Organic Blue Light-Emitting Diodes and Field-Effect Transistors Based on Monodisperse

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

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To My Wonderful Wife and Family

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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).

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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

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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.

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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

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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.

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CONTENTS

Curriculum Vitae iii

iv

vi

xi

xii

xiii

Acknowledgements

Abstract

List of Charts and Reaction Schemes

List of Tables

List of Figures

1

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4

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9

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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

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76

84

100

103

1. Introduction

2. Experimental


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


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


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

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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.

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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

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43

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LIST OF REACTION SCHEMES

5.1 Synthesis procedures for model compounds. 141

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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

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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.

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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.

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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.

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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.

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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.

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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.

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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.


minute for F(Pe)5. G, Glassy; I, Isotropic.


minute for F(MB)5. G, Glassy; N, Nematic; I, Isotropic.


minute for F(Pr)4F(EH)1. G, Glassy; N, Nematic; I, Isotropic.


minute for F(MB)4F(Pr)1. G, Glassy; N, Nematic; I, Isotropic.


minute for F(Pr)5F(MB)2. G, Glassy; N, Nematic; I, Isotropic.


minute for F(Pr)5F(EH)2. G, Glassy; N, Nematic; I, Isotropic.


minute for F(MB)7. G, Glassy; SmA, Smectic A; N, Nematic; I,

Isotropic.


minute for F(MB)6F(EH)1. G, Glassy; N, Nematic; I, Isotropic.

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minute for F(MB)9. K, Crystalline; N, Nematic.


minute for F(MB)8F(EH)1. G, Glassy; N, Nematic.




minute for F(MB)7F(DMO)2. G, Glassy; N, Nematic; I, Isotropic.



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.



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.



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.


monodomain annealed film of F(MB)10F(EH)2 with a TPBI

electron-transport layer. The data were collected without


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electron-transport layer. The data were collected with a linear









monodomain annealed film of F(Pr)5F(MB)2 with a TPBI














monodomain annealed film of F(MB)5 with a TPBI electron-

transport layer. The data were collected without polarization

analysis.

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transport layer. The data were collected with a linear polarizer with

its transmission axis parallel to the nematic director defined by the

buffing direction.



transport layer. The data were collected with a linear polarizer with

its transmission axis perpendicular to the nematic director defined

by the buffing direction.



















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minute for FTO-1. G, Glassy; N, Nematic; I, Isotropic.






minute for F8T2. G, Glassy; K, Crystalline; N, Nematic; I,

Isotropic.


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.

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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.

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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


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


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


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


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


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


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


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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a) b)

2. Nalwa, H. S. (Ed.) Handbook of Advanced Electronic and Photonic Materials

and Devices, Academic Press, San Diego, (2001). 3. a) Grell, M., Bradley, D.D.C. Adv. Mater. 11, 895 (1999). b) Lussem, G.,

Wendorff, J. H. Polym. Adv. Technol. 9, 443 (1998). 4. a) Neher, D. Macromol. Rapid Commun. 22, 1365 (2001). b) Scherf, U., List,

E. J. W. Adv. Mater. 14, 477 (2002). c) O’Neill, M., Kelly, S. M. Adv. Mater. 15, 1135 (2003).

5. a) Sirringhaus, H., Brown, P. J., Friend, R. H., Nielsen, M. M., Bechgaare, K.,

Langeveld-Voss, B. M. W., Spiering, A. J. H., Janssen, R. A. J., Meijer, E. W., DeLeeuw, D. M. Nature 401, 685 (1999). b) Dimitrakopoulos, C. D., Malenfant, P. R. L. Adv. Mater. 14, 99 (2002).

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7. Era, M., Tsutsui, T., Saito, S. Appl. Phys. Lett. 67, 2436 (1995). 8. Sariciftci, N. S., Lemmer, U., Vacar, D., Heeger, A. J., Janssen, R. A. J. Adv.

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H.-G., Scherf, U., Yasuda, A. Adv. Mater. 11, 671 (1999). b) Miteva, T., Meisel, A., Knoll, W., Nothofer, H. G., Scherf, U., Müller, D. C., Meerholz, K., Yasuda, A., Neher, D. Adv. Mater. 13, 565 (2001). c) Nothofer, H-G., Meisel, A., Miteva, T., Neher, D., Forster, M., Oda, M., Lieser, G., Sainova, D., Yasuda, A., Lupo, D., Knoll, W., Scherf, U. Macromol. Symp. 154, 139 (2000).

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15. Müllen, K., Wegner, G. Electronic Materials: The Oligomer Approach, Wiley

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V. Adv. Mater. 14, 1439 (2002). 17. Fuhrmann, T., Salbeck, J. Adv. Photochem. 27, 83 (2002). 18. Gill, R. E., Meetsma, A., Hadziioannou, G. Adv. Mater. 8, 212 (1996). 19. Nierengarten, J.-F., Guillon, D., Heinrich, B., Nicoud, J.-F. Chem. Commun.

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21. Lincker, F., Bourgun, P., Masson, P., Didier, P., Guidoni, L., Bigot, J-Y.,

Nicoud, J-F., Donnio, B., Guillon, D. Org. Lett. 7, 1505 (2005) 22. Deeg, O., Bäeuerle, P. Org. Biomol. Chem. 1, 1609 (2003). 23. Lowe, C., Weder, C. Synthesis 9, 1185 (2002). 24. Gimenez, R., Pinol, M., Serrano, J. L. Chem. Mater. 16, 1377 (2004). 25. Zhang, H., Shiino, S., Shishido, A., Kanazawa, A., Tsutsumi, O., Shiono, T.,

Ikeda, T. Adv. Mater. 12, 1336 (2000). 26. a) Grafe, A., Janietz, D. Mol. Cryst. Liq. Cryst. 411, 1519 (2004). b) Geng, Y.,

Chen, A. C. A., Ou, J. J., Chen, S. H., Klubek, K., Vaeth, K. M., Tang, C. W. Chem. Mater. 15, 4352 (2003). c) Chi, C., Lieser, G., Enkelmann, V., Wegner, G. Macromol. Chem. Phys. 206, 1597 (2005).

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27. Jandke, M., Hanft, D., Strohriegl, P., Whitehead, K. S., Grell, M., Bradley, D. D. C. Proc. SPIE 4105, 338 (2001).

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O’Neill, M., Tsoi, W. C., Vlachos, P. Adv. Mater. 17, 1368 (2005). 29. a) Geng, Y., Trajkovska, A., Katsis, D., Ou, J. J., Culligan, S. W., Chen, S. H. J.

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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







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



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


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


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


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



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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107

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


SS

SS


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



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


SS

SS


SS

SS


SS

SS


SS

SS


F8T2 [c]


SS

n

F8T2 [c]


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.







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)


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


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



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.

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300 350 400 450 500 550 600



PFO

PFO

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Glassy-AmorphousFilm

Monodomain Glassy-Nematic Film

a)

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FTO-1

FTO-1

FTO-3

FTO-3

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F8T2

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Abs

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Wavelength, nm



b)

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a)

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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

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−2.0

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−40

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urre

nt (1

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)

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−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

)

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−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

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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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137

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


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

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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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Documents

Organic Blue Light-Emitting Diodes and Field-Effect Transistors Based on Monodisperse