9 ± 0.5 Heptamers F(Pr)5F(M B)2 149 364 81 15.5 ± 0.7 0.83 ± 0.01 14.1 ± 0.2 F(Pr)5F(EH)2 116 319 78 11.0 ± 0.9 0.77 ± 0.03 10.5 ± 0.7 F(M B)7 [i] 110 320 N/A N/A N/A N/A F(M B)6F(EH)1 101 315 85 13.7 ± 0.4 0.81 ± 0.01 13.1 ± 0.7 Nonamers F(M B)9 [j] N/A N/A N/A N/A N/A N/A F(M B)8F(EH)1 112 > 375 88 14.8 ± 0.3 0.82 ± 0.01 13.2 ± 1.5 F(M B)7F(EH)2 108 > 375 92 15.1 ± 0.6 0.82 ± 0.01 15.8 ± 0.8 F(M B)7F(DM O)2 82 273 87 13.1 ± 0.6 0.80 ± 0.01 12.2 ± 0.4 Dodecamer F(M B)10F(EH)2 123 > 375 90 19.9 ± 0.5 0.86 ± 0.01 17.1 ± 0.9
[a] Transition temperatures gathered from the heating scans at 20 oC/min of samples preheated to 360 oC or 375 oC and quenched to –30 oC. Tg represents the inflection point across the glass transition, while Tc represents nematic to isotropic transition as identified by POM. [b] Thickness of spin cast films annealed to final values of absorption and emission dichroism, determined by spectroscopic ellipsometry. [c] Rabs and Rem evaluated at the absorption maximum and first emission maximum, respectively. [d] Orientational order parameter, Sabs = (Rabs – 1)/(Rabs + 2). [e] Thin film ΦPL determined as reported previously34, 35. [f] Heating and cooling scans showed a melting point at 318 oC and crystallization at 266 oC, respectively, with no mesomorphism under POM. [g] Heating and cooling scans showed a glass transition at 72 and 62 oC, respectively, without any mesomorphism. [h] Thin films showed crystal nuclei in addition to mesophase formation upon annealing. [i] Smectic A phase, which prevented the preparation of monodomain glassy-nematic films, was observed between Tg and 157 oC, where a transition to nematic occurred. [j] Heating scan showed a crystal melting to a nematic fluid at 260 oC, while cooling scans showed crystallization from the nematic melt at 160 oC; no glass transition was detected in either scan.
Five fluorene units was the minimum oligomer length for which
mesomorphism was observed in this study, while the aliphatic substituents at the 9, 9-
51
position were significant in determining the morphology and the nature of the phase
transition. For example, neither F(Pr)5 nor F(Pe)5 show liquid crystalline
mesomorphism, with F(Pr)5 displaying a crystalline melting point, Tm, at 318 oC and
F(Pe)5 possessing a glass transition temperature, Tg, of 72 oC. The longer pentyl
pendants in F(Pe)5 apparently suppressed crystallization in favor of vitrification. A
pentamer with branched 2-methylbutyl pendants, F(MB)5, showed a nematic
mesophase between Tg at 93 oC and the clearing point, Tc, at 166 oC, excellent
properties for the realization of ordered solid glassy-nematic films. In contrast to the
nematic mesophase and high glass transition temperature observed for F(MB)5, a
penta(fluorene) with 2-ethylhexyl pendants at the 9, 9-position reported elsewhere
showed a smectic mesophase with a low glass transition temperature of only 28 oC26c.
Instead of restricting this study to homo-oligomers, co-oligomers in which fluorene
units along the conjugated backbone are functionalized with different alkyl chains
were also evaluated. As a first example, insertion of 2-ethylhexyl pendants on the
central fluorene unit of F(Pr)5 resulted in F(Pr)4F(EH)1, a well-behaved glassy-
nematic material with a Tg of 114 oC and a Tc of 178 oC, demonstrating the utility of
the co-oligomer concept to produce stable glassy-nematic morphologies. On the
other hand, replacing 2-methylbutyl side chains on the central fluorene unit of
F(MB)5 with n-propyl units in F(MB)4F(Pr)1 decreased Tg and Tc by 10 and 43 oC,
respectively. These observations of the effects of pendant structure on morphology
and the nature of the phase transitions in pentamers aided in the design of higher
oligomers with an objective of realizing high glass transition temperatures, broad
nematic mesophase ranges, and good solubility to facilitate the preparation of ordered
52
solid glassy-nematic films by spin coating and thermal annealing, followed by
cooling across Tg without encountering crystallization. Furthermore, higher
oligomers are of interest to reveal the effects of chain length in the emerging
structure-property relationships.
All heptamers in Table 2.1 exhibit nematic mesophases with temperature
ranges of 200 oC or more. Shorter aliphatic pendants apparently contributed to higher
transition temperatures, as demonstrated by F(Pr)5F(MB)2 with a Tg of 149 oC and a
Tc of 364 oC. The influence of pendant structure on the phase behavior is apparent
through comparison of F(MB)7 to F(MB)6F(EH)1, which only differ in the
substituents on the central fluorene unit. The homo-heptamer F(MB)7 displays a
smectic mesophase above Tg, with a transition to nematic at 157 oC. This behavior is
similar to a hepta(fluorene) F(EH)7 reported elsewhere25c, with the less bulky 2-
methylbutyl pendants in F(MB)7 resulting in a substantially higher Tg and Tc of 110
and 320 oC vs. 42 and 246 oC, respectively. The more complicated mesophase
behavior for F(MB)7 prevented the preparation of an ordered solid film by annealing
at temperatures 5 to 20 oC above the glass transition temperature presumably because
of the high viscosity of the smectic phase. In contrast, the co-oligomer
F(MB)6F(EH)1 showed only nematic mesomorphism between Tg and Tc, facilitating
the preparation of an ordered solid glassy-nematic film. The results suggest that
pendant mixing in longer co-oligomers is important to realize nematic, rather than
smectic, mesomorphism in higher oligo(fluorene)s. The variation in pendant
structure appears to effectively disrupt the formation of smectic layering, allowing the
nematic mesophase to persist across a broad temperature range.
53
In the higher oligomers shown in Chart 2.2, co-oligomers with longer,
branched aliphatic side chains are more likely to show stable glassy-nematic behavior
with good solubility and an absence of thermally activated crystallization. For
example, homo-nonamer F(MB)9 showed a melting point Tm and a crystallization
temperature Tk of 260 oC and 160 oC in DSC heating and cooling scans. No glass
transition or clearing point could be detected for this nematic material. As in the
heptamer F(MB)7, the nematic phase for the nonamer F(MB)9 apparently cannot
bypass a more ordered arrangement upon cooling, preventing the preparation of
monodomain glassy-nematic films. In contrast, co-nonamers F(MB)8F(EH)1 and
F(MB)7F(EH)2 exhibited high glass transition temperatures and exceptionally stable
nematic mesophases with clearing points above 375 oC. A co-dodecamer,
F(MB)10F(EH)2, showed phase behavior similar to the co-nonamers but with a
slightly elevated Tg of 123 oC because of the longer conjugated backbone. The
mesophase behavior was confirmed by the typical nematic threaded textures observed
in hot stage polarizing optical microscopy of micrometer-thick films prepared
between a microscope slide and a cover slip without surface treatment, as shown in
Figure 2.2 for two representative oligomers.
54
Figure 2.2: Nematic threaded textures observed in hot stage polarizing optical
microscopy (POM) for micrometer-thick films prepared between a microscope slide
and a cover slip without surface treatment. a) F(Pr)5F(MB)2 and b)
F(MB)7F(EH)2.
a) b)
Thin Film Preparation and Optical Characterization
Thin films were prepared by spin coating from filtered 1 wt % solutions in
UV-grade chloroform on fused silica substrates that had been previously coated with
an optimized, uniaxially rubbed polyimide alignment layer. Films were vacuum-
dried overnight at room temperature before thermal annealing at a temperature 5 to 20
oC above Tg for 15 to 30 minutes in an Argon-purged glass vessel. Macroscopically
ordered solid films were obtained by cooling to below Tg, preserving the uniaxial
nematic alignment. UV-Vis linear dichroism and linearly polarized
photoluminescence spectra were collected for the aligned films (see Experimental).
Absorption and emission dichroic ratios were defined as Rabs ≡ A||/A⊥ and Rem ≡ I||/I⊥
in which A and I are absorbance and emission intensity at λmax, respectively, and
subscripts || and ⊥ specify polarization directions parallel and perpendicular to the
nematic director defined by the buffing direction. The glassy-nematic films were
55
characterized as monodomain due to the absence of nematic textures or defects in
polarizing optical microscopy under 500 × magnification. Furthermore, annealed
films were evaluated as non-crystalline by electron diffraction, supported by the
typical diffraction patterns presented in Figure 2.3. It is noted that the glassy-nematic
morphology was stable in thin films stored at room temperature for more than 2 years
without encountering crystallization.
Figure 2.3: Electron diffraction patterns of annealed glassy-nematic solid films
displaying diffuse rings characteristic of a non-crystalline morphology. a)
F(Pr)5F(MB)2 and b) F(MB)7F(EH)2.
a) b)
Transmission mode spectroscopic ellipsometry33 was employed to determine
the film thickness, ordinary and extraordinary refractive indices no and ne, and
anisotropic absorption coefficients α|| and α⊥. Films were modeled as optically
biaxial, with results indicating that the refractive index normal to the film surface was
identical to no, typifying uniaxial alignment. The extracted refractive indices and
anisotropic absorption coefficients for F(Pr)5F(MB)2 are shown in Figure 2.4. The
56
refractive indices from Figure 2.4a were required for the calculation of the
photoluminescence quantum yield of ordered thin films following previously reported
procedures (see Experimental)34, 35. In Figure 2.4b, the anisotropic absorption
coefficients determined by ellipsometry are compared to those independently
measured through linear dichroism with a spectrophotometer. The good agreement
between the two sets of data provided confidence in the ellipsometric approach.
Figure 2.5 compiles the optical properties for representative oligo(fluorene)s of
varying chain length. The absorption dichroism at λmax and the optical birefringence
at λ = 532 nm, where both absorption coefficients are zero, appear to increase with
the aspect ratio of the conjugated oligomer.
Figure 2.4: Ellipsometric characterization for a monodomain film of F(Pr)5F(MB)2.
a) Ordinary (no) and extraordinary (ne) refractive indices; b) Absorption coefficients
parallel and perpendicular to the nematic director defined by the buffing direction.
The absorption coefficients extracted by ellipsometric modeling agree well with those
measured independently by UV-Vis linear dichroism with a spectrophotometer.
1.0
1.5
2.0
2.5
3.0
300 400 500 600
Inde
x of
Ref
ract
ion
Wavelength, nm
ne
no
a) F(Pr)5F(MB)281 nm film
0.0
2.0
4.0
6.0
300 400 500 600
observed
observedextracted
extracted
Abs
orpt
ion
Coe
ffic
ient
, 10 5
cm
1
Wavelength, nm
α {α {
b)
⊥
||
57
Figure 2.5: Anisotropic optical properties for representative oligo(fluorene)s of
increasing chain length. ne and no and subscripts || and ⊥ are as defined for Figure
2.4. a) F(MB)5; b) F(MB)7F(EH)2; and c) F(MB)10F(EH)2.
0.0
1.0
2.0
3.0
4.0
5.0
1.0
1.5
2.0
2.5
3.0
300 350 400 450 500 550 600Abs
orpt
ion
Coe
ffic
ient
, 10 5
cm
1
Refractive Index
Wavelength (nm)
α||
α⊥
ne
no
a)
0.0
1.0
2.0
3.0
4.0
5.0
1.0
1.5
2.0
2.5
3.0
300 350 400 450 500 550 600Abs
orpt
ion
Coe
ffic
ient
, 10 5
cm
1
Refractive Index
Wavelength (nm)
α||
α⊥
ne
no
b)
0.0
1.0
2.0
3.0
4.0
5.0
1.0
1.5
2.0
2.5
3.0
300 350 400 450 500 550 600Abs
orpt
ion
Coe
ffici
ent,
10 5
cm
1
Refractive Index
Wavelength (nm)
α||
α⊥
ne
no
c)
58
Spectroscopic Characterization of Dilute Solutions and Pristine and Annealed
Films
An important feature of all oligo(fluorene)s in Charts 2.1 and 2.2 is solubility
in common organic solvents such as chloroform, methylene chloride, toluene, and
THF, which facilitated film processing by spin coating. A comparison between the
absorption and fluorescence spectra of pristine films and dilute solutions for
representative oligomers is given in Figure 2.6. The absorption spectra showed
nearly identical peak wavelengths with a slight broadening observed for pristine
films. The emission peaks for pristine films were red-shifted by approximately 10 nm
from those of the corresponding solutions, likely due to a more planar excited state
geometry and/or an increase in the dielectric constant of the medium. The different
intensities of the 0-0 and 0-1 vibronic bands of the fluorescence spectra for solutions
and films are probably caused by self-absorption in the films. Ground state
aggregation and excited state interactions were not discernible in pristine films based
on the similarities between dilute solutions’ and pristine films’ absorption and
emission spectra.
59
Figure 2.6: Normalized absorbance and photoluminescence spectra (with excitation at
370 nm) of dilute solutions in UV grade chloroform (10-5 M in fluorene units) and
pristine films for representative oligo(fluorene)s of increasing chain length. a)
F(MB)4F(Pr)1; b) F(Pr)5F(MB)2; c) F(MB)7F(DMO)2; and d)
F(MB)10F(EH)2.
0.00
0.50
1.00
300 400 500 600
Pristine FilmDilute Solution
0.0
0.5
1.0
Nor
mal
ized
Abs
orba
nce N
ormalized E
mission
Wavelength, nm
F(MB)4F(Pr)1a)
0.00
0.50
1.00
300 400 500 600
Pristine FilmDilute Solution
0.0
0.5
1.0
Nor
mal
ized
Abs
orba
nce N
ormalized E
mission
Wavelength, nm
F(Pr)5F(MB)2b)
0.0
0.5
1.0
300 400 500 600
Pristine FilmDilute Solution
0.0
0.5
1.0
Nor
mal
ized
Abs
orba
nce N
ormalized E
mission
Wavelength, nm
F(MB)10F(EH)2d)
0.0
0.5
1.0
300 400 500 600
Pristine FilmDilute Solution
0.0
0.5
1.0
Nor
mal
ized
Abs
orba
nce N
ormalized E
mission
Wavelength, nm
F(MB)7F(DMO)2c)
The anisotropic absorption and emission spectra for annealed films of
representative oligo(fluorene)s are compiled in Figure 2.7. The orientational order
parameter based on absorption, Sabs = (Rabs – 1)/(Rabs + 2), emission dichroic ratio
(Rem = I||/I⊥), and film thickness are also included. Both Sabs and Rem, evaluated at
λmax, increased with an increasing aspect ratio for the given chain lengths and pendant
60
structures, in agreement with the optical dichroism and birefringence results from
spectroscopic ellipsometry. A slight blue-shift for absorption perpendicular to the
chain alignment relative to that parallel to the nematic director was evident in
monodomain films. Furthermore, the fluorescence spectra parallel and perpendicular
to the chain alignment showed substantially different vibronic progression. Both of
these effects may be related to the presence of poorly aligned transition dipole
moments within the distribution about the nematic director38. Despite this, the polar
plot shown in Figure 2.8 for F(Pr)5F(MB)2 confirms that the absorption transition
dipole moment is approximately parallel to the nematic director as defined by the
buffing direction. The fact that the absorption and emission dichroic ratios are
comparable verifies that the transition dipole moments are approximately co-linear.
Results for all monodomain films prepared for this study, including
photoluminescence quantum yields and experimental uncertainties, are compiled in
Table 2.130. Co-oligomers with branched aliphatic pendants were more likely to
show the desired stable glassy-nematic morphology than homo-oligomers. The
insertion of bulky pendants, however, tended to decrease Tg and Tc, effects that could
be partially counteracted through increased oligomer length. The high order
parameters of 0.67 to 0.86 achieved through thermal annealing under mild conditions
of temperature and time demonstrated the ease of processing the glassy-nematic
oligo(fluorene) films to monodomain. The observed emission dichroism of 17.1 ±
0.9 for F(MB)10F(EH)2 compares favorably to photoluminescence polarization
ratios achieved for poly(fluorene) films12, 13, 39. The thin films emit blue light with
61
high quantum yields ranging from 43 to 60 %, with no obvious relationship to the
oligomer length or the pendant structure.
Figure 2.7: Linear dichroism and linearly polarized photoluminescence spectra (with
unpolarized excitation at 370 nm) of 80 to 90 nm thick, thermally annealed
monodomain films. Subscricpts || and ⊥ specify polarization directions parallel and
perpendicular to the nematic director defined by the buffing direction (See
Experimental). a) F(MB)5; b) F(Pr)5F(MB)2; c) F(MB)7F(EH)2; d)
F(MB)10F(EH)2.
0.0
0.5
1.0
1.5
300 400 500 6000.0
1.0
2.0
3.0
Abs
orba
nce
Em
ission Intensity, a.u.
Wavelength, nm
I
IA
A
= 0.73
F(MB)585 nm film = 11.1 at 415 nm
a)
Rem
Sabs
⊥
⊥
||
||
0.0
0.5
1.0
1.5
2.0
300 400 500 6000.0
1.0
2.0A
bsor
banc
e
Em
ission Intensity, a.u.
Wavelength, nm
I
I
A
A
= 0.82
F(Pr)5F(MB)281 nm film = 14.2 at 419 nm
b)
Rem
Sabs
⊥
⊥
||
||
0.0
0.5
1.0
1.5
300 400 500 6000.0
1.0
2.0
3.0
Abs
orba
nce
Em
ission Intensity, a.u.
Wavelength, nm
I
IA
A
= 0.81
F(MB)7F(EH)292 nm film = 16.1 at 420 nm
c)
Rem
Sabs
⊥⊥
||
||
0.0
0.5
1.0
1.5
2.0
300 400 500 6000.0
1.0
2.0
3.0
Abs
orba
nce
Em
ission Intensity, a.u.
Wavelength, nm
I
IA
A
= 0.84
F(MB)10F(EH)290 nm film = 18.1 at 422 nm
d)R
em
Sabs
⊥
⊥
||
||
62
Figure 2.8: UV-Vis absorption polar plot for a ~ 70 nm thick, monodomain
F(Pr)5F(MB)2 film.
0.00.4
0.81.2
1.6
0
30
6090
120
210
240270
300
330Abs
orba
nce
@ λ
max
Thermal Stability of Emission Color and Polarization Properties
Despite the favorable optical and electronic properties of blue light-emitting
poly(fluorene)s, the stability of the emission color is typically poor when thin films
are exposed to heat, UV light, or air. The instability, manifested as an appreciable
emission peak in the green to yellow, was attributed to excimer formation or
oxidation at polymer chain ends in early work40, 41. More recently, the formation of
emissive keto-defects along the polymer backbone has gained support as the primary
mechanism for spectral instability42. To address this potential problem with
oligo(fluorene)s, long-term thermal annealing of thin films was performed. This
experiment was also designed to address the issue of Joule heating in operating
OLEDs43, where temperatures up to 90 oC can be realized under continual forward
bias44. The stability of morphology, emission color, and dichroic ratio were evaluated
by annealing a 70 nm film of F(MB)8F(EH)1 at a temperature of 120 oC, 5 oC above
its Tg, for up to 96 hours. The results are presented in Figure 2.9. Pronounced
63
hypochromism through increased π-orbital interactions45 was present regardless of
the annealing time, as 30 minutes and 96 hours are compared in Figure 2.9a. The
absorbance at 370 nm for films annealed for 30 minutes and 96 hours diminished by
32 and 37 %, respectively, which contributed to the decrease in integrated
photoluminescence shown in Figure 2.9b. The larger decrease for the
photoluminescence, 53 % in both cases, suggests that new non-radiative decay
pathways were formed during thermal annealing, but also indicates that the emission
intensity was not continually degraded during prolonged heating above Tg. The
normalized emission spectra for the pristine and 96 hour annealed film presented in
Figure 2.9c demonstrate the temporal stability of emissive color for a film subjected
to prolonged heating. As shown in Figure 2.9d, annealing above Tg results in an
increase of Rem of the higher energy fluorescence peak from 1.1 for pristine films to
14.4 in the 96 hour annealed film. The dichroic ratio of the 96 hour annealed film is
within experimental uncertainty of the value for F(MB)8F(EH)1 films annealed for
15-30 minute at the same temperature (see Table 2.1).
64
Figure 2.9: Effects of prolonged thermal annealing of a 70 nm film of
F(MB)8F(EH)1 at a temperature of 120 oC (i.e. 5 oC above Tg) under Argon. a)
Increased π-orbital interactions due to molecular alignment and close packing
manifested by hypochromism; b) Accompanying reduction in integrated
photoluminescence intensity; c) Normalized photoluminescence spectra of pristine
and 96 hour annealed films, displaying the stability of the blue emission color; and d)
Increase in emission dichroic ratio from 1.1 to 14.4 upon annealing for 96 hours.
Subscripts || and ⊥ are as defined in Figure 2.7.
0.0
1.0
2.0
3.0
400 500 600
PristineAnnealed 30 minAnnealed 96 hr
Emis
sion
Int
ensit
y, a
.u.
Wavelength, nm
b)
0.0
1.0
2.0
3.0
4.0
400 500 600
Pristine
Pristine
Annealed
Annealed
Em
issio
n In
tens
ity, a
.u.
Wavelength, nm
d)
I|| {
{I⊥
0.0
0.3
0.6
0.9
1.2
400 500 600
PristineAnnealed96 hours
Nor
mal
ized
Em
issi
on
Wavelength, nm
c)
0.0
0.2
0.4
0.6
0.8
300 400 500 600
Pristine, Film 1Pristine, Film 2Annealed 30 minAnnealed 96 hr
Abs
orba
nce
Wavelength, nm
a)
4. SUMMARY
Based on the ease of formation of monodomain glassy-nematic thin films and
a buffing procedure that ensured both optimal and reproducible performance of
65
polyimide alignment layers, relationships between molecular structure and
thermotropic and optical properties of a series of monodisperse oligo(fluorene)s were
elucidated. Differential scanning calorimetry and electron diffraction were used to
evaluate morphology in the bulk and in thin films, while the nature of the phase
transitions were confirmed with polarizing optical microscopy. UV-Vis spectroscopy
of solutions and thin films, linear dichroism, linearly polarized photoluminescence,
and spectroscopic ellipsometry were employed to determine photophysical and
optical properties. The key experimental results are summarized as follows:
(1) Co-oligomers and branched aliphatic pendants improve morphological
stability and promote desirable mesophase behavior compared to homo-oligomers
and linear alkyl chains, respectively. The glass transition temperature appears to
depend on both oligomer length and pendant structure, while the clearing point
generally increases with the number of fluorene units with the pendant playing a less
prominent role.
(2) The variation in oligomer length and pendant structure resulted in
morphologically stable glassy-nematic materials with high transition temperatures,
with Tg up to 149 oC and Tc beyond 375 oC. These ideal thermotropic properties,
along with good solubility in common organic solvents, facilitated the preparation of
monodomain glassy-nematic thin films by spin-coating and thermal annealing at
temperatures 5 to 20 oC above Tg for up to 30 minutes, followed by cooling to room
temperature. The mild annealing conditions resulting in high order parameters and
fluorescence dichroic ratios indicate the ease of processing thin films to monodomain.
66
(3) Photoluminescence quantum yields for ordered thin films ranged from 43
to 60 %, showing the potential of oligo(fluorene)s for use in polarized OLEDs.
Optical birefringence, absorption anisotropy, and emission dichroism appeared to
increase with increasing molecular aspect ratio. The emission dichroic ratio of more
than 17 achieved for F(MB)10F(EH)2 compares well with the best
photoluminescence results achieved for aligned poly(fluorene) films prepared by spin
coating and annealing under more demanding conditions of temperature and time.
Annealed thin films were shown to be non-crystalline by electron diffraction.
Furthermore, monodomain glassy-nematic thin films do not crystallize upon storage
at room temperature for more than two years.
(4) Annealing a thin film of F(MB)8F(EH)1 above Tg for 96 hours did not
result in crystallization or pronounced changes in the emission color. The observed
loss in integrated photoluminescence intensity was attributed to a decrease in the
absorbed photoexcitation by hypochromism along with the formation of new non-
radiative relaxation pathways. Thin films that were annealed for 15 to 30 minutes
and 96 hours exhibited identical emission dichroic ratios within experimental
uncertainty.
67
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71
Chapter 3
Strongly Polarized, Efficient Deep Blue Electroluminescence Based on
Monodisperse Glassy-Nematic Oligo(fluorene)s
1. INTRODUCTION
Electroluminescence (EL) from organic materials was first observed by
Bernanose et al in the early 1950s1. The drive voltage required to produce visible
light emission resulted in low power efficiency, limiting research interest in this area
until Tang and Van Slyke showed that the voltage could be reduced and the emission
efficiency improved dramatically with amorphous thin films of vacuum-evaporable
organic materials2. The realization of electroluminescence from conjugated polymer
thin films by Burroughes et al3 was another landmark achievement in the field of
organic electronics. This work generated particular interest because of the possibility
of preparing semiconductor devices by solution processing, promising for large area
applications at low cost. In the past 20 years, organic light-emitting diodes (OLEDs)
have become the focus of intense academic and industrial research aimed towards the
development of new flat panel display and solid-state lighting technologies. Progress
in materials and device science has resulted in OLEDs spanning the visible spectrum
with improved efficiencies and device lifetimes4-7. Both fluorescent and
72
phosphorescent OLEDs emitting white light with very high efficiencies have also
been considered for displays containing color filters or as large-area light sources8-10.
Despite substantial progress in OLEDs, the dominant flat panel display
technology remains the liquid crystal display (LCD)11. Essentially, each pixel of a
transmission-mode LCD acts as an electro-optic switch, in which the light generated
by a backlight may be transmitted to the viewer or absorbed depending on its
polarization state. A backlight is required for transmission-mode LCDs because they
are not emissive by nature. Electro-optic switching in an LCD is enabled by the
optical and dielectric anisotropy of the mesomorphic fluid state12. The interaction of
the electric vector of a linearly polarized beam, entailing the inclusion of polarizing
elements in the display, with the oriented liquid crystal molecules determines if a
pixel is on or off. Furthermore, a full-color, twisted-nematic transmission mode LCD
also contains an array of absorbing color filters that result in low brightness and high
power consumption because only a fraction of the light generated by the backlight is
transmitted to the viewer. The backlight element is also bulky and adds to the
thickness and weight of the LCD. To overcome the limitations of weight, thickness,
and power consumption, OLED backlights for LCDs have been proposed13. This
combination has additional potential for transflective LCDs, which may further
reduce power consumption by operating in reflective mode under bright ambient
lighting conditions14. However, a vast majority of OLEDs emit unpolarized light due
to the random orientation of emission dipole moments in the plane of a light-emitting
organic film prepared by vacuum evaporation or spin coating. A linear polarizer
would therefore absorb at least 50 % of the OLEDs’ light output, diminishing the
73
potential reduction in power consumption over a traditional backlight. The direct
emission of linearly polarized light from an OLED could conceivably overcome this
limitation. A full-color array of linearly polarized OLED pixels with high
efficiencies and polarization ratios would make one polarizer and the color filters
redundant, substantially improving the light throughput while reducing the weight
and thickness of the LCD. In addition to the potential benefits to LCDs, linearly
polarized electroluminescence could improve OLED displays where polarizers are
often employed to enhance contrast15. Furthermore, linearly polarized OLEDs may
be applied in projection displays or stereoscopic imaging systems16, 17. The broad
array of potential benefits to display technology provides motivation for research and
development in materials and devices capable of linearly polarized light emission.
In principle, linearly polarized light emission can be realized through uniaxial
orientation of the emission transition dipole moment18. In many conjugated materials
with extended π-systems, the transition moment is approximately parallel to the
conjugated backbone. Several techniques to align the emission dipole moments of
conjugated materials in thin films have been successful in producing linearly
polarized OLEDs, including direct stretching or rubbing, vacuum evaporation on
orienting surfaces, Langmuir-Blodgett film preparation, the friction transfer
technique, and molecular self-assembly facilitated by nematic mesomorphism19-24.
With the exception of the use of nematic liquid crystalline conjugated polymers, these
approaches typically exhibited poor OLED performance, including high drive
voltages and low efficiencies. Liquid crystallinity has been the most extensively
pursued means to generate linearly polarized emission. In particular, mesomorphic
74
poly(fluorene)s have shown promise for linearly polarized OLEDs because of high
quantum yield blue emission, good charge transport properties, and favorable
morphological behavior25. For the realization of strongly polarized emission from
nematic liquid crystals, an alignment layer to facilitate monodomain uniaxial
orientation is required. Traditional alignment layers such as buffed polyimide are
electrically insulating, a drawback for polarized OLEDs where charge-transporting
ability is desired to maintain low driving voltages and high power efficiencies.
Polyimide that had been physically doped with small molecular weight hole-
transporting moieties was employed as a conductive alignment layer in early work26-
28. Devices prepared in this manner exhibited moderately polarized
electroluminescence, but suffered from low efficiencies and a tendency towards phase
separation in the alignment layer29. Similarly doped photoalignment layers have also
shown good polarization properties and higher efficiencies in polarized OLEDs
comprising aligned poly(fluorene)s as an emissive layer30. A poly(fluorene) device
including a buffed poly(p-phenylene vinylene) (PPV) alignment layer gave a high
polarization ratio of 25:1 as a result of recombination close to the rubbed interface,
but efficiencies were still low31. In typical polymer LEDs,
poly(ethylenedioxythiophene) doped with poly(styrene sulfonic acid) (PEDOT/PSS)
is used as a hole injection layer32. Uniaxially rubbed PEDOT/PSS is capable of
aligning small molecular weight liquid crystals and light-emitting side chain liquid
crystalline polymers33, 34, suggesting the applicability of PEDOT/PSS in polarized
OLEDs.
75
In the previous Chapter, the thermal, optical, and photophysical properties of
monodisperse conjugated oligomers were explored and related to the molecular
structures. Glassy-nematic oligo(fluorene)s showed high photoluminescence
quantum yields and strongly anisotropic emission in thin films aligned on insulating
polyimide alignment layers, suggesting their potential for use in polarized OLEDs35.
In this Chapter, polarized OLEDs based on thin films of oligo(fluorene)s on
uniaxially rubbed PEDOT/PSS conductive alignment layers are explored with the
following objectives:
(1) To evaluate the performance of uniaxially rubbed, optimized PEDOT/PSS
as an alignment layer for conjugated oligomers through film preparation, thermal
annealing, and characterization of absorption and emission dichroism through
polarized spectroscopy. The film continuity was also monitored with tapping mode
atomic force microscopy.
(2) To incorporate PEDOT/PSS conductive alignment layers and
monodomain, uniaxially aligned thin films of monodisperse glassy-nematic
oligo(fluorene)s in polarized OLEDs.
(3) To optimize the polarized OLED performance through device design and
experimentation, including the insertion of an electron-transport layer and variation of
the molecular structure and the thickness of the emitting layer.
76
2. EXPERIMENTAL
Materials and Morphology
The chemical structures of monodisperse oligo(fluorene)s and PEDOT/PSS
are depicted in Chart 3.1. Synthesis, purification, and structural identification of the
conjugated oligomers has been reported previously35. Morphology and the nature of
the phase transition were characterized with hot stage polarizing optical microscopy
(DMLM, Leica, FP90 central processor and FP82 hot stage, Mettler Toledo). Thermal
transition temperatures were determined by differential scanning calorimetry (Perkin-
Elmer DSC-7) with a continuous N2 purge at 20 mL/min. Samples were preheated to
360 oC for F(MB)5 and F(MB)10F(EH)2 to avoid thermal decomposition or to 375
oC for the more stable F(Pr)5F(MB)2, followed by quenching at – 200 oC/minute to
– 30 oC before collecting the reported second heating and cooling scans at ± 20
oC/minute.
77
Chart 3.1: Molecular structures of monodisperse glassy-nematic oligo(fluorene)s and
PEDOT/PSS.
F(MB)5
F(Pr)5F(MB)2
F(MB)10F(EH)2
OO
S+S
OOn
SO3-
n
HO3S
m
PEDOT/PSS
Substrate Preparation
Optically flat fused silica substrates (25.4 mm diameter × 3 mm thick,
transparent to 200 nm, Esco Products) were scrubbed with 0.05 micron alumina
micropolish (Gamma Micropolish II, Buehler), rinsed with de-ionized water,
ultrasonicated in a dilute aqueous detergent (Isoclean, Cole Parmer) at 60 oC for one
hour, thoroughly rinsed with NanoPure water, and finally dried overnight in a laminar
flow hood prior to use. An aqueous dispersion of electronic grade PEDOT/PSS
78
(Baytron P VP AI 4083, H. C. Starck) was applied by spin-coating at 2500 RPM
followed by drying at 120 oC for 30 minutes under argon. Uniaxial rubbing (buffing)
of PEDOT/PSS was performed with a home-built buffing machine comprised of a
velvet cloth on a rotating drum. The height of the drum, and hence the buffing
strength, was controlled with a micrometer. Buffing optimization was performed for
PEDOT/PSS as described in Chapter 2 for polyimide alignment layers using a 0.8 wt.
% solution of F(Pr)5F(MB)2. The results of a typical PEDOT/PSS buffing
optimization are depicted in Figure 3.1. Upon optimization, the alignment ability of
PEDOT/PSS is competitive with insulating polyimide based on the
photoluminescence dichroic ratios obtained for F(Pr)5F(MB)2, 15.7 vs. 16.7,
respectively. All substrate preparation was performed in a class 10,000 clean room
with fume hoods rated as class 100 environments.
79
Figure 3.1: Typical results for PEDOT/PSS alignment layer buffing optimization
performed with a 0.8 wt % solution of F(Pr)5F(MB)2. Rem refers to the dichroic
ratio at the first emission maximum (λ = 420 nm). The micrometer reading that
controls the buffing height is indicated, and Parallel and Perpendicular refer to the
orientation of the first polarizer relative to the nematic director.
0.0
1.0
2.0
3.0
400 500 600
ParallelPerpendicular
ParallelPerpendicular
Em
issi
on I
nten
sity
, a.u
.
Wavelength, nm
4.655
4.70
Rem = 7.6
Rem = 15.7
Preparation and Characterization of Thin Films
Thin oligo(fluorene) films were prepared in the clean room by spin-coating
from filtered UV-grade chloroform solutions at a spin rate of 4000 RPM followed by
overnight drying in vacuo. The solution concentration was varied to modify the
oligomer film thickness. Thermal annealing was performed in an Argon-purged
glassware (Blaessig Glass Specialties) at temperatures 5 to 10 degrees above the glass
transition temperature of the oligo(fluorene) for up to 1 hour. Polarizing optical
microscopy revealed that annealed thin films on optimized alignment layers were
monodomain, or defect-free under 500 × magnification.
Absorption and linear dichroism of monodomain films were evaluated using a
UV-Vis-NIR spectrophotometer (Lambda-900, Perkin-Elmer) and linear polarizers
80
(HNP’B, Polaroid). Fluorescence spectra, with and without polarization analysis,
were collected using a spectrofluorimeter (Quanta Master C60SE, Photon Technology
International) with a liquid light guide directing a 370 nm unpolarized excitation
beam onto the samples in a straight-through geometry. Emitted light normal to the
film surface was detected. The linearly polarized photoluminescence (LPPL) was
characterized by controlling the nematic director of the film relative to a pair of linear
polarizers placed in front of the detector. The film was oriented with its nematic
director vertical or horizontal. For each film orientation, the emission spectrum was
collected twice, changing the orientation of the first linear polarizer such that
polarized emission parallel and perpendicular to the nematic director defined by the
buffing direction was detected. The results from the two film orientations were
averaged to minimize error. Experimental error because of polarization differences in
the dual grating monochromator detector was further reduced by inserting a second
polarizer with its transmission axis at 45° between the first polarizer and the detector
for all measurements.
The continuity of annealed films was evaluated with tapping mode atomic
force microscopy (Nanoscope III, Digital Instruments). Annealed oligo(fluorene)
films for AFM were prepared as described above on quarter-wave flat (at 632.8 nm)
fused silica substrates (12.7 mm diameter by 3 mm thick, Almaz Optics) that had
been previously cleaned, spin coated with PEDOT/PSS alignment layers, and
uniaxially rubbed as described in the preceding sections.
Film thickness was determined with transmission mode spectroscopic
ellipsometry (V-VASE, J. A. Woollam Co.) following literature procedures36. This
81
non-destructive optical technique allowed the determination of the real and imaginary
parts of both the ordinary and extraordinary refractive indices along with the film
thickness.
OLED Preparation and Characterization
OLEDs were prepared on 2.5” square, patterned indium tin oxide (ITO)
coated glass substrates (Polytronix, surface resistivity ~ 75 ohms/square) that had
been subjected to a thorough cleaning procedure. The substrates were exposed to
oxygen plasma for 4 minutes prior to use (PlasmaPreen). PEDOT/PSS films were
spin coated, dried, subjected to buffing optimization, and uniaxially rubbed, followed
by spin coating, vacuum drying, and thermal annealing of oligo(fluorene) films as
described above. Prior to device completion, a small region of the PEDOT/PSS
conductive alignment layer and the annealed oligo(fluorene) film were removed with
acetone and a Q-tip to allow the cathodes to make electrical contact with the patterned
ITO pads. Devices were completed by vacuum evaporation in a multiple source
thermal evaporator, with all depositions at a base pressure of 5 × 10-6 Torr or lower.
The device structures employed in this study are depicted in Figure 3.2a-c.
Preliminary devices were prepared using a LiF/Ca/Al cathode described in the
literature37. Alternatively, devices were constructed with an electron-transport layer
(ETL) of tris(8-hydroxyquinoline) aluminum (Alq3) or tris(N-phenylbenzimidazol-2-
yl)benzene (TPBI). The molecular structures of the electron-transport materials are
shown in Figure 3.2d. In devices with TPBI layers, a thin layer of LiF was employed
to enhance electron injection and transport38. Evaporation rates were controlled by a
82
deposition controller and a quartz crystal microbalance (Infineon IC/5). Organic
materials were evaporated at a rate of 4 Å/s, while LiF was deposited at a rate of 1
Å/s. Mg:Ag cathodes were co-evaporated at rates of 10 and 0.5 Å/s, respectively,
while Ca and Al were evaporated sequentially at rates of 10 Å/s. Metals were
deposited through a shadow mask resulting in OLED active areas of 0.1 cm2.
Devices were encapsulated in a dry nitrogen glove box prior to testing under ambient
conditions using a source-measure unit (Keithley 2400) and a spectroradiometer
(PR650, PhotoResearch). The efficiencies of linearly polarized devices were
measured without polarization analysis. Polarized EL spectra were measured with the
spectroradiometer and a pair of linear polarizers following the same procedure
described for polarized PL. Representative raw OLED data with and without
polarization analysis for devices prepared for this Chapter can be found in Appendix
2.
83
Figure 3.2: OLED device configurations (thicknesses in Ångstroms) and molecular
structures of electron transport materials. PEDOT and oligo(fluorene) film
thicknesses were determined by spectroscopic ellipsometry after spin coating and
thermal annealing, while other layers were measured with a quartz crystal
microbalance during vacuum deposition. a) Polarized OLED without an electron-
transport layer; b) OLED with an Alq3 ETL; c) Polarized OLED with a TPBI ETL,
in which the oligo(fluorene) thickness was varied; d) Molecular structures of Alq3
and TPBI.
ITO
PEDOT (400)
Oligo(fluorene) (500)
LiF (5) Ca (100) Al (1200)
ITO
PEDOT (400)
Oligo(fluorene) (700)
Alq3 (300)
Mg:Ag 20:1 (2200)
ITO
PEDOT (400)
Oligo(fluorene)
LiF (5) TPBI (300)
Mg:Ag 20:1 (2200)
N
N
N
N
N
N
N
ON
OAl
NO
Alq3
TPBI
84
3. RESULTS AND DISCUSSION
Material System and Phase Behavior
In the previous Chapter, the thermotropic, optical, and photophysical
properties of a series of monodisperse oligo(fluorene)s were evaluated and related to
the molecular structures35. In particular, it was found that absorption dichroism,
optical birefringence, and emission polarization increased with molecular aspect ratio.
Based on that work, three representative oligo(fluorene)s with different chain lengths
were chosen for incorporation in polarized OLEDs because of their favorable
morphological properties. The molecular structures of F(MB)5, F(Pr)5F(MB)2, and
F(MB)10F(EH)2 are presented in Chart 3.1. Differential scanning calorimetric
thermograms for the three oligomers at heating and cooling rates of ± 20 oC per
minute are shown in Figure 3.3. F(Pr)5F(MB)2 was preheated to 375 oC before
cooling to – 30 oC prior to collecting the reported second heating and cooling scans,
while the other oligo(fluorene)s were heated to only 360 oC to avoid thermal
decomposition. The elevated glass transition temperatures of 93, 149, and 123 oC for
F(MB)5, F(Pr)5F(MB)2, and F(MB)10F(EH)2, respectively, are appropriate for
OLEDs which can approach operating temperatures of 90 oC under continued forward
bias at high current densities39. Furthermore, the heating and cooling scans presented
no evidence of crystallization, ideal behavior for the preparation of aligned films by
thermal annealing above Tg followed by cooling while preserving the macroscopic
order of the self-assembled nematic fluid in the solid state.
85
Figure 3.3: Differential scanning calorimetry thermograms for the three
oligo(fluorene)s used in OLED studies. G Glassy; N, Nematic; I, Isotropic. Clearing
points and mesophase behavior were confirmed by polarizing optical microscopy.
0 100 200 300 400
End
othe
rmic
Temperature, oC
G N I
N I
I
G N I
G
G
G
G
N
NN
F(MB)5
F(Pr)5F(MB)2
F(MB)10F(EH)2
Thin Film Preparation and Optical Characterization
Thin oligo(fluorene) films were prepared by spin coating from UV-grade
chloroform solutions onto fused silica substrates that had been previously coated with
a PEDOT/PSS layer and uniaxially rubbed under optimized conditions (see
Experimental). The solution concentration was varied to control the oligomer film
thickness. After overnight drying in vacuum at room temperature, the films were
annealed at temperatures 5 to 10 oC above Tg for up to 60 minutes, followed by rapid
cooling to room temperature to arrive at monodomain films. Appropriate annealing
times were determined experimentally by ex situ monitoring of the absorption and
emission dichroic ratios. Higher molecular weight oligomers or thicker films tended
to require longer annealing times to achieve the saturation values. Annealed films
were characterized as monodomain by the absence of nematic textures or defects in
86
polarizing optical microscopy under 500 × magnification. The anisotropic absorption
and linearly polarized photoluminescence of ~ 70 nm thick monodomain films of
F(MB)5 and F(MB)10F(EH)2 are presented in Figure 3.4. Absorption and emission
dichroic ratios were defined as Rabs ≡ A||/A⊥ and Rem ≡ I||/I⊥ in which A and I are
absorbance and emission intensity at λmax, respectively, and subscripts || and ⊥
specify polarization directions parallel and perpendicular to the nematic director
defined by the buffing direction. As shown in Figure 3.4, high orientational order
parameters and photoluminescence dichroic ratios were achieved on PEDOT/PSS
alignment layers buffed under optimized conditions. The order parameter Sabs = 0.84
and photoluminescence polarization ratio RPL = 20.3 at a wavelength of 424 nm for
F(MB)10F(EH)2 demonstrates the quality of alignment that can be obtained with
uniaxially rubbed films of conductive PEDOT/PSS. The use of a PEDOT/PSS
alignment layer is expected to improve the performance of polarized OLEDs because
PEDOT/PSS is widely employed as a hole injection layer in efficient OLEDs32.
87
Figure 3.4: UV-Vis linear dichroism and linearly polarized photoluminescence
spectra for monodomain, ~ 70 nm thick films: a) F(MB)5 and b) F(MB)10F(EH)2.
The orientational order parameter Sabs and the emission dichroic ratio at the first PL
maximum Rem are also included. Subscripts || and ⊥ refer to polarization directions
relative to the orientation of the nematic director defined by the buffing direction.
0.0
0.5
1.0
1.5
300 400 500 6000.0
1.0
2.0
Abs
orba
nce
Em
ission Intensity, a.u.
Wavelength, nm
I
IA
A
= 0.84
= 20.5b) Rem
Sabs
⊥⊥
||
|| at λ = 422 nm
0.0
0.5
1.0
1.5
300 400 500 6000.0
1.0
2.0
Abs
orba
nce
Em
ission Intensity, a.u.
Wavelength, nm
I
IA
A
= 0.75
= 10.4a) Rem
Sabs
⊥⊥
||
|| at λ = 415 nm
Assessment of Film Continuity
One anticipated drawback of conjugated oligomers compared to polymers is
the possibility of diminished mechanical properties because of the low molecular
weight and the absence of chain entanglements. In particular, thermal annealing
could result in local dewetting and the formation of small holes with diameters
beyond the detection limit of optical microscopy in oligomeric thin films. This type
of pinhole defect has been shown to result in the formation of electrical shorts40 or
non-emissive dark spots41 in operating OLEDs, severely reducing device efficiency
and lifetime. Thermal annealing of thin films of relatively low molecular weight
glassy-nematic oligo(fluorene)s above Tg, as required for the production of strongly
polarized emission, could therefore be detrimental to the device performance unless
88
the film continuity can be maintained. An assessment of the continuity of thin films
of conjugated oligomers intended for use in polarized OLEDs was performed with
tapping mode atomic force microscopy (AFM). Both pristine and thermally annealed
thin films of F(Pr)5F(MB)2 on buffed PEDOT/PSS were subjected to AFM
measurements. Films were characterized under ambient conditions before and after
thermal annealing at a temperature of 154 oC, 5 oC above Tg of the hepta(fluorene).
At least four 400 μm2 areas were scanned for each film, resulting in the average RMS
roughness reported herein. Representative AFM surface roughness analysis images
for pristine and thermally annealed films are presented in Figure 3.5. Image analysis
demonstrated that the RMS roughness of the films decreased from 1.34 to 0.66 nm
upon annealing followed by quenching to room temperature to preserve alignment.
Shallow features parallel to the buffing direction were present in both pristine and
thermally annealed and quenched films, evidence that the PEDOT/PSS alignment
layer was lightly scratched when uniaxial rubbing was performed. Films that were
annealed under appropriate conditions of temperature and time were considered to be
continuous because of the absence of pinhole defects in any scanned area. The
relevance of the cooling procedure to maintaining a smooth film topology is also
demonstrated in Figure 3.5. In addition to the rapid quenching described above,
F(Pr)5F(MB)2 films annealed on buffed PEDOT/PSS alignment layers were
subjected to cooling rates of 5 oC/minute or 5 oC/hour. Image analysis revealed that a
slower cooling rate led to increased RMS surface roughness of 1.85 nm for films
cooled at 5 oC/min, while cooling at 5 oC/hour resulted in a further increase up to 3.24
nm. This relatively large surface roughness for slow-cooled films probably obscured
89
scratches that may have been present in the buffed PEDOT/PSS alignment layer. To
reduce the potential effects of surface roughness of the emitter layer on OLED
performance, all annealed oligo(fluorene) films used in polarized OLEDs were
quenched to room temperature.
Figure 3.5: Characteristic 400 μm2 tapping mode AFM surface roughness analysis
images for ~ 70 nm thick films of F(Pr)5F(MB)2 on uniaxially rubbed PEDOT/PSS:
a) Pristine film; b) Thermally annealed and quenched; c) Thermally annealed and
cooled at 5 oC/min; d) Thermally annealed and cooled at 5 oC/hr.
a) b)
c) d)
90
Polarized OLED Fabrication and Characterization
For a preliminary evaluation of device performance, OLEDs were prepared
with the device structure shown in Figure 3.2a. Indium tin oxide (ITO) coated glass
substrates were oxygen plasma treated, spin coated with an aqueous dispersion of
electronic grade PEDOT/PSS, and baked at 120 oC under Argon to dry. After
uniaxial rubbing under optimized conditions, a thin film of F(MB)10F(EH)2 was
fabricated by spin coating and thermal annealing as described above. Prior to loading
the devices in the thermal evaporator, a small region of the PEDOT/PSS and
oligo(fluorene) film was removed to allow for electrical contact between the cathodes
and the appropriate patterned ITO pads (see Experimental). The LiF/Ca/Al cathode
was prepared by sequential thermal evaporation through a shadow mask at rates of
1.0, 10.0, and 10.0 Å per second, respectively, resulting in OLED active areas of 0.1
cm2. OLEDs were encapsulated in a nitrogen glove box before testing at room
temperature under ambient conditions. Figure 3.6a shows the current density-
voltage-luminance relationships for the device, which displayed diode behavior and a
low turn-on voltage of less than 4 V. The maximum luminance yield of 0.18 cd/A
was observed at a current density of 20 mA/cm2. The linearly polarized EL spectra at
the same current density are presented in Figure 3.6b, demonstrating exceptionally
high polarization ratios REL ≡ EL||/EL⊥ of 33.8:1 and 50.1:1 at the EL maxima of 424
and 448 nm, respectively. Subscripts || and ⊥ are as defined for photoluminescence.
Up to the maximum current density employed in device characterization (100
mA/cm2), the EL spectra and the polarization ratio were independent of the current
density. Both the high polarization ratio and the low efficiency are believed to be
91
related to the location of the recombination zone, which appears to be adjacent to the
buffed PEDOT/PSS layer. Carrier recombination close to the uniaxially rubbed
PEDOT/PSS interface is expected to result in high polarization ratios because of
strong surface anchoring and low efficiency due to exciton quenching31, 42. The
previous best polarization ratio for any poly(fluorene) device was 25:1, with a
comparable efficiency (at a somewhat greener EL emission color) of 0.25 cd/A31.
Figure 3.6: Polarized OLED properties of a device comprising a ~ 50 nm thick,
monodomain film of F(MB)10F(EH)2 without an electron-transport layer: a) Current
density-voltage-luminance (IV-LV) relationships; b) Linearly polarized EL spectra at
a current density of 20 mA/cm2; subscripts || and ⊥ are as defined in Figure 3.4.
0.0
1.0
2.0
400 500 600
EL In
tens
ity, a
.u.
Wavelength, nm
RELat λ = 424, 448 nm
EL|| = 33.8, 50.1
EL⊥
0
30
60
90
120
0
40
80
120
160
0 5 10 15 20
Current DensityLuminance
Cur
rent
Den
sity
, mA
/cm
2
Lum
inance, cd/m2
Drive Voltage, V
a)
b)
92
Polarized OLEDs with an Electron-Transport Layer
In many OLED materials, the hole mobility is greater than the electron
mobility, resulting in an imbalance in charge carrier densities that reduces device
efficiency. Electron-transport layers (ETLs) are typically used to improve efficiency
by increasing the electron current and by preventing exciton quenching at the
cathode43. However, there are few reports of polarized OLEDs utilizing ETLs44, 45,
despite the generally low efficiencies for polarized devices without an ETL19-31. Tris-
(8-hydroxyquinoline) aluminum (Alq3), the structure of which is shown in Figure
3.2d, is a well-known electron-transport material in OLEDs and was incorporated in
an OLED with the device structure shown in Figure 3.2b. A 70 nm thick,
monodomain film of F(MB)10F(EH)2 was fabricated on buffed PEDOT/PSS as
described previously. A 30 nm film of Alq3 and a cathode comprising Mg:Ag at a
20:1 ratio were vacuum evaporated sequentially. The EL spectrum for the device, as
shown in Figure 3.7, is dominated by yellow-green unpolarized emission
characteristic of Alq3, with almost no evidence of the blue emission from
F(MB)10F(EH)2. Since the HOMO levels of F(MB)10F(EH)2 and Alq3 are
similar46, 47, it is expected that holes were not confined on the oligo(fluorene) side of
the F(MB)10F(EH)2/Alq3 interface. Additionally, a relatively large barrier for
electron injection from Alq3 to oligo(fluorene) further decreases the probability of
exciton formation in F(MB)10F(EH)2. These factors should result in carrier
recombination very close to the Alq3/oligo(fluorene) interface, where exciton
diffusion or Förster energy transfer to Alq3 are likely.
93
Figure 3.7: EL spectrum of an OLED comprising a ~ 70 nm thick, monodomain film
of F(MB)10F(EH)2 and an Alq3 electron-transport layer (30 nm) at a current density
of 20 mA/cm2.
0.0
0.3
0.6
0.9
1.2
400 500 600 700
EL
Inte
nsity
, Nor
mal
ized
Wavelength (nm)
In order to overcome the unpolarized green emission observed for devices
containing Alq3, an ETL with near-UV emission and a large ionization potential was
desired. The emission requirement would ensure that recombination in the vicinity of
the oligo(fluorene)/ETL interface would result in polarized oligo(fluorene) emission,
and a large ionization potential was expected to prevent an appreciable hole current
from crossing this interface. Based on unpolarized OLED results, tris(N-
phenylbenzimidazol-2-yl)benzene (TPBI, molecular structure shown in Figure 3.2d),
appeared promising when combined with a thin layer of LiF to reduce drive
voltages48. The device structure shown in Figure 3.2c was employed to test TPBI as
an ETL in polarized OLEDs. Thin films of F(MB)5, F(Pr)5F(MB)2, and
F(MB)10F(EH)2 were prepared as described above to ascertain the effects of
oligomer structure on polarized OLED performance. Additionally, the film thickness
was varied with F(MB)10F(EH)2 to further elucidate the effects of device structure
94
on polarized OLED properties. A summary of the results at a current density of 20
mA/cm2 for devices containing an electron-transport layer of TPBI is provided in
Table 3.1.
Table 3.1: Device features and polarized OLED performance data at a current density
of 20 mA/cm2 for devices with a 30 nm TPBI electron transport layer.
Device[a]
Annealing time, min.
Thickness,
nm[b]
Sabs
[c]
Voltage, V
Yield, cd/A
REL
[d] λ1 λ2
REL
[e]
integrated
CIE (x,y)
A 60 73 0.82 9.5 1.10 17.3 15.4 14.4 (0.159, 0.079)
B 60 48 0.84 6.4 0.97 13.3 19.4 15.3 (0.156, 0.070)
C 30 35 0.88 6.2 1.07 27.1 31.2 24.6 (0.159, 0.062)
D 30 70 0.80 7.2 0.98 14.7 10.5 11.7 (0.157, 0.085)
E 30 68 0.75 9.7 0.48 11.8 9.5 8.6 (0.159, 0.091)
Dodecamer, F(MB)10F(EH)2
Heptamer, F(Pr)5F(MB)2
Pentamer, F(MB)5
[a] Experimental uncertainties assessed with 5 sets of device A are as follows: film
thickness, ± 5 nm; orientational order parameter, ± 0.01; drive voltage, ± 0.2 Volts;
yield, ± 0.07 cd/A; emission peak dichroic ratios, ± 0.7; integrated dichroic ratio, ±
0.5; CIE coordinates, ± 0.002, respectively. [b] Determined by variable angle
spectroscopic ellipsometry. [c] Orientational order parameter determined from UV-
Vis linear dichroism. [d] Emission peak maxima at λ1 = 424 and λ2 = 448 nm,
dodecamer; λ1 = 420 and λ2 = 448, heptamer; and λ1 = 412 and λ2 = 444 nm,
pentamer. [e] The EL spectra were integrated from 400 to 600 nm.
95
The data indicate that the orientational order parameter Sabs and both the peak
and integrated EL dichroic ratios increase with the molecular aspect ratio for films of
comparable thickness. The lower drive voltage for Device D compared to Device A
or E may be the result of higher charge carrier mobility in F(Pr)5F(MB)2 compared
to the other oligo(fluorene)s. The less bulky aliphatic pendants in F(Pr)5F(MB)2
could be responsible for this effect. The luminance yields for F(MB)10F(EH)2 and
F(Pr)5F(MB)2 were similar, while that for F(MB)5 was approximately a factor of 2
lower. Shorter conjugated oligomers are expected to possess larger HOMO and
smaller LUMO energies relative to longer homologues, giving rise to a larger
bandgap49. The adjustment in energy levels should yield higher barriers for charge
carrier injection relative to other oligo(fluorene)s, contributing to the lower efficiency
for F(MB)5. It is noted that annealed films of all oligo(fluorene)s have similar
photoluminescence quantum yields35; therefore, this factor is not expected to
contribute to the reduced OLED efficiency for F(MB)5. All devices in Table 3.1
showed pure blue emission based on the CIE coordinates, with values close to the
National Television System Committee (NTSC) standard blue CIE coordinates of
(0.14, 0.08). The polarized electroluminescence spectra for Devices A, D, and E
compiled in Figure 3.8 show the strongly polarized blue electroluminescence that can
be achieved with aligned films of monodisperse glassy-nematic oligo(fluorene)s. It is
surmised that the higher efficiency and lower polarization ratio compared to the
polarized OLED without the ETL both result from the relocation of the recombination
zone close to the oligo(fluorene)/TPBI interface.
96
Figure 3.8: Polarized EL spectra for Devices A, D, and E containing emissive layers
of ~ 70 nm thick, monodomain films of F(MB)10F(EH)2, F(Pr)5F(MB)2, and
F(MB)5, respectively, at a current density of 20 mA/cm2. Subscripts || and ⊥ are as
defined in Figure 3.4.
400 500 6000.0
1.0
2.0
EL
Inte
nsity
, a.u
.
Wavelength, nm
ELDevice D||
EL⊥
REL = 14.7, 10.5@ λ = 420, 448 nm
400 500 6000.0
1.0
2.0
EL
Inte
nsity
, a.u
.
Wavelength, nm
ELDevice E
||
EL⊥
REL = 11.8, 9.5@ λ = 412, 444 nm
400 500 6000.0
1.0
2.0
EL
Inte
nsity
, a.u
.
Wavelength, nm
ELDevice A
||
EL⊥
REL = 17.3, 15.4@ λ = 424, 448 nm
97
Effects of Emissive Layer Thickness
Within the thickness series performed for F(MB)10F(EH)2, Table 3.1 shows
that Sabs and REL increase monotonically as the thickness of the emissive layer
decreases, similar to previous results for the absorption and PL dichroic ratios for
poly(fluorene) films on polyimide alignment layers50. The higher degree of order in
thinner films results from the stronger anchoring provided by the buffed PEDOT/PSS
alignment layer rather than the recombination zone effect observed for devices
without an ETL, as inferred from the practically identical luminance yields for
devices A, B, and C. As expected, the turn-on voltage and the drive voltage to reach
a given current density both decrease with the thickness. This is illustrated in the
current density-voltage-luminance (IV-LV) relationships shown for devices A and C
in Figure 3.9. Note that the turn-on voltage remains below 5 V despite the
incorporation of the 30 nm TPBI ETL.
Figure 3.9: Current density-voltage-luminance (IV-LV) relationships for Devices A
and C comprising monodomain F(MB)10F(EH)2 films of different thickness.
0
20
40
60
80
100
0
300
600
900
1200
0 5 10 15 20
Device CDevice A
Cur
rent
Dem
sity
, mA
/cm
2
Luminance, cd/m
2
Applied Voltage, V
98
The thickness dependence of the EL dichroism is depicted in Figure 3.10,
where the polarized EL spectra for devices B and C are presented. The second
emission peak may have a higher polarization ratio for thinner films because of
differences in light out-coupling or optical cavity effects. For Device C, the
polarization ratio, luminance yield, and CIE chromaticity coordinates are all superior
to previously reported bests achieved independently for 3 different poly(fluorene)
devices. This set of performance data for a single device compares favorably with the
separate achievements of a peak polarization ratio of 2531, an integrated polarization
ratio of 2128, and a luminance yield of 0.66 cd/A with CIE coordinates of (0.17,
0.12)30c. It is noted that a polarized poly(fluorene) device with a PEDOT/PSS hole
injecting layer and a TPBI electron transport layer was published shortly after our
initial report44a, 45. This device exhibited a luminance yield close to 1 cd/A without
disclosing the CIE coordinates and showed considerably lower EL polarization ratios
than those described here (peak ratio 8.4, integrated ratio 6.8). The low polarization
ratios could be attributed to the alignment approach employed, which requires
alignment to be transferred to the emissive poly(fluorene) layer through a removable
liquid crystal template45, although the intrinsic difficulties in aligning poly(fluorene)s
on rubbed alignment layers may also play a role51. A high polarization ratio of 33:1
was recently reported for a polarized poly(fluorene) device prepared by the friction
transfer technique52. However, the luminance yield was only 0.3 cd/A despite an EL
spectrum contaminated by long wavelength emission and the incorporation of an
electron-transport layer of bathocuproine (BCP).
99
Figure 3.10: Absolute polarized EL spectra for Devices B and C for monodomain
F(MB)10F(EH)2 films of different thickness at a current density of 20 mA/cm2.
Subscripts || and ⊥ are as defined in Figure 3.4.
400 500 600
Device C
Device C
Device B
Device B
0.0
1.0
2.0
EL In
tens
ity, a
.u.
Wavelength, nm
EL ||
EL⊥
Effects of Thermal Annealing
A further advantage of monodisperse oligo(fluorene)s over polymers is
depicted in Figure 3.11, which compares the EL performance of pristine films of
F(MB)10F(EH)2 and F(MB)5 with equivalent thicknesses to Devices C and E,
respectively. The essentially isotropic devices emit unpolarized light with practically
identical luminance yields, drive voltages, and CIE coordinates as their linearly
polarized counterparts, showing that the alignment of oligo(fluorene) films to
monodomain through thermal annealing resulted in the achievement of high
polarization ratios without adversely affecting other important device parameters.
This is in sharp contrast to many polarized OLEDs based on poly(fluorene)s, in which
thermal annealing gives rise to color instability and reduced efficiency26, 27, 31.
100
Figure 3.11: Comparison of EL spectra and performance data for essentially isotropic
and linearly polarized OLEDs for a) F(MB)5 and b) F(MB)10F(EH)2.
0.0
0.5
1.0
400 500 600
Pristine
Annealed
Nor
mal
ized
EL
Int
ensi
ty
Wavelength, nm
Voltage = 10.6 VYield = 0.53 cd/ACIE x,y = (0.158, 0.066)
Voltage = 9.7VYield = 0.48 cd/ACIE x,y = (0.159, 0.091)
a)
0.0
0.5
1.0
400 500 600
Annealed
Pristine
Nor
mal
ized
EL
Int
ensi
ty
Wavelength, nm
Voltage = 6.7 VYield = 1.04 cd/ACIE x,y = (0.159, 0.047)
Voltage = 6.2 VYield = 1.07 cd/ACIE x,y = (0.159, 0.062)
b)
4. SUMMARY
Monodisperse glassy-nematic oligo(fluorene)s were successfully employed in
strongly polarized, efficient, deep blue OLEDs. Uniaxial molecular alignment was
accomplished by spin coating on a conductive alignment layer of PEDOT/PSS that
had been uniaxially rubbed under optimized conditions to produce the highest
polarization ratios. The properties of aligned thin films were studied by UV-Vis
absorption dichroism, linearly polarized photoluminescence, spectroscopic
ellipsometry, and tapping mode atomic force microscopy. The chemical structure of
the light-emitting oligomer and the device configuration were varied to optimize the
polarized OLED performance. Superior material properties and ease of processing to
monodomain combined with appropriate device design resulted in world-record
linearly polarized blue OLED performance. Linearly polarized OLEDs may provide
benefits to LCDs and OLED displays, as well as stereoscopic imaging systems and
projection displays. Key experimental findings are summarized as follows:
101
(1) Thin films of three oligo(fluorene)s, chosen for their high glass transition
temperatures and morphological stability, were aligned to monodomain on optimized,
uniaxially rubbed conductive alignment layers of PEDOT/PSS. Absorption and
fluorescence dichroism increase with molecular aspect ratio for films of comparable
thickness. Tapping mode atomic force microscopy revealed that the RMS roughness
of thin films decreases upon annealing provided that quenching to below the glass
transition temperature rather than slow cooling is employed, and showed that aligned
thin films did not contain pinhole defects.
(2) Polarized OLEDs comprising ~ 50 nm thick films of F(MB)10F(EH)2
without an electron-transport layer show very high peak polarization ratios up to 50:1,
but relatively low luminance yields of only 0.18 cd/A. Both aspects of the device
performance were attributed to recombination close to the interface between the
rubbed PEDOT/PSS and the oligo(fluorene).
(3) An OLED comprising a ~ 70 nm thick F(MB)10F(EH)2 film and an
electron-transport layer (ETL) of Alq3 exhibited unpolarized, yellow-green emission
from the ETL. The device performance was limited by a combination of poor hole
blocking and a barrier for electron injection at the oligo(fluorene)/Alq3 interface,
resulting in emission almost exclusively from the Alq3 layer.
(4) Polarized OLEDs comprising oligo(fluorene) films with a TPBI electron-
transport layer showed EL spectra identical to the PL of the conjugated oligomers.
The EL polarization ratios for ~ 70 nm films of F(MB)5, F(Pr)5F(MB)2, and
F(MB)10F(EH)2 increased with the molecular aspect ratio. The drive voltage and
device efficiency were related to the molecular structure of the oligo(fluorene).
102
(5) The dependence of polarized OLED properties on emissive layer thickness
was explored with F(MB)10F(EH)2. Thinner films showed higher polarization
ratios due to enhanced surface anchoring by the uniaxially rubbed PEDOT/PSS.
Drive voltage decreased, as expected, while luminance yield was essentially constant
with decreasing film thickness. The luminance yields suggest that higher polarization
ratios for thinner films are not related to the location of the recombination zone.
World-record device performance was achieved with a 35 nm film, which displayed a
turn on voltage < 5 V, a luminance yield of 1.07 cd/A, an emission peak polarization
ratio > 30:1, an integrated polarization ratio > 24, and CIE chromaticity coordinates
of (0.159, 0.062), all superior to independently demonstrated values from
poly(fluorene)-containing polarized OLEDs.
(6) Thermal annealing of oligo(fluorene) films to monodomain did not
increase the drive voltage, cause color instability, or reduce the luminance yield, in
contrast to previously reported results for poly(fluorene)s.
103
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Chapter 4
Organic Field-Effect Transistors Comprising Glassy-Nematic Films
of Conjugated Oligomers and Polymers
1. INTRODUCTION
Organic materials possessing π-conjugation have been extensively explored
for photonic and electronic device applications including field-effect transistors
(FETs)1-3, light-emitting diodes4, 5, photovoltaic cells6, and lasers7. In particular,
because of their structural uniformity, monodisperse π-conjugated oligomers are
ideally suited for elucidating the relationships between molecular structure, thin film
morphology, and optoelectronic properties while holding promise for practical
applications where chemical purity is of utmost concern8-11. Moreover, monodisperse
liquid crystalline conjugated oligomers can be readily processed into macroscopically
ordered glassy films capable of polarized photo- and electroluminescence12, 13, as
described in Chapters 2 and 3. Uniaxial alignment is also an approach to realize
anisotropic charge transport across large areas in solution-processed thin films14.
In single-crystalline15-17 and polycrystalline18-23 organic materials, high carrier
mobility has been achieved primarily in highly ordered π-stacks that exist in the plane
of the film, which is also the transport direction24. Mobilities superior to that of
amorphous silicon have been realized using organic semiconductors. However,
108
single crystals are limited to small areas and are difficult to implement in practical
devices, while polycrystalline films show mobilities that vary with grain size,
crystallographic defects, and orientation and order within π-stacks2, 3, 18, 19.
Furthermore, the mobility of a solution-processed polycrystalline conjugated polymer
film depends on molecular weight, film preparation protocols, and the nature of the
dielectric interface25-27. While solution processing is preferable to vacuum
evaporation from a cost standpoint, solution-processed films generally exhibit a lower
mobility than that of vacuum-evaporated counterparts28. While amorphous materials
could overcome many of the aforementioned uncertainties, FET mobilities remain
three orders of magnitude below those attainable with crystalline materials29, 30.
As a class of ordered organic semiconductors, liquid crystals have also been
attempted to achieve high mobilities31-36. To fabricate practically useful electronic
devices, mesomorphic order must be preserved in the solid state without
crystallization. Photopolymerization has been performed to freeze liquid crystalline
order in the solid state with mobilities generally below 10–2 cm2/V-s37, 38. Field-effect
mobilities on the order of 10–2 cm2/V-s have been reported for glassy-nematic
conjugated oligomers and polymers14, 39. While photopolymerization involves UV-
irradiation and/or an initiator and inhibitor, which are potentially detrimental to
device performance, the preparation of glassy-nematic films takes advantage of the
glass transition on cooling from the mesomorphic melt.
In Chapter 2, the thermotropic, optical, and photophysical properties of a
series of monodisperse glassy-nematic oligo(fluorene)s were evaluated12b, while
Chapter 3 described the electroluminescent behavior of these materials12c. In this
109
Chapter, the charge carrier transport properties of glassy-nematic conjugated
oligomers and polymers were compared in organic field-effect transistors (OFETs).
A field-effect mobility twice as high as that for PFO was reported previously for the
monodisperse dodeca(fluorene) F(MB)10F(EH)214b. Despite the moderately high
mobilities, these materials are not considered ideal for OFETs because of a large
contact resistance for hole injection from gold electrodes arising from the HOMO
energy levels of 5.8 to 5.9 eV14b, 40, 41. In an effort to reduce the contact resistance
while maintaining the desirable properties of oligo(fluorene)s, viz. good solubility,
high glass transition temperatures, broad mesophase temperature ranges, and an
absence of thermally induced crystallization, bithiophene was chosen as a co-
monomer based on the reported HOMO levels of 5.1 to 5.4 eV for crystalline
fluorene-thiophene co-oligomers42. In this Chapter, the thermotropic and optical
properties of the resulting novel glassy-nematic fluorene-bithiophene co-oligomers
were characterized. OFETs were fabricated with glassy-amorphous, polydomain
glassy-nematic, and monodomain glassy-nematic films to measure the field-effect
mobilities. Liquid crystalline polymer analogues PFO and poly[(9,9-
dioctylfluorenyl-2,7-diyl)-co-bithiophene] (F8T2)43, along with previously reported
oligo(fluorene)s and poly(fluorene)14b, were used to probe the roles played by
oligomer chain length, pendant structure, molecular aggregation in the solid state,
orientational order parameter, and contact resistance.
110
2. EXPERIMENTAL
Materials System and Morphological Characterization
The structures and thermotropic properties of the materials employed in this
study are depicted in Chart 4.1. The synthesis, purification, and structural verification
of the monodisperse oligo(fluorene)s F(Pr)5F(MB)2 and F(MB)10F(EH)2 has been
reported previously12b. Conjugated polymers PFO and F8T2 were purchased from
American Dye Source, Inc. and used without further purification. Based on size-
exclusion chromatography (SEC) with universal calibration performed using
polystyrene standards by Dr. Thomas Mourey of the Eastman Kodak Company, PFO
and F8T2 had absolute weight-average molecular weights of 36,400 and 19,900
g/mol with polydispersity factors of 2.1 and 2.3, respectively. The synthesis,
purification, and structural verification for monodisperse fluorene-thiophene co-
oligomers FTO-1, FTO-2, and FTO-3 was performed following previously reported
procedures12d, 42 as described elsewhere44. Morphology and the nature of the phase
transition were characterized with hot stage polarizing optical microscopy (DMLM,
Leica, FP90 central processor and FP82 hot stage, Mettler Toledo). Thermal
transition temperatures were determined by differential scanning calorimetry (Perkin-
Elmer DSC-7) with a continuous N2 purge at 20 mL/min. DSC samples were pre-
heated to above their clearing points and quenched to – 30 oC prior to collecting the
reported second heating and cooling scans at ± 20 oC/min. Second heating and
cooling DSC thermograms for FTO-1 through FTO-3, F8T2, and PFO are included
in Appendix 3.
111
Chart 4.1: Molecular structures of conjugated oligomers and polymers accompanied
by thermal transition temperatures from differential scanning calorimetry (DSC) and
polarizing optical microscopy (POM). G, Glassy; N, Nematic; K, Crystalline; I,
Isotropic.
F(MB)10F(EH)2 [a]
G 123 oC N > 375 oC II > 375 oC N 111 oC G
F(Pr)5F(MB)2G 149 oC N 366 oC II 357 oC N 141 oC G
n
PFO [b]
G 59 oC, 91 oC K 164 oC N 295 oC II 285 oC N 99 oC K
SS
SS
SS
FTO-1G 105 oC N 236 oC II 231 oC N 95 oC G
SS
SS
FTO-2G 99 oC N 163 oC II 158 oC N 90 oC G
SS
SS
FTO-3G 125 oC N 164 oC II 159 oC N 115 oC G
F8T2 [c]
G 97 oC, 161 oC K 254 oC N 303 oC II 300 oC N 85 oC G
SS
n
F(MB)10F(EH)2 [a]
G 123 oC N > 375 oC II > 375 oC N 111 oC G
F(MB)10F(EH)2 [a]
G 123 oC N > 375 oC II > 375 oC N 111 oC G
F(Pr)5F(MB)2G 149 oC N 366 oC II 357 oC N 141 oC G
F(Pr)5F(MB)2G 149 oC N 366 oC II 357 oC N 141 oC G
n
PFO [b]
G 59 oC, 91 oC K 164 oC N 295 oC II 285 oC N 99 oC Kn
PFO [b]
G 59 oC, 91 oC K 164 oC N 295 oC II 285 oC N 99 oC K
SS
SS
SS
FTO-1G 105 oC N 236 oC II 231 oC N 95 o C G
SS
SS
SS
FTO-1G 105 oC N 236 oC II 231 oC N 95 oC G
SS
SS
FTO-2G 99 oC N 163 oC II 158 oC N 90 oC G
SS
SS
FTO-2G 99 oC N 163 oC II 158 oC N 90 oC G
SS
SS
FTO-3G 125 oC N 164 oC II 159 oC N 115 oC G
SS
SS
FTO-3G 125 oC N 164 oC II 159 oC N 115 oC G
F8T2 [c]
G 97 oC, 161 oC K 254 oC N 303 oC II 300 oC N 85 oC G
SS
n
F8T2 [c]
G 97 oC, 161 oC K 254 oC N 303 oC II 300 oC N 85 oC G
SS
n
[a] Nematic mesophase persisted up to 375 oC beyond which thermal decomposition
was observed. [b] Crystallization at 91 oC and melting to nematic liquid crystal at 164 oC; nematic to isotropic and reverse transitions determined by POM. [c]
Crystallization at 161 oC and melting to nematic liquid crystal at 254 oC.
112
UV-Vis Absorption and Photoluminescence Spectroscopy of Dilute Solutions
Dilute solutions of conjugated oligomers and polymers were prepared at a
concentration of 10-5 M of repeat units in UV-grade dichloromethane. Absorption
spectra were collected with a diode array UV-Vis-NIR spectrophotometer (HP
8453E). Fluorescence spectra of selected materials were collected with a
spectrofluorimeter (QuantaMaster C60SE, Photon Technology International) in a 90 o
orientation. The excitation wavelength was fixed at 370 nm.
Substrate Preparation
Optically flat fused silica substrates (25.4 mm diameter × 3 mm thick,
transparent to 200 nm, Esco Products) were scrubbed with 0.05 micron alumina
micropolish (Gamma Micropolish II, Buehler), rinsed with de-ionized water,
ultrasonicated in a dilute aqueous detergent (Isoclean, Cole Parmer) at 60 oC for one
hour, thoroughly rinsed with NanoPure water, and finally dried overnight in a laminar
flow hood prior to use. Clean substrates were coated with a thin film of a commercial
alignment layer (Nissan SUNEVER grade 610 polyimide varnish, Nissan Chemical
Industries, Japan). The concentrated polyimide varnish was diluted 8 times by mass
with a mixed solvent comprised of 20/80 v/v butyl cellosolve/ethyl cellosolve prior to
spin coating at 3500 RPM for 2 minutes. The solution was applied using a filtered
syringe (0.45 μm Teflon filters, Millipore). After spin coating, the substrates were
soft-baked at 80 oC for 10 minutes followed by hard curing at 250 oC for 1 hour. To
facilitate liquid crystal alignment, the substrates were uniaxially rubbed (buffed) with
a home-built buffing machine by passing the coated substrate 6 times beneath a velvet
113
cloth mounted on a drum, which rotated at ~ 1500 rpm. The height of the drum, and
hence the buffing strength, was controlled with a micrometer. Buffing optimization
was performed as described in Chapter 2. Some PI layers remained unbuffed to
determine the effects of annealing without macroscopic uniaxial alignment on
transistor performance.
Film Preparation and Characterization
Thin films were prepared by spin coating at 4000 RPM from 0.5 to 0.8 wt. %
solutions in UV-grade CHCl3 onto cleaned, polyimide coated fused silica substrates
described above, followed by overnight drying in vacuo. The solution concentration
was varied to maintain the film thickness between 50 and 60 nm. Thermal annealing
was performed at temperatures 5 to 10 oC above the glass transition except for PFO
and F8T2, which were annealed at 220 and 280 oC, respectively. All samples were
annealed for 30 minutes and quenched to room temperature to preserve the uniaxial
alignment and avoid crystallization. Absorption and linear dichroism were
characterized with a UV-Vis-NIR spectrophotometer (Perkin-Elmer Lambda 900) and
linear polarizers (HNP’B, Polaroid). Photoluminescence spectra were collected for
certain materials using the spectrofluorimeter with a liquid light guide directing an
unpolarized excitation beam on to the sample at normal incidence. Emitted light
normal to the film surface was collected, and the aligned films were characterized
with the nematic director oriented vertically and horizontally. The results were
averaged to correct for the polarization preference of the detector. Variable angle
spectroscopic ellipsometry (V-VASE, J. A. Woollam Inc.) was performed following
114
literature procedures45 to measure the film thickness, refractive indices, and
anisotropic absorption coefficients.
Fabrication and Characterization of Field-Effect Transistors
Gold source/drain electrodes were evaporated on both rubbed and unrubbed
polyimide layers to prepare the top-gate OFET device structure reported previously14b
and reproduced in Figure 4.1. The 40 nm thick electrodes were deposited through a
shadow mask at a rate of ~ 2 Å/s resulting in transistors with channel lengths that
ranged from 75 to 120 μm. Conjugated oligomer and polymer films were spin
coated, vacuum dried, and annealed as described above. A 500 to 800 nm thick,
poly-chloro-p-xylylene (parylene-C) film was deposited by CVD on top of the
organic semiconductor to serve as the gate insulator46. Devices were completed by
evaporation of a 30 to 40 nm thick gold gate electrode through a shadow mask
resulting in a channel width of 5 mm. The thickness of the evaporated gold layers
was determined with a calibrated quartz crystal microbalance during deposition,
while that for the gate dielectric was measured ex situ by ellipsometry. Transistors
were characterized with an Agilent 4156C precision semiconductor parameter
analyzer and the data were analyzed using the standard MOSFET equations3b. The
highest occupied molecular orbital (HOMO) energy levels were determined by Drs.
Takeshi Yasuda and Katsuhiko Fujita of Kyushu University, Japan with
photoemission spectroscopy (AC-2, Rikenkeiki) of spin-cast films on indium-tin-
oxide (ITO) electrodes.
115
Figure 4.1: Top-gate OFET device structure. Source/drain electrodes are gold, 40 nm
thick. Channel lengths ranged from 75 to 120 μm, while the channel width was 5
mm.
Glass
Rubbed PI (~ 15 nm)
S D Organic Semiconductor
(50 to 60 nm)
Parylene-C Gate Insulator (500 to 800 nm)
Gold Gate (30-40 nm)
3. RESULTS AND DISCUSSION
Thermotropic and Morphological Properties
The thermal transition temperatures shown with the molecular structures in
Chart 4.1 were determined by differential scanning calorimetry (DSC) and polarizing
optical microscopy (POM) at heating and cooling rates of ± 20 oC/min. The influence
of the molecular structure on thermotropic phase behavior of oligo(fluorene)s has
been reported previously12b and was discussed in Chapter 2. All co-oligomers
showed a nematic mesophase between Tg and Tc and an absence of thermally
activated crystallization, desirable properties for the preparation of ordered solid
films. In contrast, the F8T2 and PFO heating scans displayed crystallization and
melting transitions, necessitating heating to above the melting point for the
preparation of glassy-nematic films of both conjugated polymers. Crystallization in
116
co-oligomers was apparently suppressed by the presence of the branched alkyl side
chains at the 9, 9 position of selected fluorene units. The effect of oligomer length on
transition temperatures was noted through comparison of FTO-1 and FTO-2. The
longer FTO-1 possessed a slightly higher glass transition temperature and a
significantly broader mesophase range because of an increase in Tc by more than 70
oC over FTO-2. Furthermore, FTO-2 and FTO-3 demonstrated the effect of the
aliphatic pendant structure on the glass transition temperature. The 2-methylbutyl
pendant on the central fluorene unit of FTO-3 gave rise to a Tg of 125 oC, an
improvement of 26 oC over FTO-2 with an identical conjugated backbone but bulkier
2-ethylhexyl pendants.
UV-Vis Absorption Spectroscopy of Dilute Solutions and Pristine and Annealed
Films
Absorption spectroscopy was performed to evaluate the optical properties of
the conjugated oligomers and polymers. The UV-Vis absorption spectra were
collected in dilute (10-5 M) UV-grade dichloromethane solutions while films were
prepared by spin coating from UV-grade chloroform on both rubbed and unrubbed
polyimide followed by drying in vacuum at room temperature. The films were
characterized as glassy-amorphous based on their appearance in polarizing optical
microscopy. The UV-Vis absorption spectra for dilute solutions and pristine films are
compiled in Figure 4.2. The pristine film spectra for F(MB)10F(EH)2 and
F(Pr)5F(MB)2 are slightly broadened with a minor red shift in comparison to dilute
solutions. The PFO film spectrum showed a small feature near 425 nm, which was
117
attributed to the formation of the β-phase47. A distinct shoulder on the long
wavelength side of the film spectra was observed for FTO-1 through FTO-3 and
F8T2, which is attributable to backbone planarization or increased ground state
aggregation48. Recently, the appearance of structured absorption in films combined
with a red shift relative to featureless dilute solution spectra has been related to π-
stacking in crystalline conjugated oligomers and polymers where X-ray diffraction
can be used to verify the spectroscopic observations49. Any of these effects could be
expected to improve the carrier transport properties3, 25, 50, 51.
Figure 4.2: UV-Vis absorption spectra of dilute solutions (10–5 M repeat units in
CH2Cl2) and 50 to 60 nm thick, glassy-amorphous films of (a) F(MB)10F(EH)2,
F(Pr)5F(MB)2, and PFO, and (b) FTO-1 through FTO-3 and F8T2.
0.0
0.3
0.6
0.9
1.2
300 350 400 450 500 550 600
PFO
PFO
F(MB)10F(EH)2F(Pr)5F(MB)2
F(MB)10F(EH)2F(Pr)5F(MB)2
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Dilute Solution
Glassy-Amorphous
a)
Film
0.0
0.3
0.6
0.9
1.2
300 400 500 600 700
F8T2
F8T2
FTO-2FTO-3
FTO-2FTO-1
FTO-3
FTO-1
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Dilute Solution
Glassy-Amorphous
b)
Film
0.0
0.3
0.6
0.9
1.2
300 350 400 450 500 550 600
PFO
PFO
F(MB)10F(EH)2F(Pr)5F(MB)2
F(MB)10F(EH)2F(Pr)5F(MB)2
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Dilute Solution
Glassy-Amorphous
a)
Film
0.0
0.3
0.6
0.9
1.2
300 400 500 600 700
F8T2
F8T2
FTO-2FTO-3
FTO-2FTO-1
FTO-3
FTO-1
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Dilute Solution
Glassy-Amorphous
b)
Film
Glassy-amorphous films were subsequently thermally annealed in the fluid
state for 30 minutes prior to quenching to room temperature to preserve uniaxial
alignment in the solid state. Oligomer films were annealed at 5 to 10 oC above their
respective glass transition temperatures, while PFO and F8T2 films were annealed at
118
220 and 280 oC, respectively, well above their melting points. Annealing on
unrubbed polyimide resulted in polydomain glassy-nematic films, while films
annealed on rubbed polyimide were characterized as monodomain based on the
absence of nematic textures or defects in polarizing optical microscopy under 500 ×
magnification. To aid in the visualization of film morphology, polarizing optical
micrographs corresponding to amorphous, polydomain, and monodomain glassy-
nematic films are illustrated in Figure 4.3 for F(Pr)5F(MB)2.
Figure 4.3: Polarizing optical micrographs of three glassy films of F(Pr)5F(MB)2. a)
Glassy-amorphous; b) Polydomain glassy-nematic; and c) Monodomain glassy-
nematic.
Table 4.1 shows that higher quality alignment can be achieved for oligomers
than polymers under milder annealing conditions. It is noted that the orientational
order parameters of conjugated oligomers are primarily determined by the molecular
aspect ratios and appear to be relatively insensitive to the details of film
preparation12b, c, e, in contrast to poly(fluorene)s52. The films of FTO-1 through FTO-
3 were somewhat less ordered than oligo(fluorene) films because of their shorter
molecular lengths and relatively bulkier pendants, and hence lower aspect ratios.
119
Table 4.1: Orientational order parameters of monodomain glassy-nematic films
determined by UV-Vis linear dichroism.
Compound Sabs
[a], [b] F(MB)10F(EH)2 0.86 F(Pr)5F(MB)2 0.83 PFO 0.66 FTO-1 0.81 FTO-2 0.78 FTO-3 0.76 F8T2 0.53
[a] Orientational order parameter Sabs = (Rabs – 1)/(Rabs + 2) where Rabs refers to the
ratio A||/A⊥ where A represents the UV-Vis absorbance at λmax and subscripts || and ⊥
refer to the orientation of the nematic director defined by the buffing direction
relative to a linear polarizer. [b] Values accompanied by an experimental uncertainty
of ± 0.02.
As shown in Figure 4.4, thermal annealing of all materials to monodomain
resulted in hypochromism, a manifestation of π-orbital interactions53. Despite the
difference in chain length, both F(MB)10F(EH)2 and F(Pr)5F(MB)2 showed more
pronounced hypochromism than PFO. With the same conjugated backbone in FTO-
2 and FTO-3, the degree of hypochromism can be related to the aliphatic pendant
structure, i.e. the bulkier 2-ethylhexyl pendants in FTO-2 are responsible for the
lower degree of hypochromism than that observed for FTO-3 with shorter 2-
methylbutyl pendants on the central fluorene unit. No correlation between
hypochromism and the orientational order parameter emerged, as the degree of
hypochromism follows the descending order FTO-3 > FTO-2 > FTO-1 > F8T2.
The opposite trend was noted for the relative intensity of the UV-Vis absorption from
the aggregate band. It appears that the preexisting aggregates limited further
120
advances in molecular aggregation during thermal annealing. The extent of
hypochromism apparently reflects the progression in ground-state aggregation from
glassy-amorphous to monodomain glassy-nematic films and should not be taken as an
absolute measure of molecular aggregation. Furthermore, the spectra presented in
Figures 4.2 and 4.4 suggest that the nature of molecular aggregation responsible for
an additional peak at 475 nm in films is distinct from that responsible for
hypochromism. In the case of F(MB)10F(EH)2 and F(Pr)5F(MB)2, thermal
annealing induced molecular aggregation, as evidenced by hypochromism, without
generating a new peak.
Figure 4.4: UV-Vis absorption spectra of 50 to 60 nm thick, glassy-amorphous and
monodomain glassy-nematic films of (a) F(MB)10F(EH)2, F(Pr)5F(MB)2, and
PFO, and (b) FTO-1, FTO-2, FTO-3 and F8T2.
0.0
0.3
0.6
0.9
1.2
300 350 400 450 500 550 600
F(MB)10F(EH)2F(Pr)5F(MB)2
F(MB)10F(EH)2F(Pr)5F(MB)2
PFO
PFO
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Glassy-AmorphousFilm
Monodomain Glassy-Nematic Film
a)
0.0
0.3
0.6
0.9
1.2
1.5
250 300 350 400 450 500 550
FTO-2
FTO-2
FTO-1
FTO-1
FTO-3
FTO-3
F8T2
F8T2
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Glassy-AmorphousFilm
Monodomain Glassy-Nematic Film
b)
0.0
0.3
0.6
0.9
1.2
300 350 400 450 500 550 600
F(MB)10F(EH)2F(Pr)5F(MB)2
F(MB)10F(EH)2F(Pr)5F(MB)2
PFO
PFO
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Glassy-AmorphousFilm
Monodomain Glassy-Nematic Film
a)
0.0
0.3
0.6
0.9
1.2
1.5
250 300 350 400 450 500 550
FTO-2
FTO-2
FTO-1
FTO-1
FTO-3
FTO-3
F8T2
F8T2
Nor
mal
ized
Abs
orba
nce
Wavelength, nm
Glassy-AmorphousFilm
Monodomain Glassy-Nematic Film
b)
121
OFET Fabrication and Characterization
To evaluate the charge carrier transport, top-gate organic field-effect
transistors (OFETs) were prepared in the device structure reported previously14b and
reproduced in Figure 4.1 (see Experimental). One objective in the molecular design
of fluorene-thiophene co-oligomers was to reduce the effects of contact resistance on
the mobility. Contact resistance has been shown to reduce the mobility extracted
from OFET measurements both theoretically and experimentally40, 41. The HOMO
level for F8T2 has been reported as 5.4 to 5.5 eV41b, which is expected to reduce the
effects of contact resistance relative to PFO, F(MB)10F(EH)2, and F(Pr)5F(MB)2
with HOMO levels of 5.8 to 5.9 eV14b. Ultraviolet photoelectron spectroscopy (UPS)
showed that fluorene-thiophene co-oligomers have HOMO levels of 5.5 eV, close to
that of F8T2. The lower contact resistance was manifested in the transistor output
curves compiled in Figure 4.5. The PFO drain current deviates quite significantly
from linearity at low drain voltages, which is characteristic of contact resistance41.
The effect is less pronounced in F8T2 because of the smaller offset between the work
function of gold and the HOMO level of the copolymer. Figure 4.5 also shows the
output curves for FTO-1, demonstrating the generally lower contact resistance for the
materials containing bithiophene units as compared to fluorene compounds14b.
122
Figure 4.5: Output curves for the 50 to 60 nm thick, monodomain glassy-nematic
films of (a) F8T2, (b) FTO-1, and (c) PFO, all with chain alignment parallel to
current flow.
0 −20 −40 −60 −80 −100 Drain Voltage (V)
Dra
in C
urre
nt (1
0−6 A
)
−1.0
0.0
−2.0
−3.0
−4.0
−5.0
−6.0 −100
−80
−40
Dra
in C
urre
nt (1
0−6 A
)
0.0
−0.4
−0.8
−1.2
Drain Voltage (V) 0 −20 −40 −60 −80 −100
−100
−60
Drain Voltage (V) 0 −20 −40 −60 −80 −100
Dra
in C
urre
nt (1
0−6 A
)
0.0
−1.0
−0.5
−1.5 −100
−75
−50
−25
Gate V
oltage (V)
Gate V
oltage (V)
Gate V
oltage (V)
a) F8T2
b) FTO-1
c) PFO
−50−60
−90
−70
−80
−90
−70
123
Table 4.2 summarizes the field-effect mobilities for all p-type OFETs
fabricated for this study, and also includes the data for fluorene compounds that has
been reported previously14b. Mobility was measured for current flow parallel (μ||) and
perpendicular (μ⊥) to the chain alignment for monodomain glassy-nematic films.
Furthermore, pristine glassy-amorphous (μa) and annealed polydomain glassy-
nematic (μ×) films on unrubbed polyimide were incorporated in OFETs to determine
the effects of thermal annealing on mobility in the absence of alignment. Mobilities
were determined from the saturation regime at drain voltages of – 100 V using the
standard MOSFET equations3b.
Table 4.2: Saturation regime field-effect mobilities in cm2/V-s for OFETs comprising
monodomain and polydomain glassy-nematic, and glassy-amorphous films with
on/off ratios ranging from 103 to 104.
Compound μ|| μ× μ⊥ μa FTO-1 1.8 ± 1.3×10−3 1.1 ± 0.5×10−3 3.2 ± 1.8×10−4 5.0 ± 2.3×10−4 FTO-2 5.6 ± 2.1×10−4 2.9 ± 1.5×10−4 8.9 ± 3.4×10−5 2.0 ± 0.6×10−4 FTO-3 4.3 ± 0.5×10−4 3.0 ± 0.3×10−4 1.2 ± 0.2×10−4 2.6 ± 0.3×10−4 F8T2[a] 1.1 ± 0.5×10−2 5.4 ± 0.8×10−3 2.7 ± 1.0×10−3 3.7 ± 0.4×10−3 F(MB)10F(EH)2[b] 1.2 ± 0.2×10−2 5.1 ± 0.3×10−3 1.9 ± 0.1×10−3 1.5 ± 0.3×10−5 F(Pr)5F(MB)2[b] 1.7 ± 0.4×10−3 7.8 ± 1.8×10−4 1.9 ± 0.1×10−4 2.1 ± 1.1×10−5 PFO[b] 5.8 ± 1.8×10−3 2.2 ± 0.7×10−3 5.2 ± 1.0×10−4 3.3 ± 1.0×10−4
[a] Mobility data agree with the literature14a; [b] Data from a previous report14b.
Regardless of the difference in orientational order parameter, all the
monodomain glassy-nematic films contain intermolecular aggregates as evidenced by
UV-Vis absorption spectroscopy as summarized in Figures 4.2 and 4.4. Of all the
124
measured mobilities, μ|| is the highest because current flow parallel to the chain
alignment involves the least rate-limiting interchain hops. In the fluorene-thiophene
series, μ|| follows the descending order F8T2 > FTO-1 > FTO-2 > FTO-3, which
mirrors the trend in the aggregates’ UV-Vis absorbance in monodomain glassy-
nematic films. These results are consistent with prior reports that ground-state
aggregation gives rise to improved charge-carrier mobility in a poly(p-
phenylenevinylene) derivative51. Despite a reduced contact resistance, the μ|| values
of FTO-1 through FTO-3 are an order of magnitude less than that of
F(MB)10F(EH)2. Moreover, the μ|| value of F(Pr)5F(MB)2 is comparable to that of
FTO-1 and 3 to 4 times those of FTO-2 and FTO-3. It appears that the oligomer
chain length is the predominant factor for μ|| over molecular aggregation, orientational
order, and contact resistance. As reported previously14b, the μ|| value of
F(MB)10F(EH)2 is twice that of PFO probably because of chain entanglements,
kinks, and bends along the polymer chain, all acting as charge traps50, or an inferior
orientational order parameter (0.66 vs 0.86) involving long and winding polyfluorene
chains. Of these two parameters, the effect of orientational order can be downplayed
by the fact that F8T2 (with an orientational order parameter of 0.53) showed a μ||
value comparable to that of F(MB)10F(EH)2 (with an orientational order parameter
of 0.86). A mobility value of 2.0 × 10–2 cm2/V-s measured by the time-of-flight
(TOF) technique has been reported for a structurally similar tetrafluorene32. It should
be noted, however, that the hole mobility measured by TOF can be more than an
order of magnitude higher than the OFET value for the same material because of the
presence of charge traps at the semiconductor-dielectric interface in a transistor29.
125
The contact resistance at the semiconductor-gold interface and the lower applied field
encountered in the measurement using an OFET is also expected to contribute to a
lower mobility value than that determined with the TOF technique.
In an attempt to relate hole mobility to molecular structure, the simplifying
assumption that the length scale responsible for charge transport is proportional to the
oligomers’ extended chain length was employed. Using a semi-empirical method as
part of the Gaussian 03 package, the extended chain lengths of monodisperse
oligomers were calculated via free-energy minimization by Dr. Jane Ou of the
University of Rochester. The results are as follows: 4.1 nm for FTO-2 and FTO-3,
5.8 nm for FTO-1, 5.9 nm for F(Pr)5F(MB)2, and 10.1 nm for F(MB)10F(EH)2.
Primarily, μ|| is determined by two factors, the rate (i.e. μ⊥) and the number of
interchain hops across the channel length. This concept is illustrated in Figure 4.6 for
films comprising a short and a long oligomer. Charge injection is considered to be
independent of chain length for materials with comparable HOMO levels. Intrachain
transport is expected to be fast in both cases and to play a minor role in the overall
mobility in the absence of trapping sites along individual oligomer chains. By further
assuming the same interchain distance for all oligomers, the rate of interchain
hopping should be determined by the extent of translational overlap between
neighboring chains25, which is expected to be greater for longer oligomers than the
shorter counterparts. Furthermore, the number of interchain hops across a given
channel length decreases as the oligomer length increases. These factors should
increase both μ⊥ and μ|| for a longer oligomer. This physical picture is supported by
the observations that μ|| and μ⊥ follow the trend in the length scales characteristic of
126
oligomers (i.e. extended chain length) considering the associated experimental errors,
lending support to the simple physical picture.
Figure 4.6: A schematic diagram of charge transport in an oriented film comprising
(a) a short, and (b) a long conjugated oligomer; charge injection, parallel (i.e.
intrachain) and perpendicular (i.e. interchain) transport are denoted 1, 2, and 3,
respectively.
Source
2
13
23
2
ENematic Director
Source
2
13
23
2
ENematic Directora )
b )
Source
2
13
23
2
ENematic Director
Source
2
13
23
2
ENematic Directora )
b )
In the oriented films consisting of PFO and F8T2, chain entanglements, kinks
and bends (i.e. chain defects) are inevitable. As a consequence, the polymer’s
extended chain length (i.e. contour length) is not appropriate for the interpretation of
carrier mobility. Considering the success of oligomer’s extended chain length in
correlating the hole mobility data, polymer analogue’s persistence length seems to be
a proper candidate. Using the intrinsic viscosity-molecular weight relationship, the
127
persistence lengths for PFO and F8T2 were fit to the hydrodynamic wormlike chain
model of Yamakawa, Fujii, and Yoshizaki54 by Dr. Thomas Mourey of the Eastman
Kodak Company following a previously reported procedure55. The resulting values
of the persistence length for PFO in THF at 30 ºC, lp = 9~10 nm agrees with 9.5 nm
reported by Wu et al.56 in THF at 20 ºC and is slightly longer than 8.6 nm reported by
Grell et al.57 in THF at 40 oC. The measured value lp = 4~5 nm for F8T2 is less than
that of PFO with statistical significance. Since a polydisperse sample is
characterized by a single value of persistence length, it is legitimate to compare the
backbone rigidity of PFO to F8T2 despite the difference in their molecular weight
distributions. Therefore, PFO has a more rigid backbone than F8T2. On the basis of
persistence length alone, one would expect the oriented film of PFO to possess a
higher hole mobility than F8T2 by generalizing the hypothesis successfully tested for
monodisperse oligomers within the framework of Figure 6. The hole mobility data
presented in Table 2, however, indicate that the μ|| value in monodomain glassy-
nematic film of F8T2 is twice that of PFO, suggesting that the persistence length
measured in dilute solution is not the primary determining factor for anisotropic hole
mobility in an oriented conjugated polymer film. One plausible cause pertains to the
nature, and hence the adverse effects, of chain defects on hole mobility that may vary
from PFO to F8T2. Surface anchoring of semiflexible polymer backbones in an
oriented polymer film may also render persistence length irrelevant to charge
transport in neat film. On the other hand, hole mobilitiy in oriented films consisting
of relatively short and structurally uniform oligomers was found to correlate with the
extended chain length presumably because of the absence of chain defects.
128
Since the polymer persistence length did not correlate with field-effect
mobility, the underlying cause of the lower field-effect mobility for PFO compared to
F(MB)10F(EH)2 remained unresolved. However, the photoluminescence spectra
shown in Figure 4.7 suggest a possible explanation for the lower field-effect mobility
of the polymer. The long wavelength emission peaking near 525 nm for annealed
films indicated that keto-defects were present along the PFO chains58, in contrast to
F(MB)10F(EH)2. Keto-defects have been shown to diminish the field-effect hole
mobilities of fluorene-containing materials59. The effect of molecular length on
mobility is obvious by comparison of F(MB)10F(EH)2 to F(Pr)5F(MB)2. The
absence of spectroscopic evidence for better molecular packing or enhanced π-
interactions for one of these oligomers over the other supports the argument that
longer oligomers improve mobility by reducing the number of rate-limiting interchain
hops as well as increasing the hopping rate through enhanced translational overlap.
Figure 4.7: Photoluminescence spectra of dilute (10–5 M in repeat units) CH2Cl2
solutions and 50 to 60 nm thick, monodomain glassy-nematic films of (a) PFO and
(b) F(MB)10F(EH)2.
0.0
0.3
0.6
0.9
1.2
400 450 500 550 600
SolutionPristineAnnealed
Nor
mal
ized
Pho
tolu
min
esce
nce
Wavelength, nm
PFOa)
0.0
0.3
0.6
0.9
1.2
400 450 500 550 600
AnnealedPristineSolution
Nor
mal
ized
Pho
tolu
min
esce
nce
Wavelength, nm
F(MB)10F(EH)2b)
0.0
0.3
0.6
0.9
1.2
400 450 500 550 600
SolutionPristineAnnealed
Nor
mal
ized
Pho
tolu
min
esce
nce
Wavelength, nm
PFOa)
0.0
0.3
0.6
0.9
1.2
400 450 500 550 600
AnnealedPristineSolution
Nor
mal
ized
Pho
tolu
min
esce
nce
Wavelength, nm
F(MB)10F(EH)2b)
129
For films of both conjugated polymers and monodisperse conjugated
oligomers on an unrubbed polyimide layer, a comparison of μ× and μa reveals that
without macroscopic alignment, thermal annealing improved mobility through a
promotion of microscopic-scale molecular aggregation, as reported for amorphous
conjugated polymer films30. That μ× follows the trend in μ⊥ suggests that hole
mobility in polydomain glassy-nematic films is determined by the rate of interchain
hopping within each domain while the domain boundaries present the same resistance
to hole transport regardless of the molecular structure. In glassy-amorphous films,
conjugated segments are randomly oriented, and hence the trend in μa does not follow
that in μ⊥. It appears that both molecular length and aggregation play a role in
determining μa.
4. SUMMARY
This study was motivated to identify factors affecting hole mobility in
uniaxially oriented π-conjugated systems. Thermotropic liquid crystalline fluorene-
thiophene oligomers were characterized for hole mobility using field-effect transistors
and were compared to previously reported monodisperse oligo(fluorene)s14b. Glassy-
amorphous, polydomain glassy-nematic, and monodomain glassy-nematic films were
prepared for a systematic study of the effects on hole mobility of chain length,
molecular aggregation, orientational order parameter, and the contact resistance in a
field-effect transistor through comparison with a polymer analogues. Key findings
are summarized as follows:
130
(1). Branched aliphatic pendants on selected fluorene units of co-oligomers
promoted stable glassy-nematic morphology, reducing the annealing temperature
required for uniaxial alignment. A longer co-oligomer shows a broader mesophase
range because of an increased Tc, while shorter aliphatic pendants promote higher Tg.
(2). UV-Vis absorption spectroscopy of dilute solutions and pristine and
annealed films was used to elucidate the optical properties of conjugated oligomers
and polymers. Pristine films of oligofluorenes showed no evidence of molecular
aggregation, but pronounced hypochromism resulted from thermal annealing. In
contrast, there was spectroscopic evidence for molecular aggregation in both pristine
and annealed films of fluorene-thiophene oligomers, in which thermal annealing also
led to hypochromism.
(3). Ultraviolet photoelectron spectroscopy showed the co-oligomers had
HOMO levels of 5.5 eV, very similar to F8T2. This resulted in reduced contact
resistance in OFETs compared to fluorene compounds, as evidenced by the OFET
output curves.
(4). The mobilities in polydomain glassy-nematic films were found to be
consistently higher than those in glassy-amorphous films, suggesting that molecular
aggregation on the microscopic scale is conducive to charge transport. Moreover, the
trend in μ× follows that in μ⊥.
(5). All oligomers considered, the observed anisotropic mobilities, μ|| (1.2 ×
10–2 to 5.6 × 10–4 cm2/V-s) and μ⊥ (1.9 × 10–3 to 8.9 × 10–5 cm2/V-s) with on/off
ratios of 103 to 104, are largely determined by the oligomers’ extended chain length.
An increasing characteristic length contributes to a greater extent of translational
131
overlap between neighboring chains with fewer interchain hops across a given
channel length, thereby increasing both μ|| and μ⊥ values. The anisotropic hole
mobilities in oriented conjugated polymer films do not correlate with their persistence
lengths measured in dilute solution.
132
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Chapter 5
Effect of Hole Mobility through the Emissive Layer on Temporal
Stability of Blue Organic Light-Emitting Diodes (OLEDs)
1. INTRODUCTION
Organic materials possessing π-conjugation have shown tremendous potential
for applications in optics, photonics, and electronics over the past two decades1. In
particular, the discoveries of efficient organic light-emitting diodes (OLEDs) based
on vacuum evaporated small molecules2 and conjugated polymers that can be solution
processed3 generated renewed interest in the field of organic electronic materials and
devices. The possibility of thin, flat, large area displays and solid-state light sources
provided the motivation for extensive research and development in materials and
device science, resulting in OLEDs spanning the visible spectrum with improved
efficiency and device lifetime4. Operational stability, however, is still a major
challenge to the continued development of OLED technologies5. The stability of
OLEDs containing tris-(8-hydroxyquinoline) (Alq3) emitting green light has been
extensively investigated6. However, a consensus has not been reached regarding the
mechanism of Alq3 device degradation, with experimental evidence suggesting that
both the formation of unstable Alq3+ radical cations6 and the accumulation of trapped
138
charges at the hole-transport layer/Alq3 interface7 decrease the device lifetime. In
contrast, the temporal stability of blue OLEDs remains largely unreported despite the
many classes of blue emitters that have been developed8. Several derivatives of
anthracene have been demonstrated in efficient blue OLEDs with promising
stability9. Nevertheless, the reported lifetimes for blue devices are poor compared to
state-of-the-art green and red OLEDs. Full color OLED displays would therefore
show differential aging, resulting in a loss of color fidelity over time. This can be
partially overcome by gradually increasing the current density for the blue pixels in a
full-color display, at the expense of an even higher degradation rate. Identification of
the key factors influencing device stability is crucial to the realization of superior blue
materials and devices capable of exceeding display market stability requirements.
Fluorene-based polymers10 and oligomers11 are an emerging class of blue
light-emitting materials with substantial promise for OLEDs. Several glassy-
amorphous oligo(fluorene)s have shown good charge transport properties and high
performance deep blue to UV OLEDs12, 13. The glassy-nematic oligo(fluorene)s
described in Chapter 211d were demonstrated in blue polarized OLEDs with world-
record device performance in Chapter 314. The device lifetimes for several blue
poly(fluorene)-based OLEDs have been reported without disclosure of the polymer
structure or the device configuration15. To elucidate the effects of chemical structure
on OLED performance, including stability, oligomers are better suited than polymers
because of structural uniformity and improved chemical purity.
In addition to molecular structure and chemical purity of the light-emitting
material, OLED stability can be influenced by the thermal and electrochemical
139
properties of the transport layers16, the location and extent of the recombination
zone5c, 17, and the formation and location of charge traps7, 18. Both molecular
structures and device configuration can influence some of these factors, implying that
a comparison of OLED stability for different light-emitting materials in varying
device structures is questionable in the absence of careful design, synthesis, and
purification of materials and control of device structures.
With an ultimate goal of developing stable blue light-emitting materials and
devices, the work in this Chapter was motivated by the following objectives:
(1) To design and synthesize blue light-emitting conjugated oligomers based
on anthracene, naphthalene, and fluorene for an assessment of structure-property
relationships;
(2) To fabricate OLEDs for an evaluation of device performance with a focus
on the hole transport layer/emissive layer (HTL/EML) interface and the location and
extent of the recombination zone; and
(3) To provide new understanding of the key parameters determining the
stability of blue OLEDs.
2. EXPERIMENTAL
Material Synthesis and Purification
The chemical structures of the materials prepared for this study, referred to as
model compounds, are shown in Chart 5.1. All solvents, chemicals, and reagents for
synthesis were used as received from commercial sources with the exception of THF
and 9-bromoanthracene, which had been distilled over sodium/benzophenone and
140
recrystallized from ethanol, respectively. Scheme 5.1 shows the procedures followed
for model compound synthesis. 2-bromo-9,9-bis(n-propyl)fluorene (intermediate 1)
9,9-bis(n-propyl)-fluoren-2-yl-boronic acid (2), and 2,7-dibromo-9,9-bis(2-
methylbutyl)fluorene (3) were obtained in moderate to high yields following
previously reported procedures11c. With the exception of ANF, model compounds
were prepared using a single Suzuki coupling reaction19 between the dibromide of the
center aromatic unit and intermediates 2 or 2-naphthaleneboronic acid (4), which was
used as received from Aldrich.
Chart 5.1: Molecular structures of model compounds.
ADN ANF ADF
FFF FDN
141
Scheme 5.1: Synthesis procedures for model compounds.
1 2
Br Br B(OH)2
BrBr BrBr
3
Br + B(OH)2Br
5 64
Br Br
i ii
i
Br
Br
iii iv
ADN
ADF
v4
v2
3
FDNiii4
FFFiii2
2
ANF
v
i) Benzyltriethyl ammonium chloride, NaOH (aq.); ii) 1. nBuLi, − 78 oC; 2. tri-
isopropyl borate, − 78 oC to RT; iii) Pd(PPh3)4, Na2CO3 (2 M aq.); iv) NBS, I2; v)
Pd(PPh3)4, K2CO3 (2M aq.).
General Procedure for Suzuki Coupling
Into a mixture of Br-Ar-Br (1.0 equiv.) and Ar’-B(OH)2, (2.2 equiv.), where
Ar = 9,9-bis(2-methylbutyl)fluorene or anthracene groups and Ar’ = 2-naphthalene or
9,9-bis(n-propyl)fluorene units, and Pd(PPh3)4 (5 mol %) in a flask were added THF
or toluene/ethanol and a 2.0 M aqueous solution of K2CO3 or Na2CO3, respectively
(10 equiv.; THF: water 10:4, toluene:water:ethanol 10:6:2). The reaction mixture was
refluxed under Argon for 2 days. Chloroform was added upon cooling the reaction
mixture to room temperature. The organic layer was separated and washed with brine
142
before drying over MgSO4. Upon evaporating off the solvent, the residue was
purified by column chromatography on silica gel, resulting in light yellow or white
powders.
9,10-bis-(2-naphthyl) anthracene (ADN)
The synthesis and purification of ADN has been reported previously20. A
light yellow solid was obtained in a 37 % yield. 1H NMR (400 MHz, CDCl3): δ
(ppm) 7.95-8.13 (m, 8 H), 7.76-7.78 (dd, 4 H), 7.62-7.68 (m, 6 H), 7.28-7.35 (dd, 4
H). MALD/I-TOF-MS (DCTB) m/z ([M]+): 430.0. Anal. Calcd. for C34H22 (430.2):
C, 94.85; H, 5.15. Found: C, 94.73; H, 5.11. Residual halogen levels (ppm): I, ≤ 10;
Br, ≤ 20; Cl, ≤ 50.
9,10-Bis-(9,9-bis(n-propyl)fluoren-2-yl) anthracene (ADF)
Base and solvent used were K2CO3 and THF, respectively. Eluent used for
column chromatography was hexanes:chloroform (19:1). A light yellow solid was
obtained in an 83 % yield. 1H NMR (400 MHz, CDCl3): δ (ppm) 7.95-7.98 (dd, 2 H),
7.82-7.86 (m, 6 H), 7.37-7.52 (m, 14 H), 1.95-2.07 (m, 8 H), 0.71-0.92 (m, 20 H).
MALD/I-TOF-MS (DCTB) m/z ([M]+): 674.4. Anal. Calcd. for C52H50 (674.4): C,
92.53; H, 7.47. Found: C, 92.47; H, 7.56. Residual halogen levels (ppm): I, ≤ 5; Br,
≤ 20; Cl, ≤ 50.
2,7-Bis[9,9-bis(n-propyl)fluoren-2-yl]-9,9-bis(2-methylbutyl)fluorene (FFF)
Base and solvents used were Na2CO3 and toluene/ethanol, respectively.
Eluent used for column chromatography was hexanes:chloroform (9:1). A white
solid was obtained in a 53 % yield. 1H NMR (400 MHz, CDCl3): δ (ppm) 7.65-7.82
(m, 14 H), 7.34-7.42 (m, 6 H), 2.25-2.32 (td, 2 H), 2.00-2.05 (m, 10 H), 0.67-1.05 (m,
143
32 H), 0.38-0.43 (dd, 6 H). MALD/I-TOF-MS (DCTB) m/z ([M]+): 802.5. Anal.
Calcd. for C61H70 (802.5): C, 91.22; H 8.78. Found: C, 90.84; H, 8.69. Residual
halogen levels (ppm): I, ≤ 10; Br, ≤ 20; Cl ≤ 50.
2,7-Bis(2-naphthyl)-9,9-bis(2-methylbutyl)fluorene (FDN)
Base, solvent, and eluent for column chromatography were identical to those
for FFF. A white solid was obtained in a 34 % yield. 1H NMR (400 MHz, CDCl3): δ
(ppm) 8.13 (s, 2 H), 7.75-7.99 (m, 14 H), 7.52-7.56 (m, 4 H), 2.25-2.31 (m, 2 H),
2.00-2.08 (m, 2 H), 0.63-1.03 (m, 12 H), 0.37-0.42 (dd, 6 H). MALD/I-TOF-MS
(DCTB) m/z ([M]+): 558.4. Anal. Calcd. for C43H42 (558.3): C, 92.42; H, 7.58.
Found: C, 92.18; H, 7.42. Residual halogen levels (ppm): I, ≤ 5; Br, ≤ 20; Cl, ≤ 50.
9-(2-naphthyl)-10-(9,9-bis(n-propyl)-fluoren-2-yl) anthracene (ANF)
9-(2-naphthyl) anthracene (5), was prepared by Suzuki coupling of 9-
bromoanthracene (1 equiv.) and 4 (1.2 equiv.), as shown in Scheme 5.1. Base and
solvent were identical to those for FFF. Eluents for column chromatography were
hexanes:chloroform, 9:1 followed by an additional column at 97:3. A white solid was
obtained in a 70 % yield. 1H NMR (400 MHz, CDCl3): δ (ppm) 8.56 (s, 1 H), 7.96-
8.11 (m, 6 H), 7.71 (d, 2 H), 7.47-7.63 (m, 5 H), 7.33-7.37 (m, 2 H). Bromination of
the 10 position with N-bromosuccinimide (NBS) following the literature21 gave 9-(2-
naphthyl)-10-bromoanthracene (6). Column chromatography with 6:1
hexanes:chloroform gave a yellow solid in an 86 % yield. 1H NMR (400 MHz,
CDCl3): δ (ppm) 8.66 (d, 2 H), 8.03-8.09 (m, 2 H), 7.90-7.94 (m, 2 H), 7.54-7.71 (m,
7 H), 7.36-7.40 (m, 2 H). Suzuki coupling with 1.2 equivalents of 2 using K2CO3 and
THF as the base and solvent, respectively, afforded the target compound. Eluents for
144
column chromatography were identical to those for 9-(2-naphthyl) anthracene. A
light yellow solid was obtained in a 72 % yield. 1H NMR (400 MHz, CDCl3): δ
(ppm) 7.95-8.12 (m, 5 H), 7.83-7.86 (m, 3 H), 7.75-7.77 (m, 2 H), 7.62-7.65 (m, 3 H),
7.50-7.52 (m, 2 H), 7.34-7.44 (m, 7 H), 2.01-2.06 (m, 4 H), 0.72-0.91 (m, 10 H).
MALD/I-TOF-MS (DCTB) m/z ([M]+): 552.0. Anal. Calcd. for C43H36 (552.3): C,
93.44; H, 6.56. Found: C, 93.22; H, 6.62. Residual halogen levels (ppm): I, ≤ 2; Br,
≤ 5; Cl, ≤ 10. The low bromine content in the sublimed final product and the positive
identification of only the molecular ion of the target compound in MALD/I-TOF-MS
provided evidence that undesirable bromination at additional anthracene positions
was avoided in the preparation of 6.
Molecular Structure and Chemical Purity Validation
1H NMR spectra were recorded with an Avance 400 spectrometer (400 MHz).
Elemental analysis was carried out by Quantitative Technologies, Inc. Model
compounds were subjected to temperature gradient sublimation24 prior to further
characterization or experimentation. Molecular weights were measured with a
TofSpec2E MALD/I-TOF mass spectrometer (Micromass, Inc., UK) by Dr. Andrew
Hoteling of the Eastman Kodak Company. HPLC (HP ChemStation 1100 Series,
Hypersil BDS-C18 reverse phase column) with acetonitrile:2-propanol (1:1) as the
solvent and the mobile phase provided additional support for the purity of model
compounds. A single peak was observed with UV-Vis absorption detection at 254
and 295 nm for compounds used in further experiments. Neutron activation analysis
to determine residual halogen levels was performed by Dr. Craig Swanson of the
Eastman Kodak Company.
145
Solution Electrochemistry and UV-Vis Solution Spectroscopy
Cyclic voltammetry (CV) was performed in 0.5 mM solutions of substrate
with a cell potentiostat (Bioanalytical Systems Epsilon). All oxidation measurements
were carried out in Argon-purged anhydrous CH2Cl2 containing 0.1 M tetraethyl
ammonium tetrafluoroborate (Et4NBF4) as a supporting electrolyte. Platinum disk,
platinum wire, and Ag/AgCl (sat’d) were used as the working, counter, and reference
electrodes, respectively. The HOMO energy levels were determined through a
correlation between photoemission spectroscopy and electrochemical
measurements23. The calculated HOMO level for ADN agrees well with the reported
value9a, 24, validating the correlation for model compounds. Absorption spectra of
dilute solutions were collected with a diode array spectrophotometer (HP 8453E).
The optical bandgap was estimated based on the absorption onset of 0.01 mM
solutions in UV-grade CH2Cl2.
Thin Film Preparation and Photoluminescence Spectroscopy
A thermal evaporation system was designed and constructed to enable film
preparation for an assessment of photoluminescence. Figure 5.1 shows the
configuration of the evaporator. Thin films were prepared by vacuum evaporation at
a base pressure of 5×10-6 Torr or lower. Film thickness was measured with a
calibrated quartz crystal microbalance during deposition. Fluorescence spectra were
collected with a spectrofluorimeter (Quanta Master C60SE, Photon Technology
International) with a liquid light guide directing an unpolarized excitation beam at
360 nm onto the sample surface at normal incidence. Emission was collected in a
straight-through geometry.
146
Figure 5.1: Vacuum evaporation system schematic.
CF 6” Viewport
Source
ISO 100 Power/ Thermocouple Feedthrough
Shutter
Substrates
CF 2.75” Cold Cathode Gauge
CF 1.33” Backfill Line
CF 1.33” Roughing
gauge
CF 2.75” Quartz Crystal μbalance
CF 6” to ISO 100 Adapter
CF 4.62” Power/ Thermocouple/ Linear Motion Feedthrough
Cryo- Cooled Jacket
Diffusion Pump
Pump Heater
Backing/ Roughing
Valve
Throttle Valve
Backing/ Roughing
Pump
Backing/ Roughing
Gauge
Key (in-chamber connections)
Support Power Line Thermocouple
High Vacuum Isolation
Valve
OLED Fabrication and Characterization
The OLED device structures prepared for this study are shown in Figure 5.2a,
in which CFx denotes a plasma-polymerized fluorocarbon film included to improve
hole injection25. The molecular structures of the transport materials and light-
emitting dopants incorporated in the OLEDs are shown in Figure 5.2b. Patterned
indium tin oxide (ITO) coated glass substrates (Polytronix) that had been thoroughly
147
cleaned, oxygen plasma treated, and coated with a thin layer of CFx were used as the
anode. Devices were prepared in a multiple source thermal evaporation system at a
base pressure of 5×10-6 Torr or lower. The deposition rate, controlled with a
deposition controller (Infineon IC/5) and a calibrated quartz crystal microbalance
(QCM), was 4 Å/s for all organic materials except dopants, which were co-evaporated
at appropriate rates to obtain the desired doping levels. The rubrene (Rub in Figure
5.2b) doping levels in Devices II and III were 10 and 5 %, respectively. The
magnesium:silver cathodes were co-evaporated through a shadow mask that resulted
in OLED pixels with active areas of 0.1 cm2 at rates of 10 and 1 Å/s, respectively.
All devices were encapsulated in a nitrogen glove box prior to characterization with a
source-measure unit (Keithley 2400) and a spectroradiometer (PhotoResearch PR650)
under ambient conditions. Selected raw OLED data for Devices I through III are
compiled in Appendix 4. To determine the OLED stability, defined as t1/2, the time
required for the OLED luminance to reach 50 % of its initial value, devices were
driven at a constant DC current density of 40 mA/cm2 while the emission intensity
was monitored with a silicon photodiode for 300 hours maximum. A linear
extrapolation was employed to estimate the t1/2 for devices that had not reached 50 %
of their initial luminance during this experiment.
148
Figure 5.2: a) OLED device structures with layer thickness specified in Ångstroms
and rubrene (Rub) doping levels of 5 and 10 % in Devices II and III, respectively; b)
Molecular structures of transport materials and emitting dopants; c) Transient OLED
device structure with a C-545T doping level of 2.5 %.
Model + Rub (100)
Model (300)
Alq3 (200)
ITO
CFx (10)
NPB (750)
Model (400)
Alq3 (200)
ITO
CFx (10)
NPB (750)
I II III
Mg:Ag (2200) Mg:Ag (2200)
NPB + Rub (375)
Model (400)
Alq3 (200)
ITO
CFx (10)
NPB (375)
Mg:Ag (2200)
Model + Rub (100)
Model (300)
Alq3 (200)
ITO
CFx (10)
NPB (750)
Model (400)
Alq3 (200)
ITO
CFx (10)
NPB (750)
I II III
Mg:Ag (2200) Mg:Ag (2200)
NPB + Rub (375)
Model (400)
Alq3 (200)
ITO
CFx (10)
NPB (375)
Mg:Ag (2200)
a)
b)
NN
N
ON
OAl
NO
N N
S
N
OO N
NPB Alq3 Rub BPhen C-545T
c)
Alq3 + C-545T (50 A)
ITO
Platinum (150 A)
Model (3000 A)
BPhen (200 A)LiF (5 A)
Aluminum (1000 A)
CFx (10 A)
Transient OLED Fabrication and Characterization
The transient OLED device structure employed for this study is shown in
Figure 5.2c. ITO coated glass substrates without CFx treatment were sputter-coated
with 1000 Å of chromium to lower the resistivity. Semi-transparent platinum (150 Å)
149
was sputter coated in the anode region as a low resistance hole-injecting contact. The
metallized substrate was treated with CFx prior to the deposition of organic layers as
described above. The concentration of C-545T in the thin Alq3 layer was 2.5 %.
The 5 Å LiF layer was deposited at a rate of 0.1 A/s to improve electron injection and
transport26. The aluminum cathode was deposited at a rate of 15 Å/s through a
shadow mask resulting in a device area of ~ 0.005 cm2 to improve the RC response.
Devices were encapsulated and their steady-state EL spectra were collected as
described above. In contrast to traditional transient OLED measurements27, the
frequency domain response was used to quantify the hole mobility28. A function
generator (HP 8116A) was used to apply 8.0 V DC square wave voltage pulses with a
50 % duty cycle at frequencies ranging from 1 kHz to 25 MHz. Emitted light was
collected with a silicon photodiode (UDT). The voltage pulses and the photodiode
response were recorded with a digitizing oscilloscope (Tektronix TDS 460A). The
EL intensity apparently decreased as the frequency of the voltage pulses increased as
a result of fewer holes reaching the recombination zone. The intercept of two
tangents to the normalized light intensity vs. frequency curves at the maximum light
level and 50 % of this value defined the frequency f such that the hole transit time τtr
= 1/2f. A similar approach is commonly used to extract transit times from time-of-
flight photocurrent transients29 The mobility was then calculated using the film
thickness d, the transit time, and the applied field E according to equation 5-1.
Ed
trτμ = (5-1)
In analyzing the frequency-dependent EL quenching, the electric field was assumed
to drop across the model compound layer and the built-in potential was neglected.
150
3. RESULTS AND DISCUSSION
Material Design, Synthesis, Purification, and Structural Verification
Chart 5.1 depicts the molecular structures of the model compounds used in
this study. Among these compounds, ADN is a well-known material commonly used
as a host in doped blue-emitting OLED devices with promising stability9a. As shown
in Scheme 5.1, all materials were prepared using the Suzuki coupling reaction19. The
model compounds ADN, ANF, ADF, and FFF can be considered as a set with an
increasing content of alkyl-substituted fluorene units, from zero in ADN to three in
FFF. FDN is structurally similar to ADN, with fluorene replacing anthracene as the
central aromatic unit and naphthalene moieties at the 2, 7-positions. Purification by
column chromatography gave white or yellow powders in yields of 34 to 83 %. After
further purification by temperature gradient vacuum sublimation24, model compounds
were characterized with 1H NMR spectroscopy, elemental analysis, HPLC, and
MALD/I-TOF-MS techniques to verify the chemical structure and evaluate the purity.
The MALD/I-TOF-MS spectra for the model compounds, along with the 1H-NMR
spectra of the final products and intermediates 5 and 6, are compiled in Appendix 5.
Solution Electrochemistry and UV-Vis Spectroscopy
Electrochemical properties of the model compounds in solution were
evaluated by cyclic voltammetry (CV). The half-wave oxidation potentials E1/2(ox)
were measured for 0.5 mM solutions in anhydrous methylene chloride with 0.1 M
tetraethyl ammonium tetrafluoroborate as a supporting electrolyte. The HOMO
151
levels were estimated using the E1/2(ox) values and a previously reported linear
correlation between the half-wave oxidation potential and the ionization potential
determined by photoemission spectroscopy23. The optical bandgaps were determined
from the UV-Vis absorption spectra of dilute (10-5 M) solutions in UV-grade
dichloromethane shown in Figure 5.3.
Figure 5.3: UV-Vis absorption spectra of dilute solutions (10-5 M) of model
compounds in UV-grade CH2Cl2.
0.0
0.2
0.4
0.6
0.8
300 400 500 600
FFFADNANFADFFDN
Abs
orba
nce
Wavelength (nm)
Table 5.1 shows that the model compounds with an anthracene core, ADN,
ANF, and ADF, have similar HOMO levels (5.77 ± 0.03 eV) and bandgap energies
(2.97 ± 0.02 eV). The identical HOMO levels within experimental uncertainty
suggest that the energy barrier for hole injection from NPB will be comparable for all
model compounds. Without anthracence, FFF shows a slightly higher bandgap
energy, while FDN did not show reversible oxidation behavior in the range of the
measurement, implying that the HOMO energy is larger than 6 eV.
152
Table 5.1: Measured oxidation potentials, calculated HOMO energy levels, and
optical bandgaps for model compounds and hole transport material NPB.
Compound E1/2ox
mV[a] HOMO[b]
eV Bandgap[c]
eV NPB 770 5.30 3.10 ADN 1236 5.80 2.99 ANF 1214 5.77 2.98 ADF 1190 5.73 2.96 FFF 1278 5.86 3.20 FDN[d] > 1400 > 6 3.31
[a] Relative to Ag/AgCl, determined using 0.5 mM solution in anhydrous CH2Cl2
with 0.1 M Et4NBF4 as a supporting electrolyte; [b] HOMO levels from the literature
for ADN and NPB from photoemission spectroscopy [9a, 23]; other compounds
HOMO levels determined using a linear correlation with measured E1/2(ox) 23; [c]
Estimated from the absorption onset of 10-5 M solution in UV-grade CH2Cl2 (Figure
4.4); [d] Solvent/electrolyte system oxidizes before model compound.
Film Preparation and Photoluminescence Spectroscopy
Thin films were prepared by vacuum evaporation for an assessment of the
photoluminescence. Film thickness ranged from 30 to 50 nm and was measured
during deposition using a calibrated quartz crystal microbalance. The PL spectra
shown in Figure 5.4 were collected in a straight-through geometry with an
unpolarized excitation beam at 360 nm. Identical PL spectra with emission
characteristic of anthracene30 were observed for ADN, ANF, and ADF. The emission
from FFF and FDN displayed a blue shift relative to the anthracene-containing
compounds and vibronic progression was observed, similar to other fluorene-based
light emitters10-14. The larger excited state energy manifested by shorter peak
153
fluorescence wavelength is consistent with the larger optical bandgaps for FFF and
FDN.
Figure 5.4: Normalized photoluminescence spectra of vacuum evaporated thin films
(30 or 50 nm thick) of model compounds with unpolarized excitation at 360 nm.
0.0
0.3
0.6
0.9
1.2
400 500 600
ADFANFADN
FFFFDN
Nor
mal
ized
Em
issi
on
Wavelength (nm)
OLED Fabrication and Characterization
To assess the electroluminescence properties of these model compounds, in
particular their relative stability with respect to OLED operation, three device
structures (I, II, III) were used as depicted in Figure 5.2. Device I has a three-layer
structure where the model compound forms the emissive layer sandwiched between
the NPB hole-transport layer and the Alq3 electron-transport layer. In Device II, the
NPB hole-transport layer is doped with rubrene (Rub in Figure 5.2b) at a
concentration of 10 %. This device structure is used to probe the location of the
carrier recombination zone and the range of electron and/or hole transport in the
model compound. In Device III, a portion of the emissive layer adjacent to the NPB
hole-transport layer is doped with 5 % rubrene. This device structure with a partially
154
doped emissive layer has been commonly employed to improve the efficiency and
stability of OLED devices and is used here to assess structure-property relationships
for the model compounds as a host material. For all steady-state devices used in this
Chapter, ITO modified with CFx25 is the anode and the cathode is Mg:Ag.
Table 5.2 summarizes the EL characteristics and Figure 5.5 shows the
corresponding EL spectra for Devices I through III. For comparison, these data are
grouped according to device type. Several trends correlating with the molecular
structure of the model compounds are noted: (1) the drive voltage increases in the
order of ADN < ANF < ADF; (2) the current efficiency increases in the same order;
and (3) the half-life, in contrast, decreases in the order of ADN > ANF > ADF. In
addition, distinct features are also observed in the EL spectra that can be correlated
with the molecular structures of the model compounds. For Device I comprising
ADN, ANF, and ADF, the EL emission spectra are broad and peak around 460 nm,
similar to the PL spectra for these materials shown in Figure 5.4. However, the
spectrum of the ADN device is somewhat broader than those of ANF and ADF
devices on the long-wavelength side (> 500 nm), suggesting that there is an additional
contribution to the EL emission from the green-emitting Alq3 electron-transport
layer. Device I containing either FFF or FDN exhibited featureless blue EL emission
as shown in Figure 5.6, in contrast to the structured PL shown in Figure 5.4. This
emission was found to arise primarily from NPB by comparison to the PL spectrum
of a NPB thin film, which is also shown in Figure 5.6. This behavior suggests severe
hole accumulation at the NPB/FFF or FDN interface resulting in NPB EL because of
its smaller bandgap compared to FFF or FDN. Because the emission did not occur
155
from the model compounds, the device lifetimes were omitted from Table 5.2 and
FFF and FDN were excluded from further experiments. For Devices II and III, the
EL spectra are dominated by the orange emission (> 550nm) from rubrene while blue
components attributable to the model compounds are present at much lower
intensities. The ADN devices appear different in that the relative EL emission is
higher in the green region (~ 500 nm) compared to the ANF and ADF devices,
consistent with Device I. These observations can be understood if decreasing hole
mobility of model compounds with increasing fluorene content reduces the width of
the recombination zone. The mobility decrease was confirmed through the
fabrication and characterization of transient OLEDs as described in the following.
Table 5.2: Summary of drive voltages, EL efficiencies, and CIE coordinates at a
current density of 20 mA/cm2. The device lifetime at a current density of 40 mA/cm2
is also included.
Drive Voltage
Volts Efficiency
cd/A CIE x, y Device Half-life
t1/2, hours[a] Device I ADN 8.8 0.69 (0.163, 0.137) 300 ANF 9.2 0.70 (0.151, 0.098) 135 ADF 9.5 0.85 (0.149, 0.084) 100 FFF[b] 13.4 0.37 (0.161, 0.126) N/A FDN[b] 13.5 0.61 (0.160, 0.108) N/A Device II ADN 8.2 2.98 (0.470, 0.474) 1000 ANF 8.6 3.16 (0.469, 0.461) 750 ADF 9.5 3.42 (0.473, 0.453) 280 Device III ADN 7.1 2.24 (0.437, 0.487) 1000 ANF 7.8 3.14 (0.468, 0.502) 475 ADF 9.1 3.24 (0.455, 0.494) 225
156
[a] Linear extrapolation employed for devices that did not reach their t1/2 during the
stability test conducted for 300 hours maximum. [b] Emission from NPB rather than
model compound was observed.
Figure 5.5: Normalized EL spectra at j = 20 mA/cm2 for a) Device I; b) Device II; and
c) Device III comprising ADN, ANF, and ADF, all providing evidence in support of
a varying recombination zone width.
0.0
0.3
0.6
0.9
1.2
400 500 600 700
ADNANFADF
Nor
mal
ized
Lum
inan
ce
Wavelength (nm)
a)
0.0
0.3
0.6
0.9
1.2
400 500 600 700
ADFANFADN
Nor
mal
ized
Lum
inan
ce
Wavelength (nm)
b)
0.0
0.3
0.6
0.9
1.2
400 500 600 700
ADNANFADF
Nor
mal
ized
Lum
inan
ce
Wavelength (nm)
c)
157
Figure 5.6: a) Normalized Device I EL spectra comprising FFF and FDN layers at a
current density of 20 mA/cm2. The photoluminescence spectrum of a thin film of
NPB is shown for comparison.
0.0
0.3
0.6
0.9
1.2
400 450 500 550 600
NPB
FDNFFF
Nor
mal
ized
Lig
ht In
tens
ity
Wavelength, nm
Photoluminescence
Electroluminescence
Transient OLED Fabrication and Characterization
Transient OLEDs were fabricated to measure the hole mobility. Semi-
transparent platinum anodes were sputter coated and treated with CFx to reduce the
device resistance and improve hole injection25. A high electron-mobility Li-doped
electron transport layer26 was used to ensure that the transient response be determined
by hole transport in the model compound. The transient OLED device structure is
shown in Figure 5.2c.
Contrary to typical transient OLEDs27, the frequency-domain response was
used to measure the hole mobility28. The decrease in EL intensity at increasing
frequencies of DC voltage pulses was apparently the result of fewer holes reaching
the Alq3 + C-545T layer. The hole transit time was estimated by the intercept of the
tangents to the light output vs. frequency curve. The mobility was then calculated
based on the film thickness and the applied field. To validate this approach, a 3500 Å
158
film of NPB was used as the model layer in a transient OLED. Based on the
frequency-dependent EL quenching response shown in Figure 5.7, the mobility was
9.6 × 10–4 cm2/V-s at an applied field of 2.3 × 105 V/cm, a value consistent with the
time-of-flight transient photocurrent measurement, (8.8 ± 2.0) × 10–4 cm2/V-s31. With
confidence based on the results for NPB, 3000 Å films of model compounds were
tested in transient OLEDs.
Figure 5.7: Frequency-Dependent EL Quenching for a Transient OLED Comprising a
3500 Å film of NPB upon Application of 8 V Square Wave DC Voltage Pulses at an
Increasing Frequency, 50 % Duty Cycle. The solid line represents a best fit to a
polynomial.
0.3
0.6
0.9
1.2
102 103 104 105 106 107
NPB
Nor
mal
ized
EL
Inte
nsity
Frequency (Hz)
The typical steady-state EL spectrum shown in Figure 5.8a for a device
comprising an ADN layer exhibited emission from C-545T only. The frequency-
dependent EL quenching observed for the transient OLEDs is shown in Figure 5.8b,
from which the hole mobility (in cm2/V-s) for ADN, ANF, and ADF was estimated at
3.1 × 10-4, 8.9 × 10-5, and 3.6 × 10-5, respectively. Repeated experiments with ADN
159
yielded an experimental uncertainty of ± 7 %. The observed effect of the fluorene
content in the model compounds on hole mobility is limited to the three model
compounds in amorphous films, as hole mobility is known to depend on film
morphology and molecular orientation.
Figure 5.8: a) Normalized steady-state EL spectrum of a transient OLED containing
3000 Å of ADN as the model layer; b) Frequency-dependent EL quenching of the
transient OLEDs comprising model compounds, 8 V DC square wave voltage pulses,
50 % duty cycle. The solid lines represent best fits to polynomials.
0.0
0.3
0.6
0.9
1.2
500 600 700
Nor
mal
ized
EL
Inte
nsity
Wavelength (nm)
a)
0.0
0.3
0.6
0.9
1.2
104 105 106 107
ADFANFADN
Nor
mal
ized
Inte
nsity
Frequency (Hz)
b)
0.0
0.3
0.6
0.9
1.2
500 600 700
Nor
mal
ized
EL
Inte
nsity
Wavelength (nm)
a)
0.0
0.3
0.6
0.9
1.2
104 105 106 107
ADFANFADN
Nor
mal
ized
Inte
nsity
Frequency (Hz)
b)
Additional evidence in support of the hole-transport limitation is provided in
Figure 5.9, in which the current density dependence of Device III with ADN, ANF,
and ADF is presented. For the ADN device, the spectrum is dominated by Alq3
emission at the lowest current density of 0.5 mA/cm2, suggesting that hole transport
across the ADN layer is facile even at low voltages. As the current density increases,
the primary emission originates from rubrene with an increasing contribution from
160
ADN emission at 448 nm. In contrast, the EL spectra for Device III with ANF and
ADF are essentially independent of current density from 0.5 to 100 mA/cm2.
Figure 5.9: Current density dependence of Device III EL spectrum for a) ADN; b)
ANF; c) ADF.
0.0
0.3
0.6
0.9
1.2
400 500 600 700
0.52.06.02040100
Nor
mal
ized
Lum
inan
ce
Wavelength (nm)
j (mA/cm2) =a)
0.0
0.3
0.6
0.9
1.2
400 500 600 700
0.52.06.02040100
Nor
mal
ized
Lum
inan
ce
Wavelength (nm)
j (mA/cm2) =
b)
0.0
0.3
0.6
0.9
1.2
400 500 600 700
0.52.06.02040100
Nor
mal
ized
Lum
inan
ce
Wavelength (nm)
j (mA/cm2) =
c)
161
Effects of Hole Mobility on OLED Performance: Discussion
The increase in the drive voltage from ADN to ADF devices can be attributed
to the hole transport limitation in the model compound with increasing fluorene
content. The fact that the same trend is observed for the three types of devices
indicates that the voltage across the model compound layer is significant and the
effect of the molecular structure on the electrical impedance of this layer is reflected
in the overall drive voltage on the OLED device. Based solely on the drive voltages,
it is not possible to distinguish if different electron mobilities affect the device
performance. However, if different electron mobilities were contributing to the
recombination zone width, ADF, which exhibits the least Alq3 emission in Devices I
through III, would be expected to show better electron transport than ADN or ANF.
This is inconsistent with the higher drive voltages for observed for Devices I through
III comprising ADF than for the equivalent ADN and ANF devices.
The increase in efficiency from ADN to ADF devices can also be understood
in light of the transport limitation. Recombination at the hole-transport
layer/emissive layer (HTL/EML) interface is expected to be more efficient in EL
emission than that at the emissive layer/electron-transport layer (EML/ETL) interface.
Thus, a hole-transport limitation would result in a higher EL efficiency because of a
greater extent of recombination at the HTL/EML interface. Conversely, an electron-
transport limitation would result in a lower EL efficiency. Therefore, the observed
efficiency trend agrees with ADN supporting hole transport better than ANF and
ADF.
162
The EL spectra shown in Figure 5.5 are also a result of the hole mobility
descending in the order ADN > ANF > ADF. These spectra show different degrees
of contribution to the overall EL emission from the Alq3 electron-transport layer as a
result of higher hole mobility for ADN than for ANF or ADF. For Device I
comprising undoped model compounds, the EL spectrum for ADN is distinctly
broader compared to that of ANF and ADF. This broadening in the long-wavelength
edge can be attributed to more holes approaching the Alq3 interface in the ADN
device, resulting in higher emission intensity from the Alq3 layer in the ADN device
than in the ANF and ADF devices. For Devices II and III, the EL spectra are
dominated by orange emission from rubrene with a reduced blue emission
contribution from the model compounds. However, as in Device I, emission from the
Alq3 electron-transport layer is appreciable for the ADN devices.
The differences in hole mobility is also expected to play a role in the relative
stability of the devices. The lifetimes for Devices I, II, and III, defined as t1/2, the
time required for the OLED luminance to reach 50 % of its initial value, are listed in
Table 5.2. Since ADN can support hole transport to a larger degree resulting in a
more extended recombination zone, more ADN molecules participate in the
recombination processes. Therefore, even if the degradation rate per molecule were
the same regardless of the molecular structure, the operational lifetime for the ADN
device would be longer compared to the ANF and ADF devices where the
recombination zone is more confined towards the NPB hole-transport layer due to
reduced hole mobility in ANF and ADF compared to ADN.
163
Devices II and III were designed to study the dependence of the EL
characteristics on the charge transport and recombination in the model compounds
using a doped device structure. Rubrene was chosen as the dopant because its orange
emission is easily differentiated from the blue emission of the model compound, and
thus can be used to locate the recombination zone. In Device II, rubrene was doped
into the NPB hole-transport layer whereas in Device III, it was doped into the
emissive model compound layer. For Device II, rubrene is not expected to
significantly affect the transport of holes in NPB because the HOMO level of rubrene
at 5.36 eV32 is almost iso-energetic with that of NPB, which would make it
ineffective as a hole trap. This, in fact, is the case, as inferred from the similar drive
voltage trend and magnitude for Devices I and II as shown in Table 5.2. Doping the
model compound with rubrene in Device III reduces the operating voltage compared
to Device I by providing an energetically favorable site for holes on the EML side, as
the HOMO of rubrene is substantially lower than those of the model compounds
(Table 5.1). The efficiency of Device III containing ADN is particularly low. The
relatively strong Alq3 emission implies that a significant number of excitons are
present within 20 nm of the metal cathode, which may have decreased the efficiency
because of cathode quenching33.
Table 5.2 shows consistent trends in drive voltage, efficiency, and device
lifetime across the three OLED configurations, strong evidence that the results are not
related to device structure. The EL emission near 500 nm in Figure 3 appears to arise
from Alq3, suggesting the recombination zone width follows ADN > ANF > ADF in
all the OLED devices. Carrier accumulation at the HTL/EML interface because of
164
large energy barriers or drastic changes in mobility would be expected to confine
recombination close to this interface. The HOMO levels of the model compounds,
however, suggest that the energy barriers for hole injection are nearly identical,
further reinforcing the notion that the model compounds’ hole mobility has the most
substantial impact on the recombination zone width and hence the device efficiency
and lifetime. The effects of chemical structure on transport properties and OLED
efficiency have been demonstrated without disclosing the stability34. However, this
work directly correlates the molecular structure, carrier transport properties, and
OLED lifetime for blue light-emitting materials in neat and doped OLEDs35.
4. SUMMARY
Blue light-emitting anthracene, naphthalene, and fluorene-containing model
compounds were designed and synthesized in an effort to gain new insight into the
origins of instability of blue OLEDs. Photo-physical and electrochemical properties
were characterized as parameters for investigation. Three OLED configurations were
employed to assess the EL performance and device lifetime. Key findings are
summarized as follows:
(1) Three model compounds with an anthracene core had the same HOMO
energy level within experimental error, and hence there is no variation in hole
injection barrier from NPB to the emissive layer comprising the model compounds.
Furthermore, ADN, ANF, and ADF model compounds showed the same PL spectra
characteristic of anthracene as the central unit.
165
(2) OLEDs comprising FFF and FDN showed EL emission from NPB instead
of the model compounds due to strongly localized recombination at the NPB/model
compound interface in these cases, resulting in EL from NPB because of its smaller
bandgap energy compared to FFF and FDN.
(3) The varying contribution of green emission from Alq3 as the electron-
transport layer to the EL spectra in ADN, ANF, and ADF-containing OLEDs was
attributable to the difference in hole mobility through the emissive layer, which
affected the width of the recombination zone. The hole mobilities were measured
with the transient OLED technique, yielding values of 3.1 × 10–4, 8.9 × 10–5, and 3.6
× 10–5 cm2/V-s for ADN, ANF, and ADF, respectively.
(4) The higher the hole mobility through the emissive layer, the wider the
recombination zone. A wider recombination zone was found to be responsible for a
longer device lifetime and a lower drive voltage at the expense of a lower luminance
yield, demonstrating the direct relationship between hole mobility and OLED
temporal stability for the blue light emitters studied herein.
166
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Chapter 6
Summary, Conclusions, and Potential for Future Work
1. SUMMARY AND CONCLUSIONS
Organic electronic materials and devices have been among the most actively
pursued research areas in materials science during the past twenty years. This thesis
research focused on organic light-emitting diodes (OLEDs) and field-effect
transistors (OFETs) using monodisperse conjugated oligomers comprised primarily of
fluorene, thiophene, anthracene, and naphthalene units as active materials. The
chemical purity, structural regularity, and uniformity inherent to monodisperse
systems allowed structure-property relationships to be developed and facilitated the
optimization of organic electronic device performance. Understanding how
molecular structure impacts device behavior is crucial to the design of the next
generation of organic semiconductors.
In the first part of my studies, the thermotropic and optical properties of a
novel series of monodisperse glassy-nematic conjugated oligo(fluorene)s with chain
lengths ranging from five to twelve fluorene units substituted with linear or branched
aliphatic pendants at the 9,9-position were characterized. Structural variations
included co-oligomers with different alkyl pendants along the fluorene backbone.
Five fluorene units was the minimum oligomer length for which mesomorphism was
171
observed. Both the oligomer chain length and the pendant structure were found to
affect the thermal transition temperatures and the type of the phase transitions. In
general, aliphatic pendant branching and mixing as in co-oligomers were found to
promote stable glassy-nematic phase behavior compared to linear pendants and
homo-oligomers, respectively. Nematic liquid crystals with glass transition
temperatures up to 149 oC, nematic fluid temperature ranges greater than 200 oC, and
an absence of thermally induced crystallization were realized through the proper
combination of oligomer length and aliphatic pendants. Defect-free, macroscopically
oriented glassy-nematic thin films were prepared on uniaxially rubbed polyimide
alignment layers by spin-coating, thermal annealing at temperatures 5 to 20 oC above
Tg for 15 to 30 minutes, and cooling to ambient while maintaining the characteristic
orientation of the nematic mesophase. The UV-Vis linear dichroism, optical
birefringence, and photoluminescence dichroic ratio were found to increase with the
molecular aspect ratio. The monodomain aligned thin films exhibited linearly
polarized blue photoluminescence with emission dichroic ratios up to 17:1,
competitive with the best previous results for poly(fluorene)s. Annealed thin films
showed quantum yields up to 60 %, suggesting the potential for application in linearly
polarized OLEDs. Furthermore, annealing above the glass transition for 96 hours did
not diminish the spectral quality of the blue photoluminescence through excimer or
keto-defect formation.
Following the above work, representative monodisperse glassy-nematic
conjugated oligomers with five, seven, and twelve fluorene units chosen for their
favorable morphological properties were applied to study linearly polarized blue
172
OLEDs. To maintain a low driving voltage and enhance the power efficiency,
conductive alignment layers of uniaxially rubbed PEDOT/PSS were used to prepare
monodomain thin films of the three monodisperse oligo(fluorene)s in place of the
insulating polyimide films for polarized photoluminescence. A comparison of the
anisotropic absorption and fluorescence properties indicated that PEDOT/PSS aligned
the conjugated oligomers to the same extent as optimized polyimide. The topology
and continuity of annealed thin films intended for polarized OLEDs were evaluated
with tapping mode atomic force microscopy. Films that were quenched to below Tg
to preserve macroscopic alignment exhibited RMS surface roughness below 1 nm and
were continuous based on the absence of pinhole defects. A polarized OLED device
comprising an annealed dodeca(fluorene) film without an electron-transport layer
(ETL) showed a polarization ratio of 50:1, twice that of any previously reported
poly(fluorene) device, with a comparable efficiency of 0.18 cd/A. Recombination
near the interface with rubbed PEDOT/PSS was responsible for both aspects of the
device performance, as strong surface anchoring resulted in highly polarized
luminescence at the expense of device efficiency because of exciton quenching. The
efficiency increased more than five-fold by including a 30 nm thick ETL of TPBI,
which also reduced the polarization ratio to 19:1 at the same dodeca(fluorene) film
thickness as a result of carrier recombination occurring primarily at the
oligo(fluorene)/TPBI interface. By decreasing the thickness of the light-emitting
layer to 35 nm, the polarization ratio was improved to 31:1 while maintaining the
efficiency, suggesting that the increased polarization ratio was related to stronger
surface anchoring in thinner films without affecting the location of the recombination
173
zone. Furthermore, all the linearly polarized OLEDs with an ETL of TPBI exhibited
pure blue electroluminescence, as indicated by the CIE chromaticity coordinates close
to the NTSC standard blue of (0.14, 0.08). Annealing did not result in increased drive
voltages or reduced efficiencies, in contrast to previous results for poly(fluorene)s.
The performance data for a single blue polarized device comprising a 35 nm thick
dodeca(fluorene) emitting layer surpassed previous bests in polarization ratio,
efficiency, and CIE coordinates achieved independently for three different
poly(fluorene) devices.
To evaluate the charge carrier transport properties, the same hepta(fluorene)
and dodeca(fluorene) used for polarized OLEDs had been previously incorporated in
organic field-effect transistors (OFETs) along with PFO, a corresponding
poly(fluorene). To determine the effects of chemical composition, chain length, and
pendant structure on hole mobility, field-effect transistors comprising monodisperse
co-oligomers of fluorene and bithiophene units were prepared and characterized in a
comparative study. Branched pendants in co-oligomers promoted stable glassy-
nematic phase behavior, permitting film processing at lower temperatures compared
to the polymer analogue. Molecular aggregation was detected in pristine glassy-
amorphous thin films of fluorene-bithiophene compounds with UV-Vis absorption
spectroscopy, in contrast to oligo(fluorene)s. The orientational order parameters of
films that had been thermally annealed to monodomain depended primarily on the
molecular aspect ratios. Furthermore, thermal annealing resulted in hypochromism
for all materials, a manifestation of π-orbital interaction in the solid state. Top gate,
p-type OFETs were prepared to measure the charge carrier transport properties, with
174
the output curves indicating that inclusion of bithiophene groups reduced the contact
resistance compared to oligo- and poly(fluorene) devices. Nevertheless, it was
apparent that the extended length played the predominant role in determining the hole
mobility for monodisperse conjugated oligomers. This was interpreted on the basis of
fewer rate limiting interchain hops per unit channel length for longer oligomers as
well as an increased rate of interchain hopping because of improved translational
overlap compared to shorter chains. In contrast, the field-effect mobility of
monodomain-aligned films of polymer analogues did not correlate with the
persistence lengths measured for dilute solutions. Emissive keto-defects in the
poly(fluorene) homopolymer may have contributed to a lower field effect mobility for
PFO compared to the monodisperse dodeca(fluorene).
To enable a study of OLED stability, five monodisperse anthracene,
naphthalene, and fluorene-containing oligomers were designed and synthesized. The
conjugated units in ADN, a well-known material commonly used as a host in stable
blue OLEDs, were replaced with alkyl-substituted fluorene units in a stepwise
approach, giving a series of materials with varying fluorene content. A fluorene-core
equivalent of ADN completed the series, for which the photo-physical and
electrochemical properties were characterized. Three OLED configurations were
prepared to study OLED performance, with the drive voltages, EL efficiencies, and
device lifetimes suggesting a decreasing hole mobility as the fluorene content
increased. Transient OLEDs were fabricated to quantify the hole mobilities, with
results confirming the interpretation of the steady-state OLED performance data. The
hole mobility dictated the OLED performance by adjusting the width of the
175
recombination zone, which was inferred from the contribution to the overall EL
spectra from the Alq3 electron-transport layer. A material with a higher hole
mobility resulted in a broader recombination zone, with more molecules participating
in the recombination process, thereby improving the temporal stability. This
represents a direct demonstration of the relationship between the molecular structures
of blue emitting materials, their charge carrier transport properties, and the stability of
OLEDs in which they function as neat emitting layers or hosts for fluorescent
dopants.
2. POTENTIAL AVENUES FOR FUTURE WORK
The organic electronic devices incorporating monodisperse conjugated
oligomers described in this thesis suggest many possibilities for future research aimed
towards further improving device performance. While the polarized OLEDs
discussed in Chapter 2 represent the state of the art, the device lifetime must be
improved to meet practical display requirements. The instability may be caused by
the emitting material, interfaces within the device, or both. Single-carrier devices
combined with photoluminescence could provide insight into the stability of the
emissive material1. The interface between PEDOT/PSS and oligo(fluorene) is
expected to be problematic based on recent reports, where the insertion of a hole-
transporting interlayer between PEDOT and poly(fluorene) improved device
efficiency and stability2-5. However, an interlayer for polarized OLEDs must not
interfere with monodomain alignment and should be insoluble during subsequent
fabrication steps. Cross-linkable materials represent an attractive approach,
176
especially those in which linearly polarized UV irradiation results in an insoluble film
capable of aligning liquid crystals (photoalignment)6, 7. Various crosslinked hole-
transport layers have been explored for OLEDs8-10, including one recent example for
polarized OLEDs11. OLED stability of such systems remains practically unexplored.
In particular, the choice of hole transporting moieties, the effects of film processing
(including cross-linking conditions), and the interfacial physical and electronic
properties all must be addressed. An appropriate interlayer could be expected to
improve efficiency and stability of polarized OLEDs by preventing PSS doping of the
emissive layer2, 3, reducing recombination of electrons at the anode5, and diminishing
charge accumulation within the device12.
The OLED stability study performed in Chapter 5 was motivated in part by
the desire to relate the molecular structures of blue light-emitting materials to their
OLED lifetimes. Yet, because EL emission from FFF was not realized, it remains to
be seen if fluorene homo-oligomers can be incorporated in stable OLEDs. This
question is obviously critical for the polarized OLED work discussed above. One
approach would involve a second hole transport layer, with higher HOMO and
bandgap energies, inserted between NPB and FFF. A systematic FFF doping study
could also provide insight into the degradation processes in oligo(fluorene)s.
Preliminary work13 has shown that rubrene doping as in Device III improves the
lifetime for FFF OLEDs. A second emissive dopant chosen to facilitate electron
injection across the FFF/Alq3 interface could also be included in a study of the
location and extent of the recombination zone. Structural variations in the model
compounds, particularly substitution at the 9, 9’-position of fluorene14-16, should also
177
lead to broad variation in electronic properties. Further support for the hole-transport
limitation might be provided by varying the thickness of the model compound and/or
the Alq3 ETL. Mixed-host OLEDs17, 18 could be used in an attempt to verify the
correlation between carrier mobility and device lifetime.
Research opportunities are also plentiful in the area of charge carrier transport.
Replacing rubbed polyimide with a photoalignment layer in a top-gate OFET could
allow the effects of liquid crystalline domain boundaries on field-effect mobility to be
evaluated rigorously. A single material that could form either a polycrystalline or a
glassy liquid crystalline film depending on the processing conditions would provide a
“test bed” to elucidate the effects of morphology on mobility. Furthermore, this
thesis research did not attempt to discern the field or temperature dependence of
mobility for any material studied. These measurements would allow the disorder
formalism19 for charge transport to be applied in an attempt to relate model
parameters such as zero-field mobility and positional and energetic disorder to the
molecular structures.
178
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